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Journal of Thermal Analysis and Calorimetry

, Volume 133, Issue 1, pp 123–133 | Cite as

Interaction of Ca, P trace elements and Sr modification in AlSi5Cu1Mg alloys

  • Jovid Rakhmonov
  • Giulio Timelli
  • Giulia Basso
Article
  • 161 Downloads

Abstract

Thermal and microscopy analyses were carried out to investigate the interaction of Sr modification with Ca and P trace elements in high-purity and commercial-purity Al–5Si–1Cu–Mg alloys. The results show how the addition of Sr to commercial-purity alloy induces significant changes in the nucleation and growth temperatures of eutectic Si since pre-eutectic Al2Si2(CaSr) intermetallics tend to poison AlP particles, making them inactive as nucleation sites for eutectic Si. In contrast, the addition of Sr to high-purity alloy shows no apparent influence on the characteristic temperatures of Al–Si eutectic reaction, even though the microstructural investigations reveal flake-to-fibrous transition in the eutectic Si structure. This indicates that the eutectic growth temperature, commonly used to predict eutectic modification level, is not a key feature of Sr modification, but it is indirectly caused due to the presence of additional impurities in commercial-purity alloys which affect the nucleation kinetics of eutectic Si.

Keywords

Aluminium alloys C355 Eutectic modification Foundry Microstructure Thermal analysis 

Introduction

Impurities or trace elements are substances inside a confined amount of liquid, gas, or solid, which differ from the chemical composition of the rest of the material. In Al alloys, their presence occurs either naturally or added accidentally or inevitably during the different production phases. Raw materials such as alumina, petroleum coke, coal tar pitch, and recycled anode butts used during primary Al production are the major sources of impurities [1, 2].

Certain impurity elements are either difficult or expensive to remove, and their role in mechanical properties can be important; actually, dilution seems to be the only practical, although uneconomic, method to reduce the content of trace elements in Al alloys [3]. Trace elements are often not separately specified in Al alloy specifications except as “others each” and “others total” [4].

It is widely accepted how impurities lower the quality or value of the Al alloy product, but, depending on amount, may or may not make it unfit for intended use.

Calcium and phosphorus occur generally as minor impurities or trace elements in commercial-purity aluminium, while they are voluntarily introduced into some Al–Si-based alloys in order to change the microstructure of castings [5].

Calcium causes variation on the size and morphology of eutectic Si particles, from a coarse, plate-like structure to fine, fibrous Si particles, and the modification level increases with increasing the Ca content [6, 7, 8, 9]. Furthermore, the change in the eutectic solidification mode by Ca addition affects the porosity amount and distribution; a large porosity region is generally found in the centre of the hot spot in Ca-containing alloys [10].

The use of P addition to refine primary Si is a standard practice in foundry. Phosphorus reacts with the liquid Al to form AlP particles, which act as potent nucleation substrates for primary Si [11]. In addition, AlP phase plays a key role in the nucleation of eutectic Si at lower undercooling [12]. However, the presence of Ca as impurity element can reduce the number of P-based nucleants for both primary and eutectic Si by forming Al2Si2Ca phase [13], which probably nucleates on AlP, and Ca3P2 compounds [14]. Recent works have been carried out to investigate the existing interactions between Ca and P in Al–Si foundry alloys [15].

Chemical modification of Al–Si foundry alloys by means of strontium addition is extensively used in the Al–Si foundry industry. This preliminary metal treatment before casting operations improves the mechanical properties, especially fracture toughness and elongation, by affecting the size and morphology of eutectic Si crystals [16, 17, 18].

However, the presence of trace elements and impurities can decrease the efficiency of Sr modification [19], although the mechanism by which it occurs is less clear.

Despite the extensive range of literature about the Sr modification of hypoeutectic Al–Si alloys, there is still a lack of data on combined interaction of Sr, Ca and P.

In this study, an attempt is made to investigate the effects of interaction between the trace elements Ca and P and the Sr modification on the solidification path and microstructural evolution of AlSi5Cu1 Mg alloys, one of most used alloys in the in aerospace, automotive, and household industries. These groups of alloys have become increasingly important in recent years, mainly due to their high strength at room and high temperatures [20].

Experimental

Materials and processing

A hypoeutectic primary Al–5Si–1Cu–Mg alloy (equivalent to the US designation C355.0) with different purities was used as base alloy throughout the investigations. The alloy was supplied by Slovalco (Slovak Aluminium Co., Slovakia) in the form of commercial 10-kg ingots with two different levels of trace elements, P and Ca. The alloy containing less than 5 ppm P and almost no Ca is defined as high-purity (HP) alloy, while the alloy with 10 ppm P and 30 ppm Ca is regarded as commercial-purity (CP) alloy. The chemical compositions of the CP and HP base alloys are presented in Table 1, and they were measured on separately poured samples by using a Thermo Scientific™ (Waltham, MA, USA) ARL 3460 iSpark™ optical emission spectrometer (OES) with a wavelength ranging from 130 nm to 780 nm. If present, phosphorus was below the lower detection limit (< 5 ppm) of OES in HP base alloy.
Table 1

Chemical compositions of CP and HP base alloys (mass/%)

Alloy

Si

Fe

Cu

Mg

Mn

Ti

Pa

Ca

Al

CP

5.31

0.101

1.22

0.501

0.003

0.135

0.0010

0.0030

bal.

HP

5.37

0.106

1.23

0.590

0.004

0.127

0.0001

bal.

aIf present, phosphorus is below the lower detection limit (< 5 ppm) of OES

To investigate the interaction of Sr modification with Ca and P trace elements, both HP and CP alloys were alloyed with different Sr levels (0, 65, 100, 160, and 210 ppm). A weighted Al–10Sr master alloy in the form of rod was added into the molten bath to ensure that the Sr level in the melt reached the desired content.

About 4 kg of CP and HP base alloys were melt in a SiC crucible inside an electric furnace at 750 ± 5 °C. The liquid bath was held in the furnace for 1 h and then stirred and surface-skimmed to remove the dross. In order to ensure the homogeneity of the molten metal during the entire experimentation, the chemical composition of the melt was measured on samples separately poured at the beginning and the end of every set of casting. The chemical composition did not change within every set of analysed alloy.

The molten metal was then poured into a boron nitride-coated cylindrical steel cup (outer diameter 45 mm, height 60 mm, and a uniform wall thickness of 3 mm), preheated at about 700 °C. Casting of the alloys with Sr addition was carried out following the same procedure used for the base alloys, however, after the introduction of Sr modifier into the molten bath. The contact time between the Sr modifier and the melt was about 15 min.

Thermal analysis

Figure 1 shows schematic illustration of the thermal analysis set-up. When the melt was poured into the cup, a K-type thermocouple (ø1 mm), which was fixed at the lid of the cup and covered with steel tube, was inserted into the melt. The thermocouple was positioned along the central axis of the cup at a depth of 20 mm from the lid bottom surface. The thermocouple was calibrated against the melting point of pure aluminium, assuming a temperature of 660.0 °C. The temperature and time were collected using data acquisition system with a sampling frequency of 0.25 s−1, analog-to-digital converter accuracy of 0.1 °C, and connected to a personal computer where they were processed with Alutron’s Thermal Analyzer TAW32 software. This commercial software is able to plot the cooling curve and corresponding derivative (dT/dt), and the baseline [21, 22].
Fig. 1

Schematic illustration of the thermal analysis set-up with steel cup and thermocouple. All dimensions are in mm

The derivative of the curve represents the rate of cooling of metal in the solidifying sample. When the derivative increases, this means that something has happened to slow the rate of cooling, such as the formation of a new phase that releases latent heat.

The baseline is calculated by fitting user-defined intervals of the dT/dt curve before and post-solidification according to the method of Tamminen [22]. The baseline estimates the heat transfer due to cooling effects and by assuming that no phase transformation takes place. Hence, reactions involving the release of latent heat (exothermic reactions) during solidification can be detected whenever a positive deviation from the baseline occurs.

The cooling rate of the alloys in the solidification interval was 0.17 ± 0.1 °C s−1. At least two thermal analysis tests were conducted for each experimental alloy.

Figure 2 and Table 2 display the definition method and the description of thermal analysis parameters used in the present study.
Fig. 2

a Typical cooling curve (T), associated first derivative (dT/dt) and baseline of CP base alloy; b enlargement of the eutectic arrest and the characteristic temperatures (TN, eu, Tmin, eu, TG, eu) drawn from the curve

Table 2

Definition of the characteristic temperatures during solidification from Fig. 2

Symbol

Description

T N, α

Nucleation temperature of α-Al phase

T min, α

Minimum temperature of α-Al phase

T G, α

α-Al growth temperature

T N, eu

Nucleation temperature of the eutectic

T min, eu

Minimum temperature of the eutectic

T G, eu

Eutectic growth temperature

ΔTR, eu

Recalescence undercooling of eutectic

T S

Solidus temperature

ΔTS

Solidification interval (ΔTS = TN, α − TS)

Microstructure evaluation

Samples for microstructural characterization were sectioned near the tip of thermocouple to correlate microstructural features with the solidification parameters obtained from thermal analysis. Samples were then ground and polished following standard metallographic methods. Optical microscopy (OM) and scanning electron microscope (SEM) equipped with an energy-dispersive spectrometer (EDS) were used to study the microstructural changes occurring after Sr modification.

In order to identify the chemical composition of the phases, EDS data were acquired at low accelerating voltage of 10 kV in order to reduce the beam interaction with the matrix volume. The stoichiometry of the intermetallic compounds was determined by comparing the results from the semi-quantitative EDS to the composition range of the phases reported in the literature for similar alloy systems and by thermodynamic calculations through CALPHAD modelling using JMatPro® commercial software.

Results and discussion

Commercial-purity alloys

Thermal analysis of commercial-purity alloys

The characteristics of main solidification reactions determined by thermal analysis of the studied CP alloys are presented in Table 3. The formation of primary α-Al remained unaffected upon Sr addition [17], while the solidification in Sr-modified CP alloys completed at higher temperatures by about 6 °C, if compared to the solidus temperature of CP base (unmodified) alloy. Strontium is commonly added to Al–Si alloys to influence the nucleation and growth kinetics of eutectic Si [23, 24]. The present study revealed how the formation temperature of eutectic Si shifted downward upon Sr addition (Table 3; Fig. 3).
Table 3

Characteristic temperatures (°C) from thermal analysis of CP alloys with different Sr contents

 

Addition (ppm)

T N, α

T min, α

T G, α

T N, eu

T min, eu

T G, eu

ΔTR, eu

T S

ΔTS

Base alloy

634.4

628.0

629.5

567.3

564.7

563.8

− 0.9

483.8

150.6

Sr

65

634.5

628.5

629.8

559.4

554.3

556.5

2.2

490.3

146.1

Sr

100

635.8

628.3

629.6

557.4

551.8

553.6

1.8

489.8

141.9

Sr

160

635.5

628.9

630.1

560.2

552.6

554.4

1.8

490.2

144.3

Sr

210

635.1

628.8

629.9

559.9

556.8

558.7

1.9

493.1

148.3

Fig. 3

Cooling curves in the region of Al–Si eutectic reaction recorded during solidification of CP alloys with different Sr contents

In CP base alloy, the eutectic Si reaction initiated at 567.3 °C during solidification (see Table 3). According to the predicted phase reactions assuming Scheil solidification condition and occurring during solidification of the CP base alloy, the Al–Si eutectic reaction is expected to start at 568 °C (Fig. 4). The difference in the nucleation temperatures obtained from the thermodynamic calculation and the experimental investigation can be considered as the indication of the degree of liquid metal undercooling needed for the nucleation of eutectic Si. A nucleation undercooling of ~ 1 °C indicates the presence of potent nucleation sites, which enables the eutectic Si to nucleate at low undercooling. Phosphorus, which is commonly available as an impurity in Al alloys, has been found to promote the nucleation of eutectic Si at low undercooling [23, 24, 25, 26]. It has been reported how in Al–Si alloys with P content greater than about 3 ppm, the formation of pre-eutectic AlP particles can occur; these particles then act as heterogeneous nucleation site for eutectic Si during alloy solidification [12, 26]. The CP base alloy investigated in the present study contains 10 ppm P (see Table 1), which can be credited for the nucleation of eutectic Si at low undercooling.
Fig. 4

Thermodynamic calculations of phase fraction of primary Al, eutectic Si and Al2Si2(CaSr) phases in CP alloys after 65 and 160 ppm Sr additions, respectively. Thermodynamic calculations were conducted assuming Scheil solidification conditions

Adding Sr (65 ppm) caused a significant depression of the nucleation temperature of eutectic Si (see Table 3; Fig. 3). Thermodynamic calculation of CP alloy modified with 65 ppm Sr predicted the formation of pre-eutectic Al2Si2(SrCa) intermetallics (Fig. 4), which seem responsible for the changes in the nucleation kinetics of eutectic Si reaction. It is reported how ternary Al–Si–Sr and Al–Si–Ca intermetallics can remove AlP particles from the remaining interdendritic liquid prior to Al–Si eutectic reaction [9, 25, 27], which necessitates greater undercooling for the nucleation of eutectic Si.

Upon progressive increase in the Sr content up to 160 ppm, further depression in the formation temperature of eutectic Si was observed (Fig. 3); however, when the Sr level in the alloy reached 210 ppm Sr, the Al–Si eutectic reaction occurred at slightly higher temperatures compared to CP alloys with lower Sr contents. More depression in the formation temperature of eutectic Si upon 100 and 160 ppm Sr additions can be due to complete removal or deactivation of AlP particles from the melt prior to Al–Si eutectic reaction. Thermodynamic calculation of CP alloy (Fig. 4) modified with 160 ppm Sr predicted the formation of more amount of pre-eutectic Al2Si2(SrCa) intermetallics than in CP alloy containing 65 ppm Sr.

Another solidification parameter that is affected by Sr addition is the recalescence ΔTR, eu, which is the difference between TG, eu and Tmin, eu values. As listed in Table 3, the recalescence is negative in unmodified alloy and positive in Sr-modified CP alloys. This is consistent with the literature [18, 25], and this behaviour is commonly accepted as an indication of the eutectic modification occurring by Sr or Na addition.

The growth temperature, TG, eu, of eutectic Si was also found to have strong correlation with the modification level that can be achieved upon Sr or Na addition [28, 29, 30]. In the present study, TG, eu showed similar trend with TN, eu, i.e. a continuous drop of the growth temperature occurred with progressively increasing the Sr content up to 160 ppm, while introducing 210 ppm Sr caused an upward shift of the eutectic Si growth temperature. According to thermal analysis results, the addition of Sr in the range between 100 and 160 ppm appears to be sufficient to achieve the best modification level in the present investigation conditions.

Microstructure and identification of intermetallic phases

Figure 5 shows the typical microstructures observed in CP alloys before and after Sr (100 ppm) addition. The main phases appearing in the microstructure of both modified and unmodified alloys are α-Al, eutectic Si, Al5FeSi, Al8Mg3FeSi6, Mg2Si, Al5Si6Cu2Mg8 and Al2Cu. In Sr-modified alloys, the Al2Si2(SrCa) particles were also observed throughout the microstructure (Fig. 6). Upon increasing the Sr content in the alloy, the fraction of Al2Si2(SrCa) intermetallics increased, which is also consistent with the predicted results of thermodynamic calculations.
Fig. 5

Backscattered electron images showing typical microstructures of a CP base alloy and b CP alloy after 210 ppm Sr addition. Intermetallic compounds are indicated throughout

Fig. 6

Backscattered electron image of Al–Si–Ca–Sr intermetallic observed in Sr-modified (100 ppm) CP alloy with corresponding EDS composition maps, showing the distributions of Al, Si, Sr, Ca and P

Microstructural investigations revealed how P is also bound to pre-eutectic Al2Si2(SrCa) intermetallics (see Fig. 6). These quaternary Al–Si–Sr–Ca intermetallics seem to poison AlP particles as commonly done by ternary Al–Si–Ca and Al–Si–Sr intermetallics [15, 27]; hence, the remaining liquid metal, which is free from AlP, tends to further undercool till certain solid particles become active nucleation sites for eutectic Si.

Adding Sr also promoted the formation of Mg2Si compounds over Al5Si6Cu2Mg8 intermetallics (see Fig. 5). It has been reported how Sr causes the segregation of Cu to the regions away from Al–Si eutectic during solidification, thus leading to the post-eutectic reaction involving Al2Cu in a divorced mode, rather than the post-eutectic Al–Al2Cu–Si reaction in coupled mode. Changes in Cu distribution induced by Sr addition seem also responsible for more preferential formation of Mg2Si over Al5Si6Cu2Mg8 compounds [31].

The eutectic microstructure of CP base alloy (Fig. 7a) exhibits coarse Si flakes. Adding Sr (65 ppm) induced a partial modification of eutectic structure (Fig. 7b). Further increase in Sr level, first to 100 and then to 160 ppm allowed maximum modification level of eutectic structure in the present investigation conditions (Fig. 7c, d). However, the addition of 210 ppm Sr seems redundant as microstructure reveals the evidence of over-modified eutectic structure (Fig. 7e).
Fig. 7

Typical eutectic Si structure observed in CP alloys containing a 0 ppm, b 65 ppm, c 100 ppm, d 160 ppm and e 210 ppm Sr

Statistical analysis of the size distribution of eutectic Si particles also confirms how the amount of coarser Si particles in CP alloy with 210 ppm Sr is increased at the expense of finer particles when compared CP alloy with 100 ppm Sr (Fig. 8). These findings are consistent with the thermal analysis results which predicted to the optimum modification level at Sr levels of 100–160 ppm.
Fig. 8

Size distributions of eutectic Si particles in CP alloys with different Sr levels

High-purity alloys

Thermal analysis of high-purity alloys

The characteristic temperatures of the solidification reactions determined from HP alloy with different Sr levels are listed in Table 4. Strontium addition showed no apparent influence on the precipitation kinetics of primary α-Al; however, it slightly reduced the freezing range due to the upward shift of the solidus temperature. Similar results were observed in commercial-purity alloys.
Table 4

Characteristic temperatures (°C) from thermal analysis of HP alloys with different Sr contents

 

Addition (ppm)

T N, α

T min, α

T G, α

T N, eu

T min, eu

T G, eu

ΔTR, eu

T S

ΔTS

Base alloy

636.2

626.7

626.8

560.7

556.7

558.7

2.0

482.7

153.4

Sr

65

634.9

626.9

627.7

560.7

556.4

557.7

1.3

487.0

147.8

Sr

100

635.5

627.1

627.6

560.2

556.6

558.0

1.4

487.1

148.4

Sr

160

634.8

626.6

626.9

560.8

555.5

557.7

2.2

486.3

148.5

Sr

210

634.2

627.9

628.6

561.3

557.4

559.2

1.8

489.3

144.8

Unlike in CP alloys, adding Sr to HP alloy showed no evident influence on the precipitation kinetics of eutectic Si (see Table 4; Fig. 9). Nucleation of eutectic Si in HP base alloy started at ~ 560.7 °C, which is lower by ~ 6 °C compared to that of CP base alloy. The characteristic temperatures of Al–Si eutectic reaction observed in HP base alloy are similar to those observed in Sr-modified CP alloys. This implies that Sr tends not to affect the nucleation kinetics of eutectic Si in HP alloys except at Sr addition level of 210 ppm, which slightly shifted the growth temperature of eutectic Si, TG, eu, towards higher temperatures with respect to HP alloys modified with lower Sr levels. As previously mentioned, in Al–Si alloys with P content greater than about 3 ppm, the pre-eutectic AlP becomes a thermodynamically stable phase [26]; low lattice mismatch between AlP and Si leads to the nucleation of latter phase on the former one at low undercooling [12]. Although P content in HP base alloy was not reliably quantified as it was below the lower detection limit (< 5 ppm) of OES, it is believed that the HP alloy contains very low P, and consequently, the formation of pre-eutectic AlP compounds is avoided.
Fig. 9

Cooling curves in the region of Al–Si eutectic reaction recorded during solidification of HP alloys with different Sr contents

The recalescence and growth temperature of eutectic Si have been found to have strong correlation with the modification level that can be achieved upon Sr or Na addition in Al–Si alloys; therefore, thermal analysis technique has been used in foundry up to now as a useful tool to assess the eutectic modification level [28]. However, based on the results of the present study, it can be inferred that evaluation of eutectic modification in Al–Si alloys by thermal analysis technique is highly dependent on the amount of impurities, and this technique cannot be reliable for Al–Si alloys containing limited amount of P element.

Microstructure and identification of intermetallic phases

Figure 10 shows the typical microstructures of HP base (unmodified) alloy and HP alloy modified with 100 ppm Sr. The main phases constituting the microstructure are α-Al, eutectic Si, Al5FeSi, Al8Mg3FeSi6, Mg2Si, Al5Si6Cu2Mg8 and Al2Cu. The Al2Si2Sr particles were also observable in Sr-modified HP alloys (Fig. 11). The eutectic microstructure of HP base alloy (Fig. 12a) exhibits more refined silicon flakes, with shorter and closely spaced silicon crystals with respect to that of CP base alloy (Fig. 7a). A comparison of eutectic Si size distributions between CP and HP base alloys (Figs. 8, 13) indicates how the Si particles in HP base alloy are finer than in CP base alloy. As previously mentioned, the nucleation of Si in HP base alloy required greater undercooling if compared to commercial-purity alloys. Since the nucleation rate of eutectic Si has been found to be smaller in largely undercooled melts, the solid–liquid eutectic interface is much smaller and consequently advances at a higher growth rate; this increase in the growth rate contributes to refine the eutectic structure [25].
Fig. 10

Backscattered electron images showing typical microstructures of a HP base alloy and b HP alloy after 210 ppm Sr addition. Intermetallic compounds are indicated throughout

Fig. 11

a Backscattered electron image of Al2Si2Sr intermetallic in Sr-modified (100 ppm) HP alloy and b corresponding EDS spectra; EDS composition maps, showing the distributions of Al, Sr, and Si

Fig. 12

Typical eutectic Si structure observed in HP alloys containing a 0 ppm, b 65 ppm, c 100 ppm, d 160 ppm, and e 210 ppm Sr

Fig. 13

Size distributions of eutectic Si particles in HP alloys with different Sr levels

Adding Sr (65 ppm) induced an appreciable modification of eutectic structure in HP alloy (Fig. 12b); the modification level achieved in HP alloy after 65 ppm Sr addition is better than that observed in CP alloy at the same Sr level (Fig. 7b). This indicates how the eutectic modification of HP alloy requires lower Sr amount than that needed for the modification of CP alloy. This behaviour can be due to the presence of Ca as an impurity in CP alloy. It is believed that due to the presence of Ca in CP alloy, more Sr are bound to pre-eutectic Al–Si–Sr–Ca compounds, thus lowering Sr amount involved in changing the growth of eutectic Si.

According to thermodynamic calculations of phase fractions, the amount of Al2Si2(CaSr) in Sr-modified CP alloys (Fig. 4) is comparably higher than that of Al2Si2Sr phase in Sr-modified HP alloys (Fig. 14). The formation of intermetallics based on the quaternary Al–Si–Sr–Ca system seems to involve more Sr, hence lower Sr amount in remaining liquid compared to the Sr available in HP alloy for eutectic modification.
Fig. 14

Thermodynamic calculations of phase fraction of primary Al, eutectic Si and Al2Si2Sr phases in HP alloys after 65 and 160 ppm Sr additions, respectively. Thermodynamic calculations were conducted assuming Scheil solidification conditions

Typical microstructure representing the analysed alloy with 160 ppm Sr indicates fully modified eutectic structure in CP alloy (Fig. 7d) and over-modified eutectic in HP alloy (Fig. 12d). Based on these results, it can be inferred that complete modification of HP alloy requires lower amount of Sr; the formation of intermetallics based on quaternary Al–Si–Sr–Ca system can be responsible for the addition of greater Sr content to obtain a complete eutectic modification in CP alloys.

Conclusions

Thermal analysis technique has been used to study the interaction of Sr modification with Ca and P trace elements in Al–5Si–1Cu–Mg alloy. The following conclusions can be drawn.
  • The growth temperature and the recalescence of eutectic Si reaction detected by thermal analysis show strong correlation with the actual eutectic modification level in commercial-purity alloys.

  • The addition of Sr to commercial-purity alloys affects, firstly, the nucleation stage by deactivating the potent nucleation sites, i.e. AlP particles, for eutectic Si, and then, the growth stage of eutectic Si by adsorbing into the growing front of Si crystals.

  • No correlation between characteristic temperatures of eutectic Si and actual modification level exists in high-purity alloys. As P content in high-purity alloys is low, the pre-eutectic AlP phase remained unfavourable to form during alloy solidification; hence, the absence of potent nucleation site, e.g. AlP, necessitates greater undercooling for the nucleation of eutectic Si.

  • The characteristic temperatures of Al–Si eutectic reaction in unmodified high-purity alloys is nearly the same as that of Sr-modified alloys.

  • The changes in the growth temperature and the recalescence of eutectic Si reaction occurring in commercial-purity alloys upon Sr addition are directly dependent on the nucleation kinetics of eutectic Si and insensitive to the action of Sr in the growth stage of eutectic Si.

  • The growth temperature and the recalescence of eutectic Si obtainable by thermal analysis are unreliable parameters in assessing/predicting the eutectic modification level in Al–Si alloys, particularly in high-purity alloys.

  • The presence of Ca (~ 50 ppm) in commercial-purity alloys deteriorates the efficiency of Sr modification due to the formation of pre-eutectic Al2Si2(CaSr) intermetallics.

  • The evaluation of eutectic modification in Al–Si alloys by thermal analysis technique, widely used in foundry, is highly dependent on the amount of impurities, and this technique cannot be reliable for Al–Si alloys containing limited amount of P element.

  • In metal casting industry, a deeper control of the initial Ca and P levels can significantly affect the Sr modification treatment of the molten metal; this can lead to an increasing improvement of component quality.

Notes

Acknowledgements

The authors would like to acknowledge Dr. A. Fabrizi for his precious work to perform experiments and characterizations.

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Copyright information

© Akadémiai Kiadó, Budapest, Hungary 2018

Authors and Affiliations

  1. 1.Department of Management and EngineeringUniversity of PadovaVicenzaItaly
  2. 2.Department of Mechanical EngineeringNavoi State Mining InstituteNavoiUzbekistan

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