Introduction

The Charpy impact test is a popular test for measuring the toughness of materials; it should be remembered, however, that it was developed before the theory of fracture mechanics was formulated.1 The total absorbed energy required for breaking the material is measured by the high strain rate fracture of a standard notched or unnotched specimen.2 The total absorbed energy dissipated in breaking the test specimen is then measured based on the difference in the potential energy of the striking pendulum.3 Also, the total absorbed impact energy may be represented by the area located under the load–time curve which is obtained from an instrumented Charpy impact testing machine.4 In the presence of a notch, the impact energy will be much lower than that in the case of an unnotched specimen. The unnotched impact specimen provides highly accurate results, drawing attention to the fact that the impact values depend on the microstructure rather than on the specimen configuration.2

During the past ten years, the research group at the Université du Québec à Chicoutimi has carried out an intensive program on the parameters that control the toughness of aluminum alloys and their composites using unnotched samples.5 – 10 The main results inferred from these studies show that the Sr-modified alloys exhibit higher levels of impact values for all aging times used at 220 °C (428 °F) than those obtained for the same alloys under T4 conditions. In general, the impact toughness behavior of aged hypoeutectic alloys depends to a great extent on their aging response at 180 °C (356 °F) or 220 °C (428 °F) aging temperatures. The presence of undissolved intermetallic phases, mainly Al2Cu and α-Al15(Mn,Fe)3Si2, increases alloy brittleness and thus also reduces alloy toughness. The segregation of Al2Cu caused by Sr addition is a further parameter to be considered. The combination of Sr modification and solution treatment increases the impact energy values of alloys to a noticeable degree, particularly at low levels of intermetallics. The application of Sr modification in conjunction with solution treatment significantly improves both the crack initiation and the crack propagation resistance, and consequently, results in a noticeable improvement in the overall impact energy of the alloys.

The present investigation was undertaken to study the effect of different metallurgical parameters on the microstructure and performance of 356 (Al–Si–Mg) and 413 (Al–Si–Mg–Cu) alloys covering a range of Si content and consequently the response to heat treatment. These alloys are among the alloys that are widely used in the automotive industries. The parameters studied included: alloy composition, hydrogen level, modification by strontium (Sr), grain refining (mainly by TiB2) as well as heat treatment.

Experimental Procedure

The chemical compositions of the as-received 356 and 413 alloys are given in Table 1b which are very close to those recommended by the Aluminum Association standards (Table 1a). The alloys were melted using an electrical resistance furnace. The melting temperature was maintained at 730 ± 5°C (to insure high fluidity of the liquid metal and hence sounder castings); the molten metal was degassed using pure dry argon injected into the molten metal (0.85 m3/h) by means of a rotary graphite impeller.

Table 1 (a) AA Standard Chemical Composition of as-Received Alloys,13 (b) Actual Chemical Composition of as-Received Alloys
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The degassing time/speed was kept constant at 30 min/150 rpm in order to obtain a lower hydrogen level of ~0.1 mL/100 g in the melt. Measured amounts of Al–10 %Sr and Al–5 wt%Ti–1 wt%B master alloys were added to the molten metal after degassing. Samples for chemical analysis were also extracted simultaneously for each melt/casting produced. The chemical compositions representing the average of three spectrometric analyses for each of the alloys investigated are listed in Table 2 (it should be noted that the Si in 356 alloy is slightly below the recommended level as listed in Table 1a). The melt hydrogen level was increased through addition of raw potatoes to the melt to obtain levels of 0.2 mL/100 g (low) and 0.4 mL/100 g (high) and determined using the Leco subfusion technique.

Table 2 Chemical Composition of the Alloys Used in the Present Study
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A star-like mold which provided a high cooling rate was used to prepare the impact and the hardness test specimens as well as to take metallographic measurements. Each casting provides ten impact bars which are cut from the casting and then machined to the required ASTM specifications to conduct Charpy impact tests subsequently. Hardness test specimens measuring 10 mm × 10 mm × 55 mm were cut from the casting and polished with fine sandpaper to remove any machining marks.

The hardness and impact test samples prepared for each alloy composition were solution heat-treated (SHT) at 540 °C (1004 °F) for 8 h for the 356 based alloys and at 495 °C (923 °F) for 8 h for the 413 based alloys, quenched in warm water at 65 °C (149 °F) and artificially aged for 5 h at 155, 180, 200, 220 (311, 356, 392, 428 °F) and 240 °C (464 °F). All heat treatments were conducted in a forced-air Blue M electric furnace equipped with a programmable temperature controller, accurate to ±2 °C. For each individual heat treatment condition, five test bars were used.

The hardness measurements were carried out on the as-cast and heat-treated samples using a Brinell hardness tester, applying a steel ball of 10 mm diameter and a load of 500 kgf for 30 s. An average of four readings obtained from two perpendicular surfaces was taken to represent the hardness value in each case.

A computer-aided instrumented SATEC SI-1 Universal Impact Testing Machine (SATEC Systems Inc., Model SI-1D3) was used to carry out the impact tests. A data acquisition system connected to the impact machine monitored the dynamic behavior of the test specimen and measured the load and energy values as a function of time. Samples for metallographic observations were sectioned from the impact-tested bars of all the alloys studied about ~10 mm below the fracture surface. The samples were mounted in bakelite and polished to a fine finish.

The microstructures of the polished sample surfaces were examined using an optical microscope linked to a Clemex image analysis system. The eutectic silicon particle characteristics, including area, length, aspect ratio, roundness and density, were measured and quantified. For each sample, 50 fields at a magnification of 500× were examined, so as to cover the entire sample surface in a regular and systematic manner. In addition, porosity measurements were carried out, over 50 fields per sample, at a magnification of 100×. The porosity parameters measured were percentage porosity, pore area and pore length. As a rule, the outer edges of a sample were avoided in taking these measurements so as to eliminate any distortions which might occur in the peripheral regions. Some selected samples were examined using a Hitachi SU-8000 field emission gun scanning electron microscope (FESEM). Prior to examination, samples were re-polished using ion bombardment to remove the surface oxides.

For measurements of grain size, the polished samples were chemically etched using Keller’s reagent (66 vol% HNO3, 33 vol% HCl and 1 vol% HF) and diluted with water to slow down the etching process. Once etching was completed, the sample surfaces were rinsed, dried and then photographed using a setup consisting of four projectors emitting red, blue, green and yellow light. The grain sizes were measured from the digital photographs using Sigma Scan Pro 4.0 software, employing a method similar to the line intercept method. The grain size for each sample was obtained as the average of 80 readings.

Results and Discussion

Microstructural Characterization

It is a well-established fact that grain size in Al–Si alloys is greatly influenced by the added amount of TiB2.11 , 12 Sigworth and Guzowski14 used Al–5 %Ti–1 %B master alloy to obtain minimum grain size in 356 alloys. In the present study, three concentrations of TiB2 were used, corresponding to 0, 0.0075 and 0.25 wt% Ti. The grain sizes obtained are listed in Table 3. Using low concentration of TiB2, i.e., 0.0075 % Ti, a slight decrease in the grain size was observed (from 811 to 692 µm for 356 alloys and from 2466 to 2156 µm for 413 alloys). Increasing the TiB2 concentration—measured by the amount of Ti—to 0.25 % reduced the grain size to 289 µm for 356 alloys and to 275 µm for 413 alloys (approximately 90 %) as shown in Figure 1. Figure 2 shows the grain size in CDTE alloy grain refined with 0.25 % Ti, whereas Figure 3 shows the size and distribution of TiB2 particles as viewed within a small pore using the FESEM technique. The average TiB2 particle size is about 300–500 nm.

Table 3 Grain Size of the 356 and 413 Alloys Studied
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Figure 1
figure 1

Macrostructure of grains observed in 356 alloys following addition of: (a) 0.0075 wt% Ti (DBDT alloy) and (b) 0.25 wt% Ti (DBDTE alloy).

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Figure 2
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Macrostructure of the grains observed in 413 alloys following the addition of: (a) 0.0075 wt% Ti (DCDT alloy) and (b) 0.25 wt% Ti (DCDTE alloy).

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Figure 3
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(a) Backscattered electron image showing fine TiB2 particles in DBDTE alloy and (b) corresponding X-ray image of Ti, (c) an enlarged image of the circled area in (a).

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Table 4 lists the measured parameters of eutectic Si particles in 356 alloys before and after solution heat treatment. As may be seen, the average area of the Si particles of non-modified alloys exhibits an increase after solution heat treatment. Considering the Sr-modified alloys, the Si particles revealed a tendency for spheroidization followed by coarsening. The changes in the morphology of the Si particles from acicular to fibrous explain the observed variation in Si particle measurements as exemplified in Figure 4. Figure 5a shows typical Sr-modified alloys revealing ultra-fine Si particles compared to Figure 4a. Following solution heat treatment at 540 °C (1004 °F) for 8 h, the Si particles underwent spheroidization and coarsening as displayed in Figure 5b. Figure 4c is an example of Si particles observed in non-modified 356 alloy in the as-cast condition revealing a significant variation in their size across the sample surface,15 explaining the large standard deviations reported in Table 4.

Table 4 Si Particle Characteristics of 356 Alloys (A) Before Solution Heat Treatment, (B) After Solution Heat Treatment
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Figure 4
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Changes in the morphology of non-modified eutectic Si particles in BDT alloy: (a) as-cast, (b) after solution heat treatment, (c) extracted eutectic Si particles in as-cast 356 alloys.15

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Figure 5
figure 5

Morphology of eutectic Si in Sr-modified 356 (BDS) alloy: (a) as-cast, (b) after solution heat treatment.

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Table 5 lists the measured Si particle parameters for 413 alloys before and after solution heat treatment. As in the case of 356 alloys, non-modified alloys revealed a slight increase in the particle size after solution heat treatment. However, the addition of Sr resulted in more or less uniform distribution of fibrous Si particles with spheroidization after solution heat treatment. The presence of trace elements in 413 alloy (mainly Fe, Cr and Mn) resulted in the precipitation of α-Al15(Fe,Mn,Cr)3Si2, as shown in Figure 6. In order to arrive at a clear understanding of the observed increase in the Si particle size in non-modified alloys, the solutionizing time was increased to 48 h (413 alloys). Figure 7 exemplifies the increase in the Si particle thickness following prolonged solutionizing treatment (SHT) at 495 °C (923 °F)—circled area—compared to Figure 6a. Figure 8 shows a schematic representation of the proposed mechanism of coarsening phenomenon based on the theory developed by Ostwald.16

Table 5 Si Particle Characteristics of 413 Alloys (A) Before Solution Heat Treatment, (B) After Solution Heat Treatment
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Figure 6
figure 6

Eutectic Si particle distribution in 413 alloy: (a) before Sr modification-CDT alloy note the precipitation of α-Fe in the interdendritic regions, (b) after Sr modification-DCDS alloy note the precipitation of α-Fe within the α-Al.

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Figure 7
figure 7

Eutectic Si particles in DC alloy following prolonged solutionizing at 495 °C: (a) dissolution of fine Si particles, (b) fragmentation, (c, d) coarsening.

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Figure 8
figure 8

Proposed schematic diagram for the observed increase in the particle size after solution heat treatment.

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Different types of porosity have been reported in the present study. They include gas porosity caused by the increase in the hydrogen content, shrinkage porosity due to poor feeding, a mixture of gas and shrinkage porosity (Figure 9a), oxide films or a mixture of these types.17 – 19 Table 6 lists the measured porosity parameters in 413 alloys. In addition to the porosity caused by increasing the hydrogen content (DCDH4 alloy vs. DC alloy), addition of Sr also leads to porosity formation as given in Table 6 for 413 alloys (DCDTES alloy vs. DC alloy). Obviously, the increase in the porosity due to hydrogen addition is an order of magnitude higher than that introduced by increasing the amount of added Sr. Also, over-modification appeared in the precipitation of Al4SrSi2 in the form of tetrahedral particles as displayed in Figure 9b.20

Figure 9
figure 9

(a) A mixture of gas and shrinkage porosity observed in CDH4 alloy, (b) presence of Al4Si2Sr particles in DCDTES alloy in the as-cast condition, (c) EDS corresponding to (b). Si K β = 1.84 keV, Sr L α = 1.94 keV.

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Table 6 Porosity Characteristics Observed in the 413 Alloys
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Hardness and Impact Properties

The following aging sequence of an Al–Mg–Si alloy is generally accepted as being correct:20 α → G.P. zone → metastable β′phase → equilibrium β′′ phase. The β′′ phase is considered to be extremely important because it may play a major role as an effective strengthening precipitate in the Al–Mg–Si alloys.21 The chemical composition of the β′′ phase is Si:Al:Mg = 6:3:1. In a recent study22 it was proposed that the precipitation sequence was: clusters of Si atoms → GP-I zones → GP-II zones/β′′ → β′ → β-Mg2Si. Ohmori et al.23 studied the aging process in Al–Si–Mg alloys during continuous heating. Their results show a slight hardening at temperatures 77–152 °C (171–306 °F) caused by the formation of solute atom clusters, a small amount of softening at temperatures around 177 °C (351 °F), and a sharp and large hardening at around 277 °C (531 °F), attributed to the precipitation of needle-shaped β′′, followed subsequently by a large drop in hardness due to the formation of β-Mg2Si.

Figure 10 represents the hardness data for the 356 alloys used in the present study following aging at different temperatures. Apparently the variation in the metallurgical parameters has a minor effect on the alloy strength. However, varying the aging temperatures resulted in a noticeable change in the alloy hardness due to precipitation of Mg2Si phase particles.24 In contrast to the work of Ohmori et al.,22 only a single peak was observed at 180 °C (356 °F), caused by the precipitation of both β′ and β′′ particles, followed by continuous softening due to progress in the precipitation and coarsening of incoherent β-Mg2Si.25

Figure 10
figure 10

Variation in the hardness of 356 alloys as a function of aging temperature.

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Tash et al.26 and Moustafa et al.27 found that the 319 alloy hardness first increased with an increase in aging temperature up to 180 °C (356 °F) and thereafter decreased as the aging temperature was increased. The increase in hardness with the addition of Cu and Mg may be attributed to the formation of the hard and brittle (metastable) intermetallic Al2Cu and Al2CuMg phases, as well as to an increased bonding of the Si particles with the matrix, where the thermal energy is sufficient to precipitate such intermediate phases. The decrease in hardness at aging temperatures above 180 °C (356 °F) may be due to the coarsening of the various microconstituents as well as to a decrease in cohesion with the matrix.

The coarsening of hard intermetallic phases may reduce the barrier to dislocation movement and, hence, to flow stress/hardness. Silicon alone contributes little to the strength of aluminum casting alloys, yet provides an observably effective level of strengthening when combined with magnesium to form Mg2Si. The Mg2Si phase is soluble in the alloy to a limit of ~0.7 wt% Mg and provides the precipitation-strengthening basis for the heat-treatable alloys.28

Figure 11 displays the variation in the hardness of 413 alloys as a function of aging temperature. As in the case of 356 alloys, variation in the metallurgical parameters has a slight effect on the alloy strength. However, the maximum attainable strength at 180 °C (356 °F) is about 67 BHN compared to 115 BHN in the case of 356 alloy (DBDTES alloy, degassed, Sr-modified and grain refined with 0.25 wt% Ti). It should be kept in mind that the 413 alloys contain very low amount of Mg, approximately 0.04 %, compared to 0.3 % in the case of 356 alloys. Thus, hardening by Mg2Si phase particles is more effective than hardening by Al2Cu phase precipitation which is in good agreement with the published literature.29 Figure 12 shows the difference in the microstructure of the two alloys used in the present work, whereas Figure 13 depicts the morphology and size of the precipitates observed in 356 and 413 alloys aged at 180 °C for 5 h.

Figure 11
figure 11

Variation in the hardness of 413 alloys as a function of aging temperature.

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Figure 12
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Electron images of 356 and 413 alloys in the as-cast condition: (a) backscattered-356 alloy, (b) EDS spectrum corresponding to the white circle in (a) showing peaks of Si and Mg, (c) backscattered-413 alloy, (d) EDS spectrum corresponding to the white circle in (c) showing peaks of Al and Cu.

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Figure 13
figure 13

Precipitation observed in alloys aged at 180 °C: (a) Mg2Si in BDS (356) alloy, (b) EDS spectrum corresponding to (a), (c) Al2Cu in CDS (413) alloy, (d) EDS spectrum corresponding to (c).

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Paray et al.30 reported that Sr modification improved the unnotched impact strength of both as-cast and heat-treated A356.0 and 413.0 alloys. Once the initial modification is made, there is a good retention of mechanical properties.31 The work of Li et al.32 showed that the addition of both Mg and Sr can lead to severe segregation of the Al2Cu phase in 319.2 alloys, resulting in the production of large amounts of the coarse block-like phase, compared to the finer eutectic-like form. These segregated block-like Cu-rich phase particles decrease the impact strength of the Al–Si–Cu–Mg alloys.33

Figure 14a reflects the importance of the applied metallurgical parameters on the performance of 356 alloys. In principle, all modified alloys exhibited better impact energies compared to those only grain refined. The presence of porosity in Sr-modified alloys lead to lower hardness in comparison with grain-refined alloys. Best impact results were obtained when the alloy was degassed, grain refined and properly modified. Increasing the hydrogen content and hence the porosity volume fraction resulted in toughness energy values as low as 5 J. Similar observations were noted for the 413 alloys as shown in Figure 14a where most of the Sr-modified alloys exhibited high impact energies coupled with low strength. The scattering in the results is mainly due to the presence of porosity or oxide films.34

Figure 14
figure 14

(a) Hardness–impact energy relationships in 356 alloys. (b) Hardness–impact energy relationships in 413 alloys.

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It is quite useful to examine the fracture profile on sections perpendicular to the fracture surface. This can be carried out using an optical microscope. In this way, the origin of the fracture can be examined to determine if important microstructural abnormalities are present that either caused or contributed to fracture initiation. It is also possible to determine if the fracture path at the initiation site is transgranular or intergranular and to determine whether the fracture path is specific to any phase or constituent present. Some examples of the longitudinal sections of the 413 alloy samples are presented in this section with a view further clarifying the fracture process in these samples.

The optical micrograph presented in Figure 15a shows the polished longitudinal section of the fractured non-modified CD alloy following solution heat treatment, revealing how crack initiation occurs by the cleavage and/or the transgranular fracture of the Si particles. A higher-magnification micrograph, Figure 15b, of the fracture surface corresponding to the arrowed area in (a), shows how the crack propagates further along the fracture surface by cleavage of coarse Si particles. Compared to the micrographs shown in Figure 15a, b, those of Figure 15c, d were obtained from the 413 alloy sample coded CDTES (degassed, Sr-modified and grain refined) corresponding to high energy conditions. Due to the relatively large volume fraction of elastic Al in the matrix, the tips of the dendrites are more rounded than sharp (Figure 15c, arrowed). In this case, the crack propagates through the less ductile Al–Si regions (Figure 15d, arrowed).

Figure 15
figure 15

Microstructures beneath the fracture surface in 413 alloys following solution heat treatment: (a) non-modified CD alloy, low-magnification, (b) high-magnification micrograph of (a), (c) Sr-modified CDTES alloy, low-magnification, (d) high-magnification micrograph of (c).

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Conclusions

Based on the results documented in the present study, the following conclusions may be drawn:

  1. 1.

    Increasing the amount of introduced TiB2 (measured by increasing the amount of Ti) from 0.0075 to 0.25 wt% reduces the alloy grain size by about 90 %.

  2. 2.

    In Sr-modified 413 alloys, the α-Al15(Fe,Mn,Cr)3Si2 intermetallic phase precipitates prior to precipitation of α-Al.

  3. 3.

    A new mechanism based on Ostwald theory was proposed to explain the coarsening in eutectic Si particles in non-modified alloys during solutionizing treatment.

  4. 4.

    Changing the metallurgical parameters (hydrogen content, TiB2 concentration, Sr concentration) has no significant effect on the alloy hardness with respect to the precipitation taking place during aging.

  5. 5.

    Hardening by Al2Cu precipitation (413 alloys) is much less compared to the hardening achieved by Mg2Si (356 alloys).

  6. 6.

    In contrast to hardness, impact energy is found to be sensitive to variations in the applied metallurgical parameters. Best results are obtained for degassed, Sr-modified and grain-refined alloys. Due to porosity formation associated with Sr addition, some scatter in the values is noted.

  7. 7.

    Modifying the 413 alloys with 200 ppm Sr results in a more even distribution of the spheroidized eutectic Si particles and hence improves the alloy toughness.