Metallurgical parameters controlling the microstructure and hardness of Al–Si–Cu–Mg base alloys

Research highlights ► Microstructure and hardness of Al–Si–Cu–Mg type 319 alloys. ► Formation of Mg 2 Si, Q -Al 5 Mg 8 Cu 2 Si 6 , π -Al 8 Mg 3 FeSi 6 , Al 2 Cu phases. ► Effects of alloy type, Mg content, cooling rate, heat treatment (T6, T7). ► Mg, Cu, and high cooling rate (DAS 24 μm) improve hardness in T6 condition. ► Experimental alloys show higher hardness than industrial alloys. Abstract Castings were prepared from both experimental and industrial 319 alloy melts containing 0–0.6 wt% Mg. Test bars were cast in two different cooling rate molds, a star-like permanent mold and an L-shaped permanent mold, with DASs of 24 μm and 50 μm, respectively. The bars were tempered at 180 °C (T6 treatment) and 220 °C (T7 treatment) for 2–48 h. The results showed that Mg content, aging conditions, and cooling rate have a significant effect on the microstructure of both experimental and industrial alloys and, consequently, on the hardness. The addition of Mg resulted in the precipitation of the β -Mg 2 Si, Q -Al 5 Mg 8 Cu 2 Si 6 , π -Al 8 Mg 3 FeSi 6 and of the block-like θ -Al 2 Cu phases. The Mg and Cu, as well as the higher cooling rates improved the hardness values, especially in the T6 heat-treated condition, whereas the addition of Sr decreased these values. Keywords A. Non-ferrous metals and alloys F. Microstructure E. Mechanical 1 introduction The addition of up to 0.5 wt% Mg to molten 319-type alloys leads to the precipitation of a Mg-rich phase, Mg 2 Si, in the form of rounded black particles dotted along the sides of the eutectic Si particles. It then becomes possible to observe a noticeable fragmentation of the eutectic Si ( i.e. modification) and the transformation of a large proportion of the β -Al 5 FeSi iron intermetallic phase into a Chinese-scriptlike phase having a composition close to that of Al 8 Mg 3 FeSi 6 [1] . The Sr–Mg interaction is known to reduce the porosity volume fraction of Al–Si–Cu alloys [2] . The addition of both Mg and Sr can lead to severe segregation of the Al 2 Cu phase in 319.2 alloys, resulting in large amounts of the coarse block-like phase, compared to the finer eutectic-like form [3] . The Mg content also affects the type and total volume fraction of Fe-bearing phases, especially in Be-free alloys. The iron-rich intermetallic phases in the low-Mg alloy are almost exclusively small β -phase (Al 5 FeSi) plates. Large π -phase (Al 8 Mg 3 FeSi 6 ) particles, however, are dominant in the high-Mg alloy, together with a small proportion of the β -phase [4–6] . Cooling rate has a direct effect on the shape, size, and distribution of the microstructural phases, as well as on the aluminum dendrites, eutectic Si, and pore size. Iron intermetallics also vary greatly in size and composition with the cooling rate and magnesium content [7] . The cooling rate is thus the most critical variable in controlling the size and distribution of the intermetallic phases and porosity [8] . The presence of copper in Al–Si–Cu alloys leads to the formation of the Al 2 Cu copper intermetallic. It should be noted that the block-like Al 2 Cu does not dissolve during heat treatment [9] . If iron is also present in the alloy, it forms intermetallic compounds during solidification, including the platelet β -Al 5 FeSi and the script-like α -Al 15 (Mn, Fe) 3 Si 2 phases. In the solidification process, β -Al 5 FeSi platelets are active sites for the nucleation of the Al 2 Cu phase. The addition of Mg leads to the formation of the π -Al 8 Mg 3 FeSi 6 phase. Under cooling rates close to equilibrium conditions for multi-component 3xx alloys and at ∼540 °C, the Mg 2 Si and π -Al 8 Mg 3 FeSi 6 phases begin to precipitate. When the temperature is lowered to 490–530 °C, precipitation of the Al 2 Cu and Q -Al 5 Mg 8 Cu 2 Si 6 phases occurs [10] . A complete dissolution of the iron intermetallics seems to be less likely since the solubility of iron in the aluminum matrix is negligible, thus, the iron intermetallics only transform from one phase to another, though a certain amount of fragmentation and spheroidization may occur [11–13] . Solution heat treatment of Al–Si–Cu–Mg alloys is used to homogenize the alloy, to change the morphology of the interdendritic phases, and to dissolve precipitation-hardening constituents, such as Al 2 Cu, Al 2 MgCu and Mg 2 Si [14] . Yang [15] recommended a solution heat treatment temperature guideline for experimental and industrial 319 alloys to minimize the occurrence of incipient melting. The best combination of strength and ductility is obtained when the as-cast material is solution heat-treated at 515 °C for 8–16 h, followed by a 60 °C warm water quench. A higher solution temperature results in the partial melting of the copper phase at the grain boundaries [16,17] . Aging is the final stage in the heat treatment of cast aluminum alloys. The effect of precipitation heat treatment on mechanical properties is greatly accelerated by heating the quenched Al–Si–Cu alloys in the range of 95–205 °C [18] . The incipient melting of the Al 5 Mg 8 Cu 2 Si 6 and Al 2 Cu phases of the 319 alloys took place when the high-Mg version of 319 was solution heat-treated at temperatures above 505 °C for sufficiently long periods [19] . The hardness of Al-Si alloys depends chiefly on the Mg and Cu content, as well as on the aging conditions. The addition of Mg to Al–Si–Cu alloys containing either β -Fe or α -Fe intermetallics, or both, produces a noticeable increase in hardness at all aging temperatures in both the non-modified and Sr-modified conditions [20,21] . The hardening during aging occurs through the cooperative precipitation of Al 2 Cu and Mg 2 Si phase particles [1] . The iron-intermetallic type/morphology influences the hardness of the heat-treated 319 alloys to some extent. In both the α -Fe and β -Fe intermetallic-containing 319 alloys, the Sr-modified alloys exhibit lower values for hardness, compared to the non-modified alloys [22] . The hardness of the 319 alloys increases with the addition of Mg and the α -Fe intermetallics volume fraction, although it decreases with Sr modification and aging parameters. The contribution of Mg to increasing the hardness of heat-treated 319 alloys containing β -Fe intermetallics is more noticeable than it is for the 319 alloys containing α -Fe intermetallics [20] . The addition of Mg and Cu improves the hardness and strength of the 319 alloys; the Cu-containing alloys, however, are less sensitive to the presence of Sr [22] . The 319 alloy hardness first increases with an increase in aging temperature up to 180 °C, and thereafter decreases as the aging temperature is increased [23] . The current work was undertaken to investigate the effects of Mg content, aging conditions, and cooling rate on the microstructure and hardness of Al-Si–Cu–Mg base alloys. The characteristics of eutectic Si particles are herein examined and explained for both of the cooling rates. The addition of magnesium and copper, as well as the use of higher cooling rates improved the hardness values, especially in the T6 heat-treated condition, whereas the addition of Sr decreased the hardness. 2 Experimental procedures Both industrial-commercial and experimental 319 alloys were used as a basis for the current study. The industrial B319 alloy was received in the form of 12.5 kg ingots. Samplings for chemical analysis were taken from each alloy melt which had previously been prepared. Table 1 shows the average chemical composition of both the experimental and the industrial alloys used. Measured Mg and Sr additions were made to the melt. Strontium was added in the form of the Al–10%Sr master alloy to obtain 200 ppm Sr levels, whereas Mg was added in the form of the pure metal. Prior to casting, the molten metal was degassed for 15 min using pure, dry argon to remove the hydrogen and inclusions. The various 319 alloys were used to prepare castings from which test bars were obtained for hardness testing. With this aim in mind, the molten metal was poured into a permanent mold which had been preheated to 450 °C. Each mold-type was selected for a particular intrinsic quality to prepare hardness test samples and to take metallographic measurements. The star-like mold was used for its high cooling rate, and for which each casting provided ten bars, while the L-shaped mold was chosen for its low cooling rate. After the test bars were cut from the casting, they were machined to the required specifications. The test bars were prepared for each alloy composition and divided into sixteen sets, with one set kept in the as-cast condition and one set solution heat-treated at 495 °C for 8 h, then quenched in 65 °C warm water. Seven sets of bars were solution heat-treated at 495 °C for 8 h, then quenched in warm water at 65 °C and artificially aged at 180 °C for 2, 4, 6, 8, 12, 24, and 48 h, respectively; the remaining seven sets were solution heat-treated at 495 °C for 8 h, quenched in 65 °C water and artificially aged at 220 °C for 2, 4, 6, 8, 12, 24, and 48 h, respectively. For each heat treatment, five test bars were used. Hardness test bars measuring 10 mm × 10 mm × 55 mm were cut from the casting. The specimen surfaces were polished with fine sandpaper to remove any machining marks. 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. From each of these samples prepared for metallographic characterization, two samples measuring 10 × 10 mm, one as-cast and one solution heat-treated, were sectioned off to study each alloy condition. The microstructures of the polished sample surfaces were examined using an optical microscope linked to a Clemex image analysis system. The secondary dendrite arm spacing (SDAS) of both mold samples was measured applying the line intercept method. The eutectic Si particle characteristics, including area, length, aspect ratio, roundness, and density, were measured and quantified. Impact test bars measuring 10 mm × 10 mm × 55 mm were prepared from the casting obtained from the star-like mold. The surfaces of these samples were polished to remove any machining marks; the samples were used in the un-notched condition. All the samples were heat-treated using the same conditions as those applied to the samples used for hardness measurements. A computer-aided instrumented SATEC SI-1 Universal Impact Testing Machine (SATEC Systems Inc., Model SI-1D3) was used to carry out the impact testing. The total absorbed energy was determined, where the average values of the energies obtained from the five samples tested for each alloy condition was taken to represent the impact energy value for that particular condition. A JEOL JXA-8900L electron probe microanalyzer with energy dispersive X-ray spectroscopic (EDS) and wavelength dispersive spectrometric (WDS) facilities was used mainly to detect and analyze the chemical composition of the various intermetallic phases which formed during solidification. The volume fraction of Al 2 Cu was also carried out using the same set-up, while X-ray images of various elements constituting the phases identified were also obtained to determine the distribution of these elements within the phases themselves. In these cases, the samples were cleaned in high purity ethanol using an ultrasonic agitator before submitting them to examination in the electron probe microanalyzer. A number of selected aged samples were examined for precipitation identification using a field emission gun scanning electron microscope (FEGSEM) operating at 2 kV. The polished surfaces were etched using a solution containing 1% HF + 99% H 2 O. 3 Results and discussion 3.1 Microstructural analysis Tables 2–5 summarize silicon particle characteristics for both the experimental and industrial base alloys investigated in the as-cast and solution heat-treated conditions. Table 2 lists Si particle measurements for the D1 experimental Mg-free base alloy using the star-like mold with the higher cooling rate and a DAS of 24 μm. It will be observed that the Si particle area decreases from 29.60 to 5.22 μm 2 , while its length decreases from 17 to 5.16 μm, and its aspect ratio decreases from 4.31 to 2.96, whereas its roundness ratio is observed to increase from 16.5 to 27% after Sr-modification. After solution heat treatment, the particle area decreases to 27.1 μm 2 , the particle length and aspect ratio both decrease to 12.5 μm and 3.92, respectively, although the roundness ratio increases to 26.4% for the non-modified experimental base alloy D1. Also, with respect to the solution heat-treated and Sr-modified base alloy DS1, the particle area and length increase from 5.22 to 9.4 μm 2 and 5.16 to 5.64 μm, respectively, while the aspect ratio decreases from 2.96 to 2.54, as a result of Si particle coarsening. These results are in satisfactory agreement with those of Gruzleski and Closset [7,24] , Hatch [25] , Apelian [26] , and Moustafa et al. [22,27] . Table 3 summarizes the silicon particle characteristics of the same experimental base alloy (alloy LD1) using an L-shaped mold with the lower cooling rate and a DAS of 50 μm. The silicon particle size in this case was observed to increase as a direct effect of the slow solidification rate. Table 4 summarizes the Si particle characteristics of the D7 industrial base alloy (0.3 wt% Mg) using the higher cooling rate star-like mold. Compared to the Si particle characteristics observed for the D1 experimental base alloy, the Si particle characteristics in the D7 alloy increase as a direct effect of chemical composition, namely, the increased Mg content and the impurities present in the industrial alloy. Table 5 summarizes the Si particle characteristics of the LD7 industrial base alloy (0.3 wt% Mg) using the lower cooling rate L-shaped mold. As a result of using this mold, the silicon particle size and characteristics increase as a direct effect of the slow cooling rate. Adding 0.6 wt% Mg to the Sr-modified alloy DS1, i.e. alloy DS6, illustrates how Mg impedes the modification effect produced by Sr, when it is added to the Sr-modified DS6 alloy. Moreover, Si particles are fully modified in some areas, whereas in others they are only partially modified. These observations are in satisfactory agreement with Moustafa et al. [27] and Gruzleski et al. [7] who reported that ∼1 wt% Mg refines the Si phase to a slight degree and also has a negative effect on strontium modification, that is, it changes the microstructure from a well-modified to a partially modified one, due to the precipitation of Mg 2 Sr(Si, Al). When magnesium is added to the alloy, there is a tendency for the copper phase to segregate in localized areas, leading to the formation of the block-like phase, rather than that of the fine eutectic-like Al 2 Cu phase. It will be observed that Mg does not affect the acicular nature of the Si particles, when a comparison is drawn between the Si particle size and morphology of the Mg-free D1 experimental base alloy and those of the D8 industrial alloy containing 0.6 wt% Mg. The addition of Mg to the D6 alloy also causes segregation of the copper phase, leading to its precipitation in a block-like form. It is interesting to note that in the case of the D6 alloy, the Q -Al 5 Mg 8 Cu 2 Si 6 phase also precipitates, displaying a script-like form, rather than the acicular and irregular form commonly observed in the microstructures of the D6 alloys showing examples of the other phases which precipitated during the solidification of this alloy. During solidification, iron, together with other alloying elements, goes partly into solid solution in the matrix and in part forms intermetallic compounds, including the plate-like β -Al 5 FeSi phase, the Chinese-scriptlike α -Al 15 Fe 3 Si 2 phase, and the script-like π -Al 8 Mg 3 FeSi 6 phase [28] . The cooling rate is one of the most significant variables in controlling the size and distribution of the intermetallic phases. When comparing the microstructures in Fig. 1 a and b, the size of these intermetallics increases with a decrease in the cooling rate. These observations are in satisfactory agreement with the ones reported by Samuel et al. [8] . The effects of modification, solution heat treatment, and Mg content with regard to Si particles and intermetallic compounds in the microstructures of (i) the D7 industrial alloys containing 0.3 wt% Mg, (ii) the D8 industrial alloy containing 0.6 wt% Mg, (iii) the LD6 experimental alloy containing 0.6 wt% Mg, (iv) the LD7 industrial alloys containing 0.3 wt% Mg, and (v) the LD8 industrial alloy containing 0.6 wt% Mg are all similar to those observed for the D1 and D6 experimental alloys. In addition to the copper and iron intermetallic compounds, the Mg 2 Si phase may be observed in the as-cast microstructures of both the experimental and industrial alloys containing up to 0.6 wt% Mg. In Fig. 2 , the presence of ultrafine Si particles within the Al–Al 2 Cu eutectic structure may also be observed, tending to precipitate at the end of the solidification process; the encircled area 2 in Fig. 2 b shows these ultrafine Si particles. Remnants of undissolved Al 2 Cu after solution heat treatment may also be noted in Fig. 2 a. In addition, examples of the ultrafine β -phase platelets observed in the microstructure of the solution heat-treated sample of the LD7 alloys are shown in the encircled areas numbered 1 and 2 in Fig. 2 b. Fine remnants of mostly dissolved π -Al 8 Mg 3 FeSi 6 phase are shown in the encircled area 3 of this figure also. The large standard deviation observed is due the presence of extremely large Si particles indicated by the thick solid arrow shown in Fig. 2 b adjacent to relatively fine ones indicated by the broken arrow. 3.2 Hardness Table 6 summarizes the hardness values of the experimental and industrial as-cast 319 alloys for non-modified and Sr-modified conditions, for samples obtained out of both the star-like and L-shaped molds (coded D, DS and LD, LDS, respectively). It was found that the hardness increases with an increase in both the Mg content and the cooling rate, although it decreases with Sr-modification and the use of a low cooling rate for both the experimental and industrial alloys. These observations are in good agreement with the work of Tash et al. [20,21] and Moustafa et al. [22] . The effects of Mg addition on the hardness of T6 heat-treated non-modified and Sr-modified experimental 319 alloys were investigated as a function of aging time, and shown in Figs. 3 and 4 for samples obtained from the star-like and L-shaped molds, respectively. Fig. 3 a shows the Brinell hardness values for non-modified experimental alloys, including the highest recorded Brinell hardness value of all the alloys tested here, namely, 143.52 BHN for the non-modified experimental alloy D4 containing 0.3 wt% Mg; Fig. 3 b, on the other hand, shows the Brinell hardness values for the Sr-modified alloys alone. These two groups of curves were obtained using the higher cooling rate star-like mold samples. It was found that alloy hardness values could be increased by increasing the Mg content. Thus, regular and successive increases of the Mg content in the D1 Mg-free base alloy, for example, by 0.1 wt% created the D2 alloy; by 0.2 wt% created the D3 alloy; by 0.3 wt% created the D4 alloy; by 0.4 wt% created the D5 alloy; and by 0.6 wt% created the D6 alloy, thereby increasing the alloy hardness values by ∼32%, ∼43%, ∼46%, ∼42%, and ∼42%, respectively. It was also observed that an increase in the Mg content, up to 0.4 wt%, at different aging times produces a positive effect on hardness, indicating that hardening is due to Mg 2 Si precipitation, in addition to the precipitation of Al 2 Cu. The effects of Mg in aged alloys are similar to those observed in the as-cast alloys. Aging of these Mg-containing 319 alloys at 180 °C in T6 heat treatment conditions for up to 48 h produces a sharp rise in hardness values during the first two hours of aging, followed by a broad peak or plateau spread between 2 and 12 h, as well as a noticeable period of over-aging after 12 h. Modification with 200 ppm Sr has a negative effect on hardness, as shown in Fig. 3 b. The effects of Mg-content on the Sr-modified alloys are similar to those observed for the non-modified alloys. Again, these observations are in satisfactory agreement with those reported by Tash et al. [20,21] and Moustafa et al. [22] . The lower cooling rate L-shaped mold was used to investigate the effects of cooling rate on the hardness of the experimental 319 alloys, as shown in Fig. 4 . When compared to Fig. 3 , it will be seen that the slow cooling rate decreases the hardness in both the non-modified and Sr-modified T6 heat-treated alloys. The effects of Mg addition on the hardness of the non-modified and Sr-modified experimental 319 alloys as a function of aging time were also investigated under T7 heat treatment conditions, and are shown in Figs. 5 and 6 for samples obtained from the star-like and L-shaped molds, respectively. Similar observations were noted in the case of the T6 heat-treated alloys, as shown in Figs. 3 and 4 . Aging of these Mg-containing experimental 319 alloys at 220 °C in T7 heat treatment conditions for up to 48 h produces a sharp rise in hardness during the first two hours of aging, followed by an aging peak and a noticeable period of over-aging, after 2 h of aging. These observations are in satisfactory agreement with the work of Tash et al. [20,21] . The effects of the addition of Mg on the hardness of the non-modified and Sr-modified industrial 319 alloys as a function of aging time were investigated under both T6 and T7 heat treatment conditions for samples obtained from the star-like and L-shaped molds, respectively. The results are provided in Figs. 7 and 8 . It was observed that both the Mg-containing industrial and experimental 319 alloys exhibited similar behavior with respect to the influence of Mg-content, Sr-modification, aging temperature (180 °C vs. 220 °C), and cooling rate (star-like mold vs L-shaped mold), as shown in Figs. 3–8 . The hardness curves shown in Figs. 3–8 exhibit more than one peak or a wavy form with aging time, resulting from the presence of several hardening phases, including θ -Al 2 Cu, β -Mg 2 Si, and Q -Al 5 Mg 8 Si 6 Cu 2 which contribute to the precipitation hardening of the alloys. Similar observations for this wavy form were recorded and explained for Al–Si–Cu–Mg 380 alloys by Morin [29] , on the basis of two different types of structures which are likely be obtained during the aging treatment of 319 alloys. The negative effects of Sr-modification on the properties of 319 alloys involve the segregation of the brittle, block-like Al 2 Cu in areas away from the modified Si particles. This block-like Al 2 Cu as well as the Q -Al 5 Mg 8 Cu 2 Si 6 phase are known to be insoluble. Table 7 illustrates these undissolved Cu-rich phases which control the properties and the fracture behavior of the alloys investigated. Such an observation is in satisfactory agreement with data reported by Paray et al. [9,30] . Table 7 summarizes the volume fraction of Al 2 Cu (%) for both the as-cast and solution heat-treated experimental 319 alloy in non-modified and Sr-modified conditions, observed in the star-like and L-shaped mold samples. The volume fraction of Al 2 Cu (%) for the Sr-modified alloy is higher than it is for the non-modified alloy in both the as-cast and solution heat-treated conditions. These volume fraction values confirm: (i) the negative effects of Sr modification on 319 alloys containing ∼3.5% Cu, (ii) the decrease in hardness values of Sr-modified 319 alloys in both the as-cast and solution heat-treated conditions, and (iii) the hardness data summarized in Table 6 and those plotted in Figs. 3–8 . These observations are all in satisfactory agreement with those reported by Tash et al. [20,21] , Moustafa et al. [22] , and Paray et al. [9,30] . The effects of aging temperature on the hardness of the non-modified experimental 319 alloys as a function of Mg content were also investigated for both molds and the results are shown in Fig. 9 . It was found that alloy hardness values could be increased by increasing the Mg content. Thus, increasing the Mg content up to 0.3 wt% increased hardness values noticeably. Increasing Mg content beyond 0.3 wt% up i.e. 0.6 wt% did not lead to a noticeable increase in the alloy hardness. The decline observed in the hardness values beyond 0.3 wt% Mg may be attributed to the increase in the volume fraction of both π -Al 8 Mg 3 FeSi 6 and Q -Al 5 Mg 8 Cu 2 Si 6 phases. Increasing aging temperature up to 180 °C increased hardness values which thereafter decreased for both the mold samples as the aging temperature was increased to 220 °C. Similar observations were recorded for increasing Mg content and aging temperature in Sr-modified samples and industrial 319 alloys. These observations are all in satisfactory agreement with those reported by Tash et al. [20,21] , Moustafa et al. [22] , and Mohamed [23] . 3.3 Phase precipitation In order to arrive at a better understanding of the nature of the precipitate phase that takes place during the course of aging, some aged samples were selected from the alloys investigated in the non-modified and Sr-modified conditions. The samples were examined using Field Emission Gun Scanning Electron Microscopy (FEGSEM). The as-quenched sample revealed noticeable precipitation even at magnifications as high as 60,000×. Precipitation of fine particles was observed when the sample was aged for 2 h at 180 °C, as shown in Fig. 10 a. Increasing the aging time resulted in a significant increase in the density of the precipitated phase with no change in its morphology, as displayed in Fig. 10 b corresponding to the maximum hardness achieved by this alloy. It is interesting to note that aging for 2 h at 220 °C resulted in a density of precipitates which was more or less similar to the one shown in Fig. 10 b, as is evident in Fig. 10 c. Increasing the aging time at this temperature, Fig. 10 d indicated the commencement of the change in the morphology of the precipitate from spherical to a mixture of spherical and elongated particles indicating the tendency of the precipitates to change from semi-incoherent to stable incoherent Al 2 Cu, causing a clear reduction in the alloy hardness, as described in Fig. 5 . Fig. 11 shows the EDS spectra produced from Fig. 10 d exhibiting the presence of both Mg 2 Si and Al 2 Cu precipitates. Modification with Sr seems to have relatively little effect on the precipitated phase in terms of density and morphology with respect to aging temperature and aging time, as displayed in Fig. 12 . Note the presence of a precipitate free zone along the grain boundary in Fig. 12 a. 3.4 Impact testing In order to arrive at a better understanding of the mechanical properties and to confirm the hardness values of 319 alloys, the impact toughness of both experimental and industrial 319 alloys were assessed using Charpy instrumented impact testing for the samples obtained from the star-like mold. All the samples were heat-treated for the same conditions as the hardness samples. Figs. 13 and 14 display the effects of Mg addition on the impact energy values of non-modified experimental and industrial 319 alloys as a function of aging time in the T6 and T7 heat-treated condition, respectively. It was found that alloy toughness, or the total impact energy, decreased upon increasing the Mg content. It was also observed that an increase in the Mg content from 0.3 wt% up to 0.6 wt% has more or less no significant effect on the variation in the alloy toughness similar to that reported for the hardness. The aging of Mg-containing 319 alloys under T6 (180 °C) and T7 (220 °C) heat treatment conditions for up to 48 h produced a sharp decrease in impact toughness during the first two hours of aging, followed by a broad valley or plateau spread between 2 and 24 h as well as a noticeable period of over-aging beyond 24 h. The T7 heat-treated alloys displayed higher impact toughness values than those which were T6 heat-treated. In both T6 and T7 heat-treated conditions, the experimental alloys demonstrated higher values of impact energy than the industrial alloys. These observations are presented elsewhere [31] and confirm the observations relating to hardness. The impact toughness curves shown in Figs. 13 and 14 exhibit more than one peak or a wavy form with aging time, resulting from the presence of several hardening phases as reported and explained earlier in the hardness section for the hardness curves [29] . 4 Conclusions (1) The addition of Mg to Fe-containing 319 alloys results in the precipitation of the Mg 2 Si, Q -Al 5 Mg 8 Cu 2 Si 6 and π -Al 8 Mg 3 FeSi 6 phases, where the Q - and π -phases appear in a script-like form rather than as irregularly-shaped particles. Magnesium has a slight refining effect on the Si phase and a negative effect on Sr modification, since a change in the microstructure may be observed corresponding to that between a well-modified and a partially modified alloy. (2) Magnesium and copper improve the hardness of the alloy samples tested, especially in the T6 heat-treated condition. The higher cooling rate also produces an increase in the hardness values, especially for the non-modified Mg-containing alloys. (3) The addition of Sr, however, decreases the hardness of both the Mg-free and Mg-containing alloys; this situation arises most likely from a delaying of Mg 2 Si precipitation during the aging process of the Mg-containing alloys. (4) The T7 heat-treated alloys display lower hardness values than those which were T6 heat-treated. In both T6 and T7 heat-treated conditions and for both the non-modified and Sr-modified alloys, the experimental alloys demonstrate higher hardness values than the industrial alloys. (5) The aging of the Mg-containing 319 alloys at 180 °C in T6 heat-treated conditions, as well as of the Mg-containing experimental 319 alloys at 220 °C in T7 heat-treated conditions, produce a sharp increase in hardness during the first two hours of aging. At 180 °C, this initial increase is followed by a plateau from 2 to 12 h, with a noticeable period of over-aging beyond 12 h. At 220 °C, the initial rise is followed by an aging peak and a noticeable period of over-aging beyond 2 h. References [1] F.H. Samuel P. Ouellet A.M. Samuel H.W. Doty Effect of Mg and Sr additions on the formation of intermetallics in Al-6 wt Pct Si–3.5 wt Pct Cu-(0.45)–(0.8) Wt Pct Fe 319–type alloys Metall Mater Trans A 29A 1998 2871 2884 [2] N. Roy A.M. Samuel F.H. Samuel Porosity formation in Al–9%Si–3%Cu alloy systems: metallographic observations Metall Mater Trans A 27 1996 415 429 [3] A.M. Samuel P. Ouellet F.H. Samuel H.W. Doty Microstructural interpretation of thermal analysis of commercial 319 Al alloy with Mg and Sr additions AFS Trans 105 1997 951 962 [4] J. Barresi M.J. Kerr H. Wang M.J. Couper Effect of magnesium, iron and cooling rate on mechanical properties of Al–7Si–Mg foundry alloys AFS Trans 117 2000 563 570 [5] C.H. Caceres C.J. Davidson J.R. Griffiths Q.G. 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