1 Introduction

In-situ aluminum matrix composites (AMCs), in which reinforcement phases are formed in Al matrices during processing, have received considerable attention for high-performance applications owing to their low densities, excellent mechanical properties, and low production costs [1]. Reactions between Al melts and salts, gases, and powders have been used to form various reinforcement phases (e.g., TiB2, TiC, and Al2O3) showing good interfaces with Al matrices and homogeneous reinforcement-particle distributions [1,2,3,4]. Reinforcement phases also have been produced by the solidification of alloyed melts [5].

The Al–Mg2Si system is an in situ AMC in which the reinforcement phase is naturally formed during solidification of alloyed melt and is a promising material owing to the low density, high hardness, and good physical properties of the Mg2Si phase as well as low production costs [6,7,8,9,10,11]. In addition to modifying the main reinforcement Mg2Si phase [6, 7], many investigators have improved the mechanical properties of Al–Mg2Si-based AMCs by forming additional reinforcement phases through the addition of alloying elements such as Cu [8], Zn [9], Ti [10], and Zr [11].

Recently, efforts have been made to explore the chemical compositions of other in situ AMCs [12,13,14,15]. Yang et al. [12] used suction casting to produce several multiphase-reinforced in situ AMCs. They reported that Al80Li5Mg5Zn5Sn5 and Al80Li5Mg5Zn5Cu5 showed various reinforcement phases (e.g., Mg2Sn, Li2MgSn, and Al2Cu) and excellent compressive strengths (in the range 836–879 MPa) and fracture strains (in the range 16%–17%). We used conventional mold casting to produce multiphase-reinforced Al70Mg10Si10Cu5Zn5 [13] and Al68Mg10Si10Cu5Zn5Ni2 [14, 15]. Such in situ AMCs contained more than 30% reinforcement phases and exhibited compressive properties superior to those of conventional Al alloys (A356, A390, and piston alloy) up to 200 °C [13]. We also found that the compressive strength of Al68Mg10Si10Cu5Zn5Ni2 was improved by low-temperature (i.e., < 120 °C) aging owing to the formation of fine clusters and precipitates [14, 15].

There are numerous combinations of alloying element types and contents; thus, the effects of alloying elements on the microstructures and mechanical properties of AMCs still need to be explored to develop advanced in situ AMCs showing the desired reinforcement phases. Therefore, we investigated the microstructures and mechanical properties of seven in situ AMCs prepared with different combinations and contents of Li, Mg, Si, Cu, Zn, Sn, and Ni. Furthermore, the effects of the annealing temperatures and cooling rates on the microstructures and mechanical properties of as-cast AMC were examined.

2 Experimental Procedures

Seven alloy ingots were produced with the nominal chemical compositions (at%) listed in Table 1. For Alloys 1–5, 0.5 kg of each melt was produced in a high-frequency induction-melting furnace, using high-purity (i.e., 99.9 wt%) metals (i.e., Al, Mg, Zn, and Sn) and various master alloys (i.e., Al–30Cu, Al–15Si, and Al–15Li). The alloy melts were poured at 700 °C (i.e., above their corresponding liquidus temperatures (TL) in the range 556–684 °C calculated using Thermo-Calc software [16] and the TCAL3 database [see Table 1)] into a copper cylindrical mold (Φ 20 × 50 mm), which had been preheated in the range 100–110 °C.

Table 1 Chemical compositions, liquidus/solidus temperatures, and densities of as-cast alloys
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Meanwhile, 1.5-kg melts of Alloys 6 and 7 were produced in an induction-melting furnace, using high-purity metals (i.e., Al, Mg, and Zn) and master alloys (i.e., Al–50Cu, Al–25Si, and Al–10Ni) and were poured at 700 and 850 °C (considering the high TL (= 754 °C) of Alloy 7), respectively, into a copper book mold (245 × 200 × 70 mm3), which had been preheated in the range 130–140 °C. A protective gas mixture (SF6:CO2 = 1:10) was flowed at 660 mL/min over the Alloy 7 melt to prevent oxidation. Samples of the as-cast Alloy 6 ingot were annealed in the range 100–440 °C for 2 h and then air-cooled (AC) to room temperature (RT). Additionally, the sample annealed at 440 °C was water-quenched (WQ) and then naturally aged (NA) at RT.

The chemical compositions of the seven alloy ingots were examined using inductively coupled plasma optical emission spectrometry (ICP-OES, Thermo Fisher Scientific, ICAP-6500), and the results are listed in Table 1. Note that the measured chemical compositions of Li (i.e., in the range 0.4–0.6 wt%) were lower than the corresponding nominal ones (i.e., in the range 1.1–1.3 wt%) owing to the high reactivity of Li. The density of each alloy ingot was measured using an analytical balance (Mettler Toledo, AG285) according to Archimedes’ method.

Specimens for the microstructural observations were machined from the centers of the cylindrical and book-shaped ingots. The specimens were mechanically polished, and their microstructures were observed using an optical microscope (OM, Nikon, MA200). The area fraction of the secondary phases and secondary dendrite arm spacing (SDAS) were measured from three OM images taken at 500 × magnifications using an image analyzer (IMT, i-Solution). To identify the secondary phases in the as-cast ingots, the samples were analyzed using an X-ray diffractometer (XRD, Rigaku, D/Max-2500 V/PC) equipped with a Cu target (λ = 1.5406 Å). The samples were scanned at 1°/min in the Bragg angle (i.e., 2θ) range 20°–100°. The types, morphologies, and chemical compositions of the secondary phases were further characterized using a scanning electron microscope (SEM, ZEISS, MERLIN) and an energy-dispersive X-ray spectroscope (EDS, Oxford, X-Max 100) operated at 5 kV.

The hardness of the Al matrices, secondary phases, and overall composites was measured using a Vickers hardness tester (Mitutoyo, HM-122) operated under loads of 0.025, 0.1 or 0.3, and 0.5 kg, respectively. Compression tests were conducted on two or three cylindrical specimens (Φ 8 × 12 mm) per alloy at 25 ± 2 °C, using an Instron 5982 universal testing machine. The specimens were strained at 5 × 10−4 s−1.

3 Results and Discussion

3.1 Thermodynamic Calculations

Figure 1a, b represent the mole fractions (calculated using Thermo-Calc software and the measured alloy chemical compositions) of secondary phases formed in the seven alloys solidified under nonequilibrium (i.e., Scheil) and equilibrium conditions. The theoretically predicted total mole fractions of the secondary phases ranged from 11.2% to 51.9%, depending on alloy composition. The total phase fractions of the alloys solidified under non-equilibrium and equilibrium conditions were similar despite the alloys showing some different secondary phases.

Fig. 1
figure 1

Mole fractions of secondary phases formed in alloys solidified under a nonequilibrium (i.e., Scheil) and b equilibrium conditions. c Equilibrium mole fractions of secondary phases in alloys at 25 °C

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Various secondary phases were expected to form in the alloys during solidification, owing to the high contents and various types of alloying elements. Alloy 1 showed θ-Al2Cu, Mg2Sn, Si, and Sn, while Alloy 2 showed θ-Al2Cu, Q-Al5Cu2Mg8Si6, Mg2Si, Si, and V-Mg2Zn11. The main secondary phases in Alloy 3 were θ-Al2Cu, AlLiSi, Q-Al5Cu2Mg8Si6, and Mg2Si while those in Alloy 4 were V-Mg2Zn11, η-Mg(Zn,Cu,Al)2, Al2CuLi, and AlCuLi. Alloy 5 showed the lowest total mole fractions of Mg2Sn, Sn, Zn, and AlLi. Note that Alloys 1 and 5 showed Sn and Mg2Sn phases formed during the final stage of low-temperature solidification (see solidus temperatures (TS) in Table 1).

Alloy 6 showed all the secondary phases in Alloy 2 owing to their similar alloying elements. Alloy 6, however, showed higher mole fractions of Mg2Si and Si as well as additional Ni-containing phases such as δ-Al3CuNi and γ-Al7Cu4Ni. Therefore, Alloy 6 showed a higher total mole fraction than Alloy 2. Alloy 7 exhibited the highest total mole fraction of Mg2Si, η-Mg(Zn,Cu,Al)2, S-Al2CuMg, and V-Mg2Zn11. Note that Alloys 3, 6, and 7 showed primary phases, which had formed before the Al matrices, in mole fractions of 6.7% (AlLiSi), 8.1% (δ-Al3CuNi, Mg2Si, and Si), and 27.4% (Mg2Si), respectively.

Figure 1c shows the calculated equilibrium mole fractions of the secondary phases in the alloys at 25 °C. The total mole fractions of the secondary phases in the alloys at 25 °C were higher than those of the secondary phases formed during alloy solidification, most likely owing to changes in phase stability and to the formation of precipitates at low temperatures.

3.2 Microstructures

Figure 2 shows XRD patterns indicating various secondary phases in the as-cast alloys. The XRD peaks were identified using the secondary-phase crystallographic data listed in Table S1. The secondary phases detected in the XRD patterns were consistent with the main phases expected to form during alloy solidification, despite some discrepancies between the theoretically predicted and experimentally obtained minor phases. The alloy microstructures were observed using both OM and SEM to further investigate the secondary phases in the as-cast alloys.

Fig. 2
figure 2

XRD patterns of as-cast Alloys: a 1, b 2, c 3, d 4, e 5, f 6 and g 7

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Figure 3 shows both the OM and SEM images of the microstructures of as-cast Alloys 1 and 2. The secondary phases were distinguished based on their chemical compositions (see Table S2). Alloy 1 exhibited block-like Sn and Mg2Sn phases, eutectic colonies of θ-Al2Cu and Si phases, and a small amount of the Q-Al5Cu2Mg8Si6 phase (Fig. 3a, b). Meanwhile, Alloy 2 exhibited block- or Chinese-script-like Mg2Si, eutectic Si, θ-Al2Cu, Q-Al5Cu2Mg8Si6, and Zn phases (Fig. 3c, d). The Mg2Sn and Sn phases in Alloy 1 and θ-Al2Cu and Zn phases in Alloy 2 appeared attached together because they had sequentially formed during alloy solidification. Most of the secondary phases in Alloys 1 and 2 were present in the inter-dendritic regions. This means that the secondary phases had formed with Al matrices during alloy solidification, which is consistent with the theoretical prediction that Alloys 1 and 2 would not show any primary phases.

Fig. 3
figure 3

OM and SEM images of as-cast Alloys a, b 1 and c, d 2

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Figure 4 shows the OM and SEM images of the microstructures of Alloys 3, 4, and 5, i.e., the three Li-containing alloys. The inserts of the SEM images show the corresponding EDS maps of the constituent elements in the Li-containing phases. Since Li could not be detected by SEM–EDS, only the element mappings of the other constituent elements in the Li-containing phases are displayed. The microstructure of Alloy 3 consisted of AlLiSi, θ-Al2Cu, Mg2Si, and Q-Al5Cu2Mg8Si6 phases, as shown in Fig. 4a, b. The AlLiSi phase showed two different morphologies: a coarse, block-like primary phase initially formed during alloy solidification and a smaller phase (attached to the θ-Al2Cu phase) formed during the final stage of alloy solidification.

Fig. 4
figure 4

OM and SEM images of as-cast Alloys a, b 3, c, d 4, and e, f 5. Inserts of SEM images show corresponding EDS maps of constituent elements in Li-containing phases

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The microstructure of Alloy 4 exhibited η-Mg(Zn,Cu,Al)2, V-Mg2Zn11, and Cu-containing phases, as shown in Fig. 4c, d. The XRD pattern (Fig. 2d) confirmed that the Cu-containing phases in Alloy 4 were θ-Al2Cu and Al2CuLi phases. Figure 4e, f show that Alloy 5 consisted of AlLi, Mg2Sn, Sn, and Sn(Zn) phases. The secondary phases identified in Alloys 4 and 5 were somewhat different from those previously reported for the same alloys [12], i.e., the main phases reported in Al80Li5Mg5Cu5Zn5 (Alloy 4) and Al80Li5Mg5Zn5Sn5 (Alloy 5) were Al2Cu + AlCu3 and Mg2Sn/Li2MgSn + Sn, respectively. This discrepancy is believed to be due to the reduced Li and Mg contents of the two alloys produced in this study.

Alloys 6 (Fig. 5a, b) and 7 (Fig. 5c, d) both showed a dendritic primary Mg2Si phase. The higher-Mg-content alloy (i.e., Alloy 7) showed both a larger primary Mg2Si phase and a higher Mg2Si area fraction than Alloy 6. In addition to the primary Mg2Si phase, Alloy 6 contained δ-Al3CuNi, θ-Al2Cu, Q-Al5Cu2Mg8Si6, eutectic Si, and Zn phases while Alloy 7 contained a large amount of eutectic η-Mg(Zn,Cu,Al)2. The microstructures of the as-cast alloys (Figs. 3, 4 and 5) were very consistent with the corresponding XRD patterns (Fig. 2) and were comparable to the theoretically predicted microstructures (Fig. 1). The measured area fractions of the secondary phases in the as-cast alloys were also comparable to the theoretically predicted ones, as shown in Fig. 6a.

Fig. 5
figure 5

OM and SEM images of as-cast Alloys a, b 6 and c, d 7

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Fig. 6
figure 6

a Calculated mole fractions and measured area fractions of secondary phases and b Δχ-δ plots of as-cast alloys. Three constituent phase regions reported in multicomponent alloys [12] are indicated in b

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The phase stability of the as-cast alloys was also estimated using the differences in atomic size (δ) and Pauling electronegativity (Δχ) [12, 17] as follows:

$$\delta = \sqrt {\mathop \sum \limits_{i = 1}^{n} c_{i} \left( {1 - \frac{{r_{i} }}{{\bar{r}}}} \right)^{2} }$$
(1)
$$\Delta \chi = \sqrt {\mathop \sum \limits_{i = 1}^{n} c_{i} \left( {\chi_{i} - \bar{\chi }} \right)^{2} } ,$$
(2)

where ci and ri are the atomic fraction and atomic radius of the ith element, respectively, \(\bar{r}\) is the average atomic radius, \(\chi_{i}\) is the Pauling electronegativity of the ith component, and \(\bar{\chi }\) is the average Pauling negativity. These parameters can be found in the literature [18]. The Δχ and δ values of the as-cast alloys are plotted in Fig. 6b. According to a previous study [12], the constituent phases of multicomponent alloys are usually categorized as solid–solution (SS), intermetallic compounds (IMC), and SS + IMC hybrids corresponding to the region in which they are located in Δχ–δ maps (see Fig. 6b). Alloys 1–5 were all located in either the SS or overlapped SS/SS + IMC region. Alloy 6 was located in the center of the SS + IMC region, and Alloy 7 was located in the overlapped SS + IMC/IMC region. Overall, the Δχ–δ map was consistent with the microstructures of the as-cast alloys, considering that Δχ and δ both moved to the IMC region with decreasing the area fraction of the Al matrix.

The Al matrices of the as-cast alloys contained precipitate-forming elements, as indicated by their measured chemical compositions (see Table 2). Although Li could not be detected by SEM–EDS, the Al matrices of the Li-containing alloys were believed to contain at least some Li. The formations of the Zn-containing η-Mg(Zn,Cu,Al)2 (i.e., in Alloys 4 and 7) and Sn(Zn) (i.e., in Alloy 5) phases had reduced the Zn contents of the Al matrices. The chemical composition of Mg was much higher in the Al matrix of Alloy 7 than in those of Alloys 4 and 5, owing to the high Mg content of Alloy 7. SDAS was measured in the ranges 8–14 and 9–13 μm in Alloys 1–5 and Alloys 6 and 7, respectively, suggesting that the solidification cooling rates of both groups of alloys were similar [19]; thus, enabling us to directly compare the alloy mechanical properties.

Table 2 Chemical compositions of Al matrices of as-cast alloys (at%)
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3.3 Mechanical Properties

Figure 7a shows the hardness values of the matrices and composites of the as-cast alloys. Matrix hardness increased from 114 to 135 HV with increasing difference in mole fractions of secondary phases between 25 °C and TS (see Fig. 7b), implying that solutes or precipitates contributed to the matrix strengthening. This finding was also supported by the relationships between matrix hardness and measured chemical compositions, showing that harder Al matrices contained more alloying elements (Table 2).

Fig. 7
figure 7

a Hardness values of Al matrices and composites for as-cast alloys. Relationships between b matrix hardness and calculated phase fractions, c matrix hardness and composite hardness, and d composite hardness and measured phase fractions of as-cast alloys

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Matrix and composite hardness were positively correlated, as shown in Fig. 7c. Note that the slope (i.e., 3.5) of the best-fit line through the data points plotted for the relationship between composite and matrix hardness was much greater than 1 (indicated by the dotted line). This means that the hardness of the composite was significantly affected by the secondary phases rather than the Al matrix. Composite hardness increased from 123 to 195 HV with increasing total area fraction of the secondary phases from 26% to 41% (see Fig. 7d) because most of the secondary phases, except the Sn-containing ones, were harder than the Al matrix. The hardness values measured for Mg2Sn/Sn, Mg2Sn, Mg2Si, AlLiSi, and θ-Al2Cu/δ-Al3CuNi were 64 ± 13, 114 ± 13, 354 ± 24, 625 ± 28, and 742 ± 52 HV, respectively.

Interestingly, Alloy 7 (showing a secondary-phase area fraction of 58%) exhibited a composite hardness similar to that of Alloy 6 (showing a secondary-phase area fraction of 41%). This is because Alloy 7 contained a large amount of Mg2Si, which is softer than the major reinforcing phases (e.g., θ-Al2Cu and δ-Al3CuNi) in Alloy 6. Further, note that the composite hardness values of Sn-containing Alloys 1 and 5 were significantly lower than those of the other alloys because Mg2Sn/Sn-induced softening had complemented the strengthening induced by the other hard secondary phases.

Figure 8a shows the compressive stress–strain curves of the as-cast alloys, and the insert shows low-magnification photographs of fractured specimens of Alloys 2 and 7. The fractured specimens of Alloys 1–6 showed cracks in the direction 45° to the compression axis. Alloy 7, on the other hand, fractured prematurely and showed various cracks (deviating from the direction 45° to the compression axis) and fragmentation associated with rapid crack initiation and propagation [20]. The maximum compressive stress values increased with increasing composite hardness for all the alloys except Alloy 7, which had prematurely fractured during compression (Fig. 8b). The fracture strain decreased with increasing area fraction of the secondary phases for all the alloys, as shown in Fig. 8c. In particular, the alloys containing primary phases (i.e., Alloys 3, 6, and 7) exhibited lower fracture strains than the other alloys because their coarse primary phases readily acted as sites for initiating and propagating cracks [21]. Clearly, cracks had appeared adjacent to coarse primary Mg2Si phases in Alloy 7, as shown in the inset of Fig. 8c.

Fig. 8
figure 8

a Compressive stress–strain curves of as-cast alloys and relationships between b maximum stress and composite hardness and c fracture strain and area fraction of secondary phase in as-cast alloys. d Specific strengths of as-cast alloys

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Figure 8d shows the specific strengths of the as-cast alloys, obtained by measuring the alloy maximum compressive stresses and densities. Alloys 2, 3, 4, and 6 showed specific strengths ranging from 216 to 231 N m/g. Among them, Alloy 6 exhibited the highest specific strength owing to pronounced strengthening by the Al matrix and secondary phases. Alloys 1, 5, and 7 showed lower specific strengths of 161, 185, and 156 N m/g, respectively; the first two owing to soft/heavy Mg2Sn and Sn phases and the last one owing to a large amount of the coarse primary Mg2Si phase in the alloy, thereby leading to premature fracturing.

The results indicated that matrix and secondary-phase strengthening should be controlled simultaneously during alloy solidification to improve the mechanical properties of as-cast alloys. Further, the results implied that as-cast alloys could be thermally annealed to improve their mechanical properties by mitigating the negative effects of Al matrices and secondary phases on alloy specific strengths. Hence, the effects of thermal annealing on alloy microstructures and mechanical properties were further investigated using Alloy 6 because it exhibited the highest specific strength among the alloys.

3.4 Effects of Thermal Annealing on Microstructure and Mechanical Properties of Alloy 6

Figure 9a shows the hardness of Alloy 6 samples as-cast or thermally annealed in the range 100–440 °C for 2 h and subsequently AC, WQ, or WQ + NA at RT. For AC specimens, the matrix and composite hardness both decreased with increasing annealing temperature up to 200 °C and then both increased with increasing annealing temperature above 200 °C. The WQ specimen annealed at 440 °C exhibited the lowest hardness. However, naturally aging the WQ specimen at RT for 300 h (i.e., WQ + NA) dramatically increased its hardness. Figure 9b shows the compressive properties of Alloy 6 samples as-cast or thermally annealed at either 200 or 440 °C for 2 h and subsequently AC, WQ, or WQ + NA at RT. Annealing-induced changes in the maximum stress exhibited a trend similar to that of annealing-induced changes in hardness, while all the annealed specimens showed higher fracture strains than the as-cast alloy.

Fig. 9
figure 9

a Hardness and b compressive properties of as-cast or thermally annealed Alloy 6

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Figure 10a, b show the equilibrium volume fraction of secondary phases calculated using the chemical composition of alloy and Al matrix, respectively. These calculations provide the thermal stability of micrometer-sized secondary phases (Fig. 10a) and nanometer-sized precipitates in the Al matrix (Fig. 10b). Decreases in composite hardness and compressive stress up to 200 °C mainly originated from matrix softening because the secondary phases had hardly changed at all at such low temperatures (see Fig. 11a). As-cast Al70Mg10Si10Cu5Zn5 with a chemical composition similar to that of Alloy 6 has been reported to contain nano-sized precipitates such as θ′-Al2Cu, Q-Al5Cu2Mg8Si6, and Zn [13]. Thus, matrix softening in Alloy 6 is thought to be due to precipitate coarsening during low-temperature annealing.

Fig. 10
figure 10

Equilibrium volume fraction of secondary phases of Alloy 6 calculated using the chemical composition of a alloy and b Al matrix

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Fig. 11
figure 11

OM images of Alloy 6 samples annealed at a 200 °C and b 440 °C

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Above 200 °C, the unstable Mg2Si phase transformed into the stable Q-Al5Cu2Mg8Si6 one while the secondary phases were spheroidizing (see Fig. 11b). Any changes in secondary phases above 200 °C would reduce composite hardness and compressive stress while increasing fracture strain. Thus, increases in composite hardness and the maximum compressive stress in the range 200–440 °C were due to significant matrix strengthening, which was strong enough to overcome any secondary-phase-induced softening. It is thought that the dissolution of thermally unstable secondary phases and precipitates (e.g., θ-Al2Cu, V-Mg2Zn11, Zn) during annealing (Fig. 10) enables the formation of precipitates during AC, thereby increasing matrix hardness. Matrix strengthening became more pronounced at higher annealing temperatures because more precipitates had formed in the highly supersaturated Al matrix during cooling.

Decreases in the hardness and compressive stress of the WQ specimen probably originated from its soft supersaturated Al matrix, which had allowed fine precipitates to form at RT. A previous study [15] reported that precipitates formed in ellipsoidal Guinier–Preston zones (e.g., Zn clusters) in such alloys at RT, thereby drastically increasing the hardness and compressive stress of WQ + NA specimens.

4 Conclusions

  1. 1.

    Seven AMCs having nominal compositions Al80Mg5Si5Cu5Sn5, Al80Mg5Si5Cu5Zn5, Al80Li5Mg5Si5Cu5, Al80Li5Mg5Cu5Zn5, Al80Li5Mg5Zn5Sn5, Al68Mg10Si10Cu5Zn5Ni2, and Al50Mg30Si10Cu5Zn5 were produced by mold casting, and their microstructures and mechanical properties were investigated.

  2. 2.

    The area fractions of the secondary phases in the as-cast AMCs ranged from 26% to 58%. Various secondary phases such as Mg2Si, Si, θ-Al2Cu, Q-Al5Cu2Mg8Si6, η-Mg(Zn,Cu,Al)2, δ-Al3CuNi, Li-containing phases (i.e., AlLiSi, Al2CuLi, and AlLi), and Sn-containing phases (i.e., Sn and Mg2Sn) were present in the as-cast AMCs, depending on their chemical compositions. The Al matrices of the as-cast AMCs contained some alloying elements.

  3. 3.

    The types and amounts of secondary phases were more important than matrix strengthening in determining the composite hardness and compressive properties. Composite hardness and the maximum compressive stress increased while fracture strain decreased with increasing total area fraction of the secondary phases up to 40%. Compressive properties degraded in Al50Mg30Si10Cu5Zn5 showing numerous coarse primary Mg2Si phases. Al80Mg5Si5Cu5Sn5 and Al80Li5Mg5Zn5Sn5 also exhibited degraded mechanical properties owing to soft/heavy Sn-containing phases. Al68Mg10Si10Cu5Zn5Ni2 exhibited the highest specific strength owing to pronounced strengthening due to secondary phases and the Al matrix.

  4. 4.

    Although the composite hardness and compressive strength of Al68Mg10Si10Cu5Zn5Ni2 were deteriorated by annealing and AC, they were improved by continuous high-temperature annealing, WQ, and NA.