Microstructure and mechanical properties of Al-Mg-Si alloy fabricated by a short process based on sub-rapid solidification

1. Introduction

Aluminum is one of the most important lightweight metals, which has become inevitable and been widely used in automotive structures to reduce vehicle weight [1,2]. Specially, AA6xxx series alloys, containing Mg and Si as major elemental additions, are now being used for autobody applications. These alloys possess relatively good formability at solution state (T4), and have increased service strength at age-hardening state (T6 or T8). Traditionally, the fabrication process of Al-Mg-Si sheets is pretty long, which includes at least the following steps: slab casting under conventional solidification (CS), homogenization, hot rolling, cold rolling and solution heat treatment et al. This manufacture route is not economical, which needs long production time and a great deal of energy consumption. Moreover, severe chemical segregation and coarse eutectic microstructures will arise in slabs because of the low cooling rate during CS, which may bring negative effects on the ultimate properties. Homogenization, as a significant process to reduce or eliminate the micro-segregation and coarse eutectics, usually lasts from 6 to 24 h, with the temperature commonly no less than 550 °C [3]. Sjölander and Seifeddine [4] also proposed that homogenization resolves coarse particles formed during solidification, promoting solute diffusion to decrease the concentration heterogeneity. Accordingly, homogenization is necessary but treated as a high cost factor. A high strength EN AW 6082 alloy was researched without solution heat treatment, through adopting an appropriate forging operation that sufficient Mg and Si were got in Al matrix to offer adequate age-hardening subsequently [5]. It is in great urgent to design a new compact roadmap, which can shorten or even omit the homogenization for preparation of Al alloy sheets.

In general, solidification microstructure of alloys is directly influenced by the cooling rate. Note that the sub-rapid solidification (SRS) with a high cooling rate, ranging from 10² to 10³ °C/s, can significantly promote the grain refinement and microstructure homogeneity [6], which has a great potential to be used in industrial production. Some researchers have realized the SRS at the laboratory scale to produce Al, Mg alloys and stainless steels by using particular copper chilling moulds. For example, Zhang et al. [7] utilized the water-cooling copper mould to fabricate commercial casting Al-0.4Mg-7.0Si alloy, and the α-Al dendrites, volume fraction of eutectic phases (∼ 2.4%) and grain size (∼ 8 μm) all decrease in size scale. An Mg-6Al-1Zn-0.2 Mn alloy was prepared using copper mould based on SRS. In this work, the grain size obtained by SRS is about 1.8-13.5 μm, which is much smaller compared to that through CS (200-300 μm). Meanwhile, eutectic transformation is suppressed to a great extent in SRS, as evidenced by the decreased volume fraction of Mg₁₇Al₁₂ in this case compared to CS [8]. It was found that the quantity of austenite in 17Cr-9Ni-3Mo alloy increases from 60% to 82% as cooling rate climbing (from about 4 × 10⁴ to 2 × 10⁶ °C/s), by adopting copper mould and melt spinning, which means the constitution of microstructure changes under SRS process [9]. However, there is still much to be studied on the evolution of microstructure under SRS process.

In addition, higher mechanical properties can be obtained by SRS technique [10]. Zhang et al. [11] manufactured Mg-6Zn-1 Mn alloy by SRS (250 °C/s) and CS (10 °C/s), respectively. They found that samples produced under SRS possess has higher compressive strength of 460 MPa than CS sample (335 MPa), due to refined dendrite cell size (25 μm) in the former case. Specifically, the high cooling rate of molten metal in SRS relative to CS permits large deviations from equilibrium phase transformation, which can lead to fine metastable phases, weaker microsegregation, and promote the formation of supersaturated microstructure. Theoretically, the samples produced by SRS may have probability to be rolled directly without homogenization heat treatment.

In the present work, a novel and short route for the preparation of AA6022 alloy sheets has been proposed on basis of SRS, which omits the homogenization step. For comparison, the ordinary route for Al-Mg-Si rolling sheets involving CS, followed by homogenization was also taken into account. The mechanisms for microstructure evolution and mechanical properties based on the two routes were investigated in great detail. It is expected that the results will provide references for the further development of a short and low-cost fabricating process in aluminum autobody industry, such as twin-roll casting based on SRS technique realized by water-cooled copper rolls.

2. Experimental procedure

Fig. 1 is a schematic diagram to describe the three different manufacturing procedures in the present work. The AA6022 alloy used in this work was prepared from commercial pure Al (99.90 wt.%), Al-24.4Si master alloy, pure Mg (99.85 wt.%), Al-10 Mn master alloy, pure Fe (99.90 wt.%) and pure Cu (99.90 wt.%). The pure Al was completely melted in an electric resistance furnace, and then all the other raw materials were added to liquid Al sequentially. After being homogenized, the molten alloy was poured into two different moulds. A copper mould was used to realize the SRS cooling rate, with a cavity of 6 × 50 × 80 mm (thickness × width × length). Moreover, a steel mould was used for comparison to simulate the CS process, which was designed as cylinder shape with a cavity of Φ50 × 100 mm (diameter × length). After being measured by an optical spectrum analyzer (ARL 4460, Switzerland), the actual composition of alloy was Al-0.5Mg-1.3Si-0.07Mn-0.05Cu. Hereinto, the CS samples were cut from the center region of as-cast cylinder. Moreover, the SRS and CS samples have the same dimensions of 6 mm × 25 mm × 40 mm (thickness × width × length). All the samples were polished with a 2000 grit SiC sand paper to eliminate surface defects.

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Fig. 1.   Schematic diagram for the three different fabrication processes.

In the fabricating roadmap involving SRS, casting slabs were cold rolled to 1 mm directly without homogenization treatment, as shown in Fig. 1. While for the route based on CS, casting samples were firstly homogenized at 550 °C for 2 h in an electric resistance furnace, and then water quenched. Note that, 2 h holding time is chosen due to the small size of CS slab we utilized in this work, in reference to many literatures. It will take much longer time (6-24 h) for homogenization of industrial slabs with its increasing thickness. For comparison, samples based on CS without homogenization treatment were also prepared (CSN). The samples were cold rolled to 1 mm, and then all the rolling plates were solution heat treated at 550 °C for 30 min, followed by quenching in water. Here the three samples were denoted as SRS-T4, CS-T4 and CSN-T4, respectively. Then the rolling sheets at T4 state were subjected to a pre-strain of 10% by cold rolling to simulate practical forming process for autobody. Finally, artificial aging was carried out at 170 °C for 30 min (T6 state) to simulate the paint baking process used in automotive applications. At this step, the samples following three fabricating routes were labeled as SRS-T6, CS-T6 and CSN-T6, respectively.

K-type thermocouples and a data acquisition hardware (EM9104C, ZTIC, China) were utilized to collect the real-time data of solidification temperature at a scanning rate of 200 Hz. Microstructures of the samples were characterized by an optical microscope (OM, Carl ZEISS-Axio Imager.A2m, Germany), a scanning electron microscope (SEM, ZEISS EVO18, Germany) equipped with an energy dispersive spectrometer (EDS) analyzer (INCA-X-Max, England) and a transmission electron microscope (TEM, JEM-2100 F, Japan).Samples for OM observation were prepared by traditional grinding and polishing, followed by electro-polishing in perchloric acid solution (10 ml HClO₄, 90 ml alcohol). Then the as-polished samples were anodized in fluoroboric acid solution (5 ml HBF₄, 95 ml deionized water). Samples used for SEM observation were manually grounded and polished, and then chemically etched in HF solution (5 ml HF, 100 ml deionized water). The TEM thin foils of 3 mm diameter were prepared by twin-jet electron polishing using a mixture of HClO₄ and alcohol. Moreover, the electron backscattered diffraction (EBSD) analysis was executed on a scanning electron microscope (VEGA 3 XMU, TESCAN, Czech) equipped with an Oxford Instruments NordlysNano EBSD detector, which works with an accelerating voltage of 20 kV, a working distance of 20 mm and a scanning step size of 1.0 μm. The AZtec and Channel 5.0 software was used to collect and analyze data. The Nano Measurer 1.2 software was utilized to estimate grain size and its distribution. Compositional analysis of solute distribution was performed on as-polished samples using an electron beam microprobe analyzer (EMPA, JEOL JXA-8230, Japan) equipped with four wavelength dispersive X-ray spectrometers (WDS), which works with an working current of 1 × 10^-8 A, and a working distance of 11 mm. 144 points were collected and analyzed in a square of 110 μm × 110 μm. Furthermore, tensile tests were carried out along the rolling direction using the INSTRON 5869 testing machine (INSTRON, Boston, USA) under a strain rate of 1 × 10^-3 s^-1. At least three tensile specimens, with gauge length of 10 mm, for each condition were tested to ensure a good reproducibility.

3. Results

3.1. Solidification process

The cooling curves of molten alloys solidified during SRS and CS were measured, respectively. K-type thermocouples were placed at the center of cavity and about 20 mm high from the bottom, as shown in Fig. 2 (a). The entire cooling curves of SRS and CS samples are shown in Fig. 2 (b). Furthermore, Fig. 2 (c) and (d) illustrate the details of initial stage on the two cooling curves. The cooling rate of Al melts is evaluated based on the interval between pouring temperature and solidification onset temperature (liquidus point), which is the slope of straight-line portion of cooling curves [12]. It can reflect the real heat extraction by moulds. The cooling rate of solidification in copper mould reaches as high as about 160 °C/s, which is treated as SRS (10² -10³ °C/s). However, in terms of CS performed in steel mould, the cooling rate is just about 30 °C/s. Note that the onset and ending solidification temperature of SRS sample, which corresponds to the liquidus point and solidus point, decreases synchronously compared with the conventional case. Hereinto, the decrease of liquidus point is not significant, valuing 7 °C with increase of cooling rate, indicating that a little higher undercooling is in need for SRS. However, the variation of solidus point is substantial (decrease by 20 °C). Accordingly, the solidification interval becomes larger for SRS sample (59 °C) than CS sample (46 °C). Birol [13] has also drawn a similar result, by employing Differential Scanning Calorimetry (DSC) method. He found that as cooling rate increases from 0.04 to 0.67 °C/s, the solidus point of AA6016 alloy decreases merely 3.3 °C, while liquidus point declines by 29.8 °C, leading to the widen of solidification interval of 43.2-69.7 °C.

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Fig. 2.   (a) Schematic diagram of temperature measuring system; (b) entire cooling curves of SRS and CS samples; (c) detail of initial stage on SRS cooling curve; (d) detail of initial stage on CS cooling curve.

3.2. Microstructure characteristics

Fig. 3 (a) and (b) indicate the typical as-cast microstructure of samples following the three different solidification processes. Distinctly, the sample subjected to SRS has finer primary grains with a secondary dendrite arm spacing (SDAS) of about 12 μm, while that for CS sample is about 36 μm. The refined grain size can be attributed to the increased nucleation rates induced by larger undercooling and inhibited grain growth by high cooling rate during SRS. Fig. 3 (c) illustrates the homogenized microstructure of CS sample after being heat treated at 550 °C for 2 h. As can be seen, the homogenized sample possesses cellular grains with an average size of about 68 μm.

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Fig. 3.   Optical micrograph (OM) of (a) as-cast SRS sample, (b) as-cast CS sample and (c) as-homogenized CS sample.

Fig. 4 (a) and (b) exhibits the SEM micrographs of SRS and CS samples, and the corresponding composition distribution profiles of Fe, Mg and Si along line 1 and line 2 are shown on right. Clearly, there are two main intermetallic phases at grain boundaries after casting. The black phases at grain boundaries are composed of Mg and Si, existing in the form of eutectic Mg₂Si [4,5,14,15]. On the other hand, Al, Fe and Si are the main elements for white phase. Previous works point out that β-AlFeSi exhibits needle-like morphology in two-dimension direction [[15], [16], [17], [18]], which agrees well with our present results. Specifically, there are more white AlFeSi phase and less black Mg₂Si phase under SRS compared to CS condition. Wang et al. [14] researched two Al-0.8Mg-0.9Si-0.5Cu-0.1 Mn-0.2Fe alloys (with and without Zn content) under CS condition, and they also found large secondary particles of white AlFeSi and black Mg₂Si phases in their solidification microstructure. In the present work, it is found that the constituents of secondary phases along grain boundary change with cooling rate, and the formation of Mg₂Si phase is obviously restrained in SRS process. The mechanism for microstructure evolution with cooling rate increasing will be discussed in detail in sections below.

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Fig. 4.   Backscattered electron imaging (BSE) micrographs observed in as-cast samples of (a) SRS and (b) CS; (line 1), (line 2) corresponds to the composition distribution profiles of Fe, Mg, Si in (a) and (b), respectively.

The average partition coefficient (k) of solute elements, being acquired on basis of EPMA chemical data, can be used to estimate the extent of microsegregation. The partition coefficient of Mg (k_Mg), Si (k_Si) and Fe (k_Fe) under SRS and CS processes are calculated according to the method of GANESAN et al. [19]. The EPMA data for Mg, Si and Fe elements are sorted in ascending order. Then the solid fraction (f_s) is calculated according to rank number (R_i) and total number of points (N) by f_s = (R_i - 0.5) / N. The average partition coefficient is determined through the conventional Schiel equation:

C_s = kC₀(1 - f_s)^k-1 (1)

Afterwards, the value of k is obtained through linear regression, as follows:

ln (C_s / C₀) = (k - 1) ln (1 - f_s) + ln k (2)

where C_s is the solute concentration at each point for every solute element, and C₀ is the corresponding average content of ingot. The calculated k_Mg, k_Si and k_Fe are listed in Table 1, as a result, average partition coefficients of Mg, Si, and Fe elements under SRS condition are higher than those under CS process, particularly for Mg atoms. Consequently, SRS sample exhibits weaker microsegregation owing to higher cooling rate during solidification.

After being cold rolled with 83% reduction, the microstructure of samples following SRS and CS routes were descripted in Fig. 5 (a) and (b), respectively. Grains of the two samples are both elongated significantly by cold rolling, and the SRS sample has much smaller size contributed by the inheritance of refined solidification microstructure.

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Fig. 5.   OM micrographs of cold-rolled (a) SRS and (b) as-homogenized CS samples.

Fig. 6(a)-(d) presents the representative grain characteristics of SRS-T4 and CS-T4, as well as the corresponding grain size distribution. After being solution heat treated, the two samples have undergone recrystallization thoroughly. The SRS sample exhibits an average grain size of about 19 μm, which is 42% smaller than that of the CS sample.

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Fig. 6.   (a) and (c) EBSD inverse pole figure (IPF) maps of SRS-T4 and CS-T4; (b) and (d) correspond to grain size histogram of (a) and (c).

The TEM micrographs of SRS-T4 and CS-T4 are given in Fig. 7 (a) and (b), respectively. The precipitates are mainly AlFeSi phase based on the selected area diffraction (SAED), with mostly spherical morphology and also rod-like shape. Obviously, CS-T4 has higher density of precipitates, with area fraction of about 1.2 × 10¹² m^-2, in comparison with SRS-T4 (∼ 0.6 × 10¹² m^-2). Moreover, the precipitates of CS-T4 exhibit a narrower edge-to-edge spacing (∼ 400 nm) and smaller average radius (∼ 20 nm), compared to those of ∼ 1000 nm and ∼ 45 nm in SRS-T4.

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Fig. 7.   TEM micrographs and corresponding selected area electron diffraction (SAED) patterns obtained from the given particles: (a) SRS-T4, (b) CS-T4; TEM micrographs and high resolution TEM images (HRTEM) of the given particles: (c) SRS-T6 and (d) CS-T6.

However, the precipitates under T6 state are primarily Mg₂Si phase [20] based on the high resolution TEM images (HRTEM), with typical rod-like shape (Fig. 7 (c) and (d)). Unlike at T4 state, SRS-T6 displays a shorter edge-to-edge spacing (∼ 180 nm) and higher area fraction (∼ 3.7 × 10¹² m^-2) than those of CS-T6 (∼ 230 nm and ∼ 1.9 × 10¹² m^-2). In addition, precipitates of SRS-T6 has smaller average radius (∼ 17 nm), about 55% smaller than that of CS-T6 (∼ 38 nm).

After undergoing a 10% pre-strain and being artificial aging treated (T6 state), grains of the two samples have been elongated slightly along the rolling direction (Fig. 8). Dislocation density in grains can be estimated by ρ≈θ/bδ, as developed by Chen et al. [21] based on EBSD data. Here b is the Burger vector, equal to 0.286 nm for fcc Al, and θ is the accumulated misorientation angle (in radian) within a distance δ. The parameters of θ and δ can be extracted from the misorientation profiles in Fig. 8. The dislocation density in grains of SRS-T6 is about (7-25) × 10¹³ m^-2, while that is (1-4) × 10¹³ m^-2 for CS-T6, as measured in at least five typical grains in every sample. It can be seen that SRS-T6 has higher dislocation density.

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Fig. 8.   Typical EBSD inverse pole figure maps of (a) SRS-T6 and (c) CS-T6; (b) and (d) corresponding misorientation profiles measured along the lines in (a) and (c).

3.3. Mechanical properties

Typical tensile engineering stress-strain curves of the samples at T4 and T6 states are shown in Fig. 9 (a) and (b) respectively, with corresponding yield strength (YS), ultimate tensile strength (UTS) and elongation (El.) listed in Table 2. For comparison, relevant properties for commercial AA6022 alloys obtained from other work involving CS, are also showed there. As compared to the references, CS samples in this work have similar mechanical properties to those reported in other papers. Furthermore, it can be seen that SRS-T4 has lower yield strength (∼ 85 MPa) and higher strain-hardening exponent (∼ 0.351) than those (∼ 113 MPa, ∼ 0.340) in CS-T4, which indicates it has better formability than CS-T4. After being imposed 10% pre-strain and artificial aging (170 °C, 30 min) treatment, SRS-T6 reversely obtains slightly higher yield strength (∼ 248 MPa) and larger ultimate tensile strength (∼ 295 MPa) than those of ∼ 238 MPa and ∼ 288 MPa in CS-T6, signifying better ability of dent resistance in practical service. Moreover, the sample following the CS route without homogenization (CSN) has similar yield strength and ultimate tensile strength, but remarkable lower elongation in comparison with CS samples at T4 and T6 states, respectively. On the other hand, the yield strength increment of SRS sample from T4 to T6 state is about 163 MPa, while that for CS sample is about 125 MPa. For all samples, enhancement of yield strength from T4 to T6 states owes to both dislocation strengthening and precipitation strengthening.

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Fig. 9.   Typical tensile engineering stress-strain curves of SRS, CS and CSN samples at (a) T4 and (b) T6 states.

4. Discussion

4.1. Competitive relationship between Mg₂Si and AlFeSi phases in SRS and CS process

During a non-equilibrium solidification process, a relationship for velocity-dependent partition coefficient of solute atoms has been derived by Aziz [25]:

K_v=(k_e + δV/D)/(1+δV/D) (3)

where k_e is the equilibrium partition coefficient, V is the growth rate, δ is the diffusive-like jump of length, and D is the interface diffusion coefficient. Under SRS condition, samples have much higher growth rate because of the high cooling rate.

It can be concluded that the value of k_v is within the range between k_e and 1, mainly decided by the variable of V. When the V value is extremely small, it can be speculated that k_v ≈ k_e. On the contrary, k_v ≈ 1 if the grain growth rate increases to big enough [26]. Furthermore, the solidification is implemented in a lower temperature region for the SRS process as mentioned above, so solute atoms in liquid have smaller value of D, implying the interface diffusion is impeded to a certain extent. Taken together, the k_v value is larger under SRS condition on account of higher cooling rate compared with that of CS process, meaning that more solute atoms are partitioned into solid phase (α-Al matrix).

The evolution of Mg₂Si phase lies on the diffusion behavior of Mg atoms during solidification. According to the equilibrium binary phase diagrams of Al-Mg, Al-Si and Al-Fe systems, the maximum solid solubility of Mg (14.9%) in Al is much higher than that of Si (1.65%) and Fe (0.052%) [27]. On the other hand, according to the average partition coefficients of Mg, Si and Fe atoms listed in Table 1, k_Mg is higher than k_Si and k_Fe at both CS and SRS condition. However, there is a higher increment of k_Mg (0.15) than k_Si and k_Fe (0.04 and 0.03) for SRS sample, implying that more Mg atoms are inclined to be trapped in α-Al matrix during SRS [28], leading to less Mg₂Si phase formed at grain boundaries. Thus, more excess Si atoms can combine with Fe atoms to form higher proportion of AlFeSi phase. However, the situation in CS is not the same. Mg atoms can easily diffuse to liquid from solid/liquid (S/L) interface because of lower cooling rate, and therefore facilitate the formation of considerable Mg₂Si phase and inhibit AlFeSi phase.

4.2. Mechanism for lower YS of SRS-T4

Note that the yield strength of SRS-T4 is about 28 MPa lower than that of CS-T4. It is known that the dominant strengthening mechanisms include grain boundary strengthening (Hall-Petch, σ_GB), dislocation strengthening (σ_D), precipitation strengthening (σ_P), and solid-solution strengthening (σ_SS) [29,30]. So the yield strength σ can be written as:

σ = σ₀ +σ_GB + σ_D +σ_P + σ_SS (4)

where σ₀ is a constant decided by the intrinsic characteristics of metal matrix. In the present experiment, the two samples have the same σ₀, about 28 MPa, which is the yield strength for commercial pure Al at annealed condition [31].

The grain boundary strengthening can be calculated using the Hall-Petch equation [32,33]:

σ_GB = k_yd^-1/2 (5)

where d is the average grain size, and k_y is the Hall-Petch coefficient, approximately 0.068 MPa m^1/2 for this alloy [34]. Accordingly, the grain boundary strengthening contributes to approximate 16 MPa and 12 MPa for the SRS-T4 and CS-T4, respectively. The difference of yield strength between the two samples caused by grain boundary strengthening is pretty close, which can be neglected.

The dislocation strengthening is described by the Taylor equation [35]:

σ_D = MαGbρ^1/2 (6)

where G is shear modulus, about 27 GPa for Aluminum alloys, M is the Taylor factor, equal to 3.06 for fcc metals, α is a constant, equal to 0.2 for fcc metals, b is magnitude of the Burgers vector, equal to 0.286 nm for fcc Al, and ρ is dislocation density [30]. Since the two samples undergo the same flows of cold rolling and solution treatment (T4), they are inferred to both have low dislocation density with the same order of magnitude. This results in the similar effects of dislocation strengthening. Besides, the difference in solute strengthening effects between the SRS-T4 and CS-T4 can also be ignored, because of their little contribution to improved YS relative to other strengthening mechanisms.

Furthermore, the precipitation strengthening is estimated based on Orowan looping mechanism [36]:

σ_p=(MGb/2πλ$\sqrt{1-v}$)⋅ln(πd/4b) (7)

where υ is the Poisson’s ratio, about 0.33 for Al, λ is the average spacing between particles, and d is the average diameter of precipitates. It can be predicted that the contribution of precipitate strengthening to YS is about 25 MPa and 54 MPa for the SRS-T4 and CS-T4, respectively.

Therefore, the weaker precipitation strengthening is the major factor that causes lower YS of SRS-T4. The calculated difference of precipitation strengthening is 29 MPa, which is in well consistent with the actual deviation of YS value, about 28 MPa in this experiment.

4.3. Mechanism for higher YS increment of SRS sample from T4 to T6 states

It is worth noting that there is a higher yield strength increment of SRS sample from T4 to T6 states than that of CS sample. According to the analysis in Section 4.2, difference of grain boundary strengthening can be ignored because of the low Hall-Petch coefficient (0.068 MPa m^1/2) of Al alloys. Owing to pre-strain and heat treatment were carried out to achieve T6 state, precipitation strengthening and dislocation strengthening are introduced and treated as the major strengthening mechanisms after T6 treatment.

According to Eq. (7), precipitation strengthening of SRS-T6 and CS-T6 samples is estimated to be 130 MPa and 107 MPa, respectively. Consequently, the increment of precipitation strengthening from T4 to T6 for SRS is 105 MPa, being larger than that of CS (53 MPa). The larger rise in precipitation strengthening is induced by higher density of Mg₂Si precipitates in SRS-T6. On the other hand, the dislocation strengthening is assessed by Eq. (6) and average dislocation density, valuing 60 MPa and 24 MPa for SRS-T6 and CS-T6, respectively.

As a whole, precipitation strengthening and dislocation strengthening both contribute to the YS increment of samples from T4 to T6 state. Nevertheless, the former contribution is more significant. Owing to an extending solid solubility is obtained under sub-rapid solidification, more precipitates of tiny Mg₂Si are generated during T6 treatment, inducing an enhanced precipitation strengthening of SRS-T6.

4.4. Advancement of SRS process

All the results obtained in this work prove that the preparation process based on SRS is not only short, but also beneficial to mechanical properties of products. Under CS conditions, the aluminum billets have to be homogenized at high temperature for a long time. Since supersaturated aluminum solid solution and weaker microsegregation can be acquired by SRS process, these slabs can be rolled directly without homogenization heat treatment. Moreover, aluminum sheets at T4 sate have lower yield strength, which is conducive to subsequent drawing process, through the adoption of SRS. In addition, these sheets at T6 sate have similar or even higher yield strength compared to that of products fabricated by CS technological path, means they can completely satisfy the service requirements. The peculiar features of this SRS route, in a word, are timesaving, low energy-consumption, and can sustain high performance in service.

5. Conclusion

In present work, the microstructure evolution and mechanical properties of Al-0.5Mg-1.3Si alloy, fabricated by the routes based on SRS, are studied in detail. It was found that melts solidifies with a larger undercooling (∼ 7 °C larger than CS) and wider solidification interval (∼ 59 °C) under SRS. The SRS sample has a much smaller SDAS (∼ 12 μm) than CS sample (∼ 36 μm). The sample under SRS has higher average partition coefficients of solute elements, especially for Mg atoms, leading to weaker microsegregation owing to higher cooling rate (160 °C/s). Besides, Mg atoms tend to be trapped in Al matrix with increasing cooling rates, preventing the generation of Mg₂Si at grain boundaries, and promoting the formation of AlFeSi under SRS. Moreover, samples fabricated by SRS route (∼ 85 MPa) has lower YS than that of CS sample (∼ 113 MPa) at T4 state, indicating better formability. The lower YS in SRS-T4 sample is mainly attributed to the weaker precipitation strengthening effect. Furthermore, after being deformed and artificial aging heat treated, samples fabricated through SRS have higher enhancement of YS (163 MPa), inducing higher YS (∼ 248 MPa) at T6 state. This work certifies that SRS technique has great potential to produce high performance Al-Mg-Si series alloys in a short and low-cost process. Therefore, our present work inspires the further development of twin-roll casting based on SRS technique in industries, which can be realized by the water-cooled copper rolls.

Acknowledgements

Financial supports from The National key research and development program (No. 2016YFE0115300) and The Natural Science Foundation of China (Nos. 51790483, 51625402, 51790485 and 51801069) are greatly acknowledged. Partial financial support came from The science and technology development program of Jilin Province (No. 20190901010JC), The Changjiang Scholars Program (T2017035) and the Program for JLU Science and Technology Innovative Research Team (JLUSTIRT, 2017TD-09).