Precipitation behavior and tensile properties of A356.2 alloy with different high temperature pre-precipitation temperatures

In order to understand the effect of high temperature pre-precipitation (HTPP) temperature on the precipitation behavior and tensile properties of A356.2 alloy treated by Al-6Sr-7La master alloy, SEM, TEM and tensile tests were applied to investigate the evolution of fracture morphology and precipitates of the alloys under different HTPP temperatures. The results showed that ultimate tensile strength (UTS) and yield strength (YS) of the alloys decrease and elongation (El) increases with HTPP temperature decreasing. When HTPP temperature decreased from 510 °C to 470 °C coarsen coherent β″ phase appear in α-Al matrix, continue to decrease HTPP temperature to 450 °C the main precipitate transformed into semi-coherent β′ phase, leading to the change in mechanical properties. In addition, coarsening and transformation of the precipitate were attributed to the reduction of Si concentration which decreases with HTPP temperature decreasing. Moreover, Si nanoparticles precipitated in α-Al matrix, leading to the decrease of UTS and YS to certain extent due to reducing Si concentration during aging process.


Introduction
Due to its light weight, high specific strength, high thermal conductivity, high electrical conductivity, corrosion resistance, low thermal expansion coefficient and good fluidity, Al-Si-Mg alloy have been developed in communication field [1][2][3]. With the development of communication technology, the wall thickness of radiators for communication base station is getting thinner and thinner, which increases the possibility of fracture during transportation and installation [4]. Therefore, higher requirements are needed for Al-Si-Mg alloy mechanical performances. Coarse lamellar or needle-like eutectic silicon and large grain size in unrefined and unmodified Al-Si-Mg alloy have negative effect on its mechanical properties [1,[5][6][7]. Therefore, in order to improve its comprehensive mechanical properties, refiners and modifiers are usually applied to reduce the grain size of α-Al and improve the morphology of eutectic silicon [8][9][10][11]. Even so, solute segregation still existed in ascast alloy which reduces its mechanical properties [12].
Generally, solution treatment and aging treatment are two main methods to effectively optimize the microstructure and mechanical performances of Al-Si-Mg alloys [13][14][15][16]. Recently, researchers have focused on the adjustment of precipitation phase in Al alloys by multi-stage aging, including retrogression and re-aging treatment (RRA), interrupted artificial aging treatment and high temperature pre-precipitation treatment (HTPP). HTPP has a shorter process compared to other multi-stage aging and can be applied to large crosssectional parts, which is suitable for industrial production of Al-Si-Mg radiators [17][18][19]. It has been proved that the ultimate tensile strength (UTS) of Al-Zn-Mg-Cu alloy is significantly improved after HTPP treatment [20] and with the increase of HTPP temperature UTS gradually increased. Meanwhile, the UTS reached the maximum when HTPP temperature was close to the solid solution temperature. In addition, there is a larger range of HTPP temperature in Al-5.1Mg-3.0Zn-0.15Cu alloy to reach the maximum of UTS which is 410°C-450°C [21]. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI.
At present, the alloys using HTPP process are mainly concentrated in 7××× and 5××× aluminum alloys, few researchers focused on Al-Si-Mg alloys. Previously, Al-6Sr-7La composite refinement-modification agent was developed and its synergistic effect on microstructure and tensile properties of A356.2 alloy were investigated [22,23]. The results showed that Al-6Sr-7La possessed excellent refinement and modification effect which significantly increased the UTS, yield strength (YS) and elongation (El) of A356.2 alloy. Based on this, the effect of HTPP treatment on the precipitation and tensile properties of A356.2 alloy treated by Al-6Sr-7La refinement-modification agent was studied in this work and the precipitation mechanism of the precipitates was analyzed.
2. Materials and experimental procedures 2.1. Material and sample preparation Commercial A356.2 alloy and Al-6Sr-7La master alloy were used as the matrix and composite refinementmodifier respectively in this work. Detailed preparation processes were as follows. Firstly, about 1kg of A356.2 alloy was melted at 750°C in graphite crucible in a resistance furnace. Secondly, after the alloy was melted, slags were removed and high-purity argon gas with the flow rate of 1.8 l min −1 was used to degas the melt for 2 min. Thirdly, removed the slag and added 0.5 wt% Al-6Sr-7La master alloy into the melt, then lowered the melt temperature to 730°C and held for 3 min, after that high-purity argon gas with the flow rate of 1.8 l min −1 was used to degas the alloy for 3 min. At last, holding the melt at 730°C for 3 min, removed the slag and poured the melt into a cast iron mold (with its inner diameter of 45 mm and height of 160 mm) which was preheated to 200°C to obtain the as-cast samples. The chemical compositions of A356.2 alloy treated by 0.5 wt% Al-6Sr-7La composite refinement-modifier is shown in table 1. Figure 1 shows the heat treatment process of the alloy. A356.2 alloy treated by 0.5 wt% Al-6Sr-7La composite refinement-modifier was firstly treated at 540°C for 3h (solid solution treatment) in box type resistance furnace with temperature accuracy of ±1°C. After that, specimens were quenched into 60°C water with transfer time from box type resistance furnace to the water less than 3 s. For HTPP treatment, five factors were taken into consideration, that is, HTPP temperature (T1), HTPP time (t2), quenching temperature (T2), aging

Mechanical properties
WDW-200 material testing machine with displacement control and stretching speed of 1.5 mm min −1 was used for room temperature tensile testing. Specimens were cut in accordance with GB/T228-2010. At least 3 samples were tested for each parameter to ensure reproducibility and the average data were used in this work.

Microstructure observation
JEOL JSM-6510A with accelerating voltage of 15 kV and working current of 60 μA was used for tensile fractures observation. FEI Tecnai G2 F20 with accelerating voltage of 200 kV was used for TEM observation and specimens for TEM observation were firstly cut into 300 μm slices, then the slices were thinned to 40-50 μm using sandpapers from coarse to fine, followed by punching into Φ3 mm discs. After that the discs were electrolytically double-jet thinned using a double-jet thinner (Struers Tenupol-5, dry ice and ethanol were applied to control the temperature at −30±1°C, double-jet electrolyte was 30% nitric acid and 70% methanol solution). Then ion thinning (Gatan 691) with voltage, current and incidence angle of ion beam of 5V, 0.5mA and ±3°respectively was used to increase thin zone area. Liquid nitrogen was used for cooling throughout the process.

HTPP uniform experiment
Tensile stress-strain curves for each parameter of the HTPP uniform experiment is shown in figure 2(a). Taking sample 4 as an example, the curves consist of three stages, that is, elastic deformation stage (I), plastic deformation stage (II) and fracture stage (III). During the elastic deformation stage, stress increases rapidly with the increase of strain until the strain exceeds a certain value which reaches to the plastic deformation stage. Continue to increase the strain, necking fracture occurs. Figure 2

Tensile properties
UTS is an important index to express the tensile properties of alloys and as shown in table 3, HTPP temperature has the greatest influence on UTS. Based on the optimum process parameters about UTS (table 4), four HTPP temperatures were applied to further optimize the tensile properties of the alloy and the tensile properties are shown in figure 3. It can be seen UTS and YS gradually decrease with HTPP temperature decreasing while the change in EI is opposite. UTS for the alloys with HTPP temperatures of 510°C and 490°C are almost unanimous  ( figure 3(a)). The alloy with HTPP temperature of 510°C possesses the highest UTS (300 MPa) and YS (271 MPa) values. While the alloy with HTPP temperature of 450°C has the lowest UTS (175 MPa) and YS (116 MPa) values, which reduce by 41.67% and 57.20% respectively compared with the alloy with HTPP temperature of 510°C. While, the alloy with HTPP temperature of 510°C possesses the lowest El of 9.3%. All of these mean that HTPP temperature has an important effect on the mechanical properties of A356.2 alloy treated by Al-Sr-La composite refinement-modification agent ( figure 3(b)). Fracture morphologies of the alloys under different HTPP temperatures are shown in figure 4. There are obvious differences between them, e.g., quasi-cleavage planes and tear ridges exist in the alloys with HTPP temperatures of 510°C and 490°C, meaning mixed fracture mode of cleavage fracture and ductile fracture. Moreover, there are cracked Si particles in the dimples, indicating the occurrence of transgranular fracture. With HTPP temperature decrease to 470°C, fracture mode changes to ductile fracture with fewer cleavage planes. In addition, dimples gradually become deeper with the decrease of HTPP temperature, indicating the increase of plasticity.

Strengthening phase evolution
Generally, mechanical properties of aluminum alloys are greatly affected by precipitates in Al-Si-Mg alloys [15]. The morphology evolution of strengthening phase in alloy under different HTPP temperatures is shown in figure 5 and it is clear HTPP temperature has great influence on the precipitates, e.g., as the temperature decrease, the size of precipitates increases, and the number of precipitates decreases (figures 5(a-1)-(d-1)). According to the microscopic morphologies (figures 5(a-2)-(d-2)) and Fast Fourier Transform (FFT) results (figures 5(a-3)-(d-3)), nanoscale phase is needle-like β″ phase as the HTPP temperature decreases from 510°C to 470°C. When HTPP temperature decreases to 450°C, the precipitates change to rod-like β′ phase and a small amount of needle-like β″ phase. Generally, strengthening effect of β″ phase is better than β′ phase in Al-Si-Mg alloys [15,24,25]. Therefore, as HTPP temperature decreasing, the number of β″ phases decrease with β′ phases increasing, resulting in YS decrease.
β″ phase and β′ phase are all composed of Mg element and Si element, and their precipitation is influenced by solid solubility of Si and Mg. Mg content in the alloy is much lower than its upper limit of solid solubility, while Si content is much higher than its solid solubility limit in the range of 450°C-510°C. Therefore, as the HTPP temperature decreasing from 510°C to 450°C, supersaturation of Mg element remains basically unchanged, while supersaturation of Si element decreases, resulting in reduce the content of Si element participating aging process. Figure 6 shows that some spherical nanoparticles appear in the α-Al matrix with the HTPP temperature of 450°C. According to the HRTEM ( figure 6(b)) and FFT results (figure 6(c)), the spherical nanoparticle is determined to be Si nanoparticle. In addition, the Si nanoparticle is not found in the α-Al matrix as the HTPP temperature decreases from 510°C to 470°C. It shows that Si nanoparticle is easier to precipitate at lower HTPP temperature as the HTPP temperature decreasing from 510°C to 450°C. The Si nanoparticle can grow up during the aging process [26], which can reduce the content of Si element participating precipitation of β″ phase and β′ phase during aging process. Therefore, due to the decreasing of HTPP temperature and the precipitation of Si nanoparticle, the amount of Si element participating in the precipitation of β″ phase and β′ phase during subsequent aging process decreases, leading to β″ phase and β′ phase decrease. Generally, in the aging process of Al-Si-Mg alloy, the shorter the peak aging of the same alloy, the larger precipitate under the same aging time and aging temperature. Some studies have shown that the reduction of Si concentration can shorten peak aging time [27,28]. Therefore, lower concentration of Si atoms as the HTPP temperature decreasing from 510°C to 470°C, shorter peak aging time, resulting in β″ phase coarsening.
In Al-Si-Mg alloy, the accepted precipitation sequence is: super-saturated solid solution (SSSS) → clusters/ GP zones → β″ → β′ → β + Si [29][30][31]. In this study, the precipitation phase is mainly β′ phase with a small amount of β″ phases when HTPP temperature is 450°C. Studies have shown that there is no β′ phase in Al alloys at temperatures above 400°C [29,32,33]. Therefore, β′ phase is not precipitated during HTPP process. Meanwhile, it has been shown that β″ phase will transform into β′ phase at the temperature of above 200°C [29,32,34,35]. However, in this work, aging temperature is lower than 200°C, indicating that precipitation of β′ phase is not due to high temperature overaging. When supersaturation of Si atom is lower in alloy matrix, β′ phase is easier to nucleate and grow than β″ phase during aging stage [36]. Therefore, it is inferred that precipitation of β′ phase is caused by the change of Si concentration. This can be explained by considering the nucleation probability or rate J of β″ phase and β′ phase, expressed by equation (2) [37,38].
where ΔG is nucleation energy barrier, k B is Boltzmann constant, T is temperature. In addition, ΔG can be expressed by equation (3) [37,38].
where C Si is supersaturation of Si atom in matrix before aging, C eq Si is characteristic concentration value of Si atom, below which precipitate phase cannot precipitate. Moreover, Ω is related with the interfacial energy between nuclei and matrix, which can be assumed as constant for a certain phase to nucleate during aging [39].
Ω value of β′ phase is much higher than that of β″ phase, C eq Si of β′ phase requires smaller Si concentration than that of β″ phase and C eq Si is related to temperature [39]. During aging process after different HTPP, the aging temperature and time all are same, soT and C eq Si is constant value. Therefore, combined with equation (3), the relationship between C Si and ΔG can be expressed qualitatively, as shown in figure 7. In figure 7, ΔG of β″ phase and β′ phase decrease with increasing C , Si and there must be a focal point (C 0 Si ) between the two curves, which is due to that higher Ω value of β′ phase resulting in faster descent speed of corresponding ΔG. When C Si <C , 0 Si ΔG of β′ phase is lower, otherwise that of β″ phase is higher. It can be seen from equation (2) that the smaller the G value, the larger the J value, and corresponding precipitate is firstly nucleate and grow in supersaturated matrix. Thence, when C Si is lower than C , 0 Si β′ phase firstly precipitate, rather than transitioning from β″ phase. Therefore, precipitation of many β′ phases during aging stage after HTPP process of 450°C is attributed to very low C Si value before aging stage.

Corresponding relationship between precipitation phase and YS
In this work, the change of alloy YS is mainly attributed to precipitation hardening of the strengthening phase [32,40,41]. According to the Orowan model, YS of the alloy can be expressed by equation (4)    where YS s is YS of alloy, ppt s D is precipitation hardening, G m is shear modulus of the matrix, b is Burgers vector of the matrix, d p is average particle size and l is the average particle spacing. Also, l is can be calculated by (5).
where V p is the average volume fraction of particles. So YS value of the alloy can be expressed by V p and d p , as shown in (6). ( ) s = -According to equation (6), with V p value increasing and d p value decreasing, YS becomes higher (d p >e). In this work, with HTPP temperature decreasing from 510°C to 470°C, V p value decreases with d p value increasing, resulting in YS decreasing. Although V p and d p are similar at the HTPP temperature of 450°C and 470°C, the strengthening phase contains not only a small amount of β″ phases but also many β′ phases at the HTPP temperature of 450°C. The semi-coherent β′ phase reduce YS of the alloy, because low degree of lattice distortion between the β′ phase and the matrix weakens the pinning effect on dislocations compared with coherent β″ phase [15,24,25].

Conclusions
The effect of HTPP temperature on the precipitation and tensile properties of A356.2 treated by Al-Sr-La composite refinement-modification agent was investigated, following conclusions can be obtained: (1) HTPP treatment has a significant effect on the tensile properties of A356.2 alloy, where HTPP temperature has the most significant effect according to the regression analysis results of the uniform test.
(2) UTS and YS decrease with HTPP temperature decreasing from 510°C to 450°C while the change in EI is opposite. When HTPP temperature is 510°C, UTS, YS and El of alloy can reach to 300 MPa, 271 MPa and 9.3% respectively. During the heat treatment process, the precipitates are mainly β″ phase and β′ phase which affect the tensile properties of alloy based on their morphology and quantity.
(3) The solid solubility of Si element decreases during HTPP process with HTPP temperature decreasing. Moreover, Si nanoparticles precipitated in α-Al matrix reduce the content of Si element participating precipitation of β″ phase and β′ phase during aging process. Therefore, the β″ phase gradually coarsens and the number decreases as the temperature of HTPP decreases, while the number of β′ phase increases, leading to the decrease of UTS and YS.