The size of the dendritic cells (dendritic parameter, d), as well as of the intermetallics, was experimentally determined using metallography of high-contrast microstructural images processed with the appropriate software, ImageJ (National Institutes of Health, Bethesda, MD, USA). Figure 3b shows the volume fraction of the Al3Ni phase in Al-8%Zn-3%Mg-xNi(x) alloy depending on temperature under equilibrium conditions. The microstructure of the FC sample (Figure 4a) can be qualified as closest to equilibrium.
A further increase in the cooling rate resulted in the formation of intermetallic bands indicating a proximity of the solidification path to the eutectic. It is striking that the majority of Zn and Mg are dissolved in the (Al) matrix even during cooling at 133 K/s. In the microstructure of the 1 mm cast sample (Figure 6), the hard casting condition resulted in the formation of shrinkage cavities along the grains.
Microstructural characteristics and properties of sputter formed Zn-rich Al-Zn-Mg-Cu alloy under different aging conditions. Mater. Fabrication of as-cast Al matrix composite reinforced by Al2O3/Al3Ni hybrid particles via in-situ reaction and evaluation of its mechanical properties. Mater. Effect of nickel on the structure, mechanical and casting properties of aluminum alloy type 7075. Mater.
The influence of the Al3Ni and Mg2Si eutectic phases on the casting properties and hardening of the Al-7%Zn-3%Mg alloy. Mater. Investigation of aging and tensile properties of Al-Zn-Mg/Al3Ni eutectic composite. Met. Effect of Zn/Mg ratio on microstructure and mechanical properties in Al-Zn-Mg alloys. Mater.
Inﬂuence of Continuous Casting Speeds on Cast Microstructure and Mechanical Properties of
- Experimental Procedure 1. Materials
The tensile strength and strain of the samples were obtained from the tensile stress-strain curve. It can also be clearly observed that the average sizes of primary Si are the largest at a speed of 1.0 mm/s, regardless of the observed location. Furthermore, the distributions of primary Si phases are non-homogeneous regardless of the casting speed.
No sharp difference was also observed between the upper and lower parts of the cast samples, regardless of the casting speed. All observations were obtained from longitudinal cross-sections of the samples as shown in Figure 6. SDAS is one of the most important factors affecting the mechanical properties of the Al-Si alloy family.
Furthermore, the SDAS of OCC samples is less than that of GC samples, regardless of the casting speed. Moreover, the average tensile strength of OCC specimens cast at 4.0 mm/s is also much higher (more than 2 times) than that of GC specimens. Nevertheless, the fracture strains of OCC samples are slightly higher compared to GC samples, regardless of the casting rates.
Notably, the improvements in the tensile strength of the OCC samples are due to the refinement of the SDAS, which can be a barrier to the movement of dislocations. SEM fractographs of tensile samples cast by GC and OCC are shown in Figure 12. Furthermore, the SDAS dendrites of α-Al in OCC samples is smaller than that in GC samples, regardless of the casting speed.
Study of mechanical properties of Al–Si–Cu alloy (ADC12) produced by different casting processes. Mater.
Analysis of Diﬀerent Solution Treatments in the
Materials and Methods
One hundred and twenty specimens were taken from the intermediate beam area of a number of 200 mm diameter slabs in the as-cast condition. There were 2 specimens after each solution treatment, one was used for metallographic inspection and material hardness calculation, while the other was used for differential scanning calorimetry (DSC). 3) The remaining 4 specimens were analyzed in the cast state. This treatment also allows the dissolution of the Mg2Si phase and the transformation of β-Al5FeSi particles into α-Al8(FeMn)2Si particles.
After a dissolution treatment at 600◦C with water cooling, which enabled the maximum dissolution of the Mg2Si phase and the transformation of β-Al5FeSi particles into α-Al8(FeMn)2Si particles. However, as the residence time of the treatment temperature increased, the atomic content of Mn in theα-Al(FeMn)Si particles also increased. The endothermic peak before the exothermic peak A indicates the dissolution of GP sites formed during natural aging.
Solution treatment at 600 °C with a residence time of 4 h resulted in the largest delay in the formation of metastable transition phases (peak A) and their subsequent dissolution (peak B). This paper analyzes the effect of solution treatment on the transformation of β-AlFeSi particles into α-(FeMn)Si particles and their possible effect on different aging treatments carried out in the range of 150–180 °C. 1) The Fe/Si atomic ratio increased with an increase in the solution treatment temperature from 550 to 600 ◦C. In the transformation of β-Al5FeSi particles into α-Al8(FeMn)2Si, a greater dissolution of Si and Fe atoms in the matrix was observed when the solution treatment was performed at 600 ◦C. The dissolution of Fe was slightly more pronounced when the retention times were increased from 2 to 4 hours. 3) At a solution temperature of 550 °C, the atomic ratio (Fe+Mn)/Si remained practically constant.
A model of beta-AlFeSi to alpha-Al(FeMn)Si transformation in Al-Mg-Si alloys. Mater. Development of Fe-based intermetallic phases in Al-Si hypoeutectic casting alloys: Influence of the Si and Fe concentrations and solidification rate.J. The dependence of beta-AlFeSi to alpha-Al(FeMn)Si transformation kinetics in Al-Mg-Si alloys on the alloying elements.
Simulation of the effect of composition on precipitation in Al 6xxx alloys during continuous heating DSC.J.
Simulation on the Eﬀect of Porosity in the Elastic Modulus of SiC Particle Reinforced Al
Square and circular morphologies of the particle reinforcement were selected to clarify the effect of the confinement shape in the elastic moduli. First, two validation studies are performed by comparing the FEA results with experimental data reported for the system Co-WC  and experimental results for porous refractory spinel . In the case of matrix reinforcement, the composite matrix is made of cobalt (Co), whose elastic modulus E2=207[GPa], Poisson ratioη2=0.31 and shear modulus G2=79[GPa].
For the Co-WC composite, experimental data of the elastic moduli as a function of the volume fraction of reinforcement varying at high concentrations of filler are available in the literature . Therefore, these data make it possible to investigate the reliability of the numerical model when considering high volume fractions. In the case of the porous material, the elastic modulus and Poisson's ratio of MgAl2O4Espinel= 268.2[GPa] and ν =0.262 respectively.
In the case of UCs models, the center of mass of the reinforcing particles coincides with the center of the square representing the matrix. The average strains, stresses, and strains in the composite are related to the limit displacements UC and REV by Gauss' theorem. Therefore, using appropriate boundary conditions to generate uniform stress and strain in a homogeneous medium, the relationship between the actual heterogeneous composite and the homogeneous medium is given by averaging the stress and strain tensors over the surface of the unit cell .
The eﬀective modulus of elasticity of the metal matrix composite (Ec) is given by substituting the two preceding equations in Hooke's lawσij=Ecεij. Figure 1 shows the boundary conditions and the periodic microstructures considered for the UCs models. The inclusion volume fraction is varied by changing the inclusion size in a parametric study where the circular reinforcement is only able to produce volume fraction up to 0.71%, while the square reinforcement a maximum volume fraction of 0.91%.
This latter study was performed to compare the mechanical behavior predicted by the UCs models with that predicted by the structure of real composite material.
Results and Discussions
Figure 5 shows the von Mises stress surface curves in the microstructure when the particle and matrix are fully dense. For a square SiC particle (Figure 5a), the maximum stress value in the particle is 0.4 Pa, this stress is mainly concentrated at the corners of the particle. In both cases, the stress in the particle is greater than the stress in the matrix.
Therefore, at low volume content of reinforcement, the shape of the particles has little effect on the modulus of elasticity of the composite. Considering the lower H–S limit, the circular shape in the numerical model underpredicts the modulus of elasticity of the Al-SiC composite. Figure 8 shows the von Mises stress surface curves in the microstructure when the Al matrix is porous and the SiC particle is fully dense.
Figure 10 shows the surface diagrams of von Mises stress in the microstructure when the Al matrix and the SiC particle are both porous. Figure 11 shows the surface diagrams of von Mises stress in the microstructure when the pore is located at the particle-matrix interface. Whereas, in the case of the circular particle, it carries much higher stress than matrix when the particle is porous, and the matrix is completely dense (Figure 9b).
Then, the transfer load is lowered and this is reflected as a reduction in the elastic modulus of the composite material (Figure 12, black line with circular markers). Figure 12 shows the results collected from the parametric study to demonstrate the effect of the porosity on the elastic modulus of the SiC particle reinforced Al matrix. The lower deterioration of the elastic modulus of the composite is for the porous particle and fully dense matrix.
Therefore, the data indicate that the porosity in the composite matrix is the main cause in the damage of the elastic modulus in MMC. Table 2 shows how increasing the % porosity in the matrix reduces the elastic modulus of the composite. Figure 14a,b show the surface plots of the von Mises stress in the Al matrix and SiC particles, respectively.