Figures 1A shows the schematic representation of the free vibration test. A specimen cantilever beam was subjected to dynamic testing. The frequency spectrum of each specimen was triggered by hitting the free end of the bottom side with a rubber tip hammer (Type 8206, Bruel & Kjaer) and received by a piezoelectric accelerometer (Type 4507-B, Bruel & Kjaer, Naeuram, Denmark). A LAN-XI pulse analyzer was used to collect the vibration signals (type 3050 A-60, Bruel & Kjaer). The Pulse Lab Shop and MEscope (Vibrant Technology Inc. Scotts Valley, CA, United States) The software was used to process the signal after acquisition. The FFT was computed using a frequency range of 10 k Hz and a resolution of 3,200 lines. In a linear mode, ten data blocks were averaged together. All samples were analyzed using an average of five samples for each alloy. In order to calculate the damping ratio (ζ), the first resonant frequency (F_n), and the dynamic modulus (E_d), the following equations were used depending on the free vibration decay curves of a cantilever beam (Moustafa, 2018; Ahmed et al., 2021; Moustafa and Almitani, 2021). δ = 1 n ln x o x n (1) ζ = δ 4 π 2 + δ 2 (2) E d = 16 π 2 M + 0.236 m L 3 f n 2 w t 3 (3) E l o s s   m o d u l u s   = E d × η (4) Where X _o and X _n are the amplitude located form the time domain at a number of cycle n as shown in Figures 1B. M is the mass of the accelerometer and mounting clips of 6.0 × 10⁻³ kg, the distributed mass of the cantilever beam m is 10.0 × 10⁻³ kg), l is the beam length (0.13 m), f _n is the first resonant frequency that determined from both method the logarithmic decrement curve calculations and the frequency response function (FRF), w (0.01 m) and t of (0.003 m) are the width and thickness of the cantilever beam and η is the loss factor = complex loss modulus/storage modulus. The loss factor is obtained from the half bandwidth method from FRF curve (Hammad and Moustafa, 2020; Basha et al., 2022).

The nucleation and development of α-Al during solidification determine the grain size of the aluminum alloy. It has been shown that boron, when combined with an excess of titanium, is more efficient than boron alone (at values of 0.005%–0.1%) in the role of grain refiner during solidification. In commercial grain refineries, titanium and boron are often found in a 5-to-1 ratio. In this step, grain refiner is added to the aluminum alloy melt to create massive heterogeneous nucleation sites of α-Al. Solution components that cause constitutional supercooling can also accelerate heterogeneous nucleation. When α-Al nuclei form, the solution components (or insoluble particles) can clump and hinder their development. As a result of the high nucleation rate and the constraints imposed on grain development, the grain size is reduced. Optical microstructure pictures of the standard A242 aluminum alloy, as well as the modified A242 + TiB aluminum alloy, are shown in Figure 2. When TiB modifiers were added to the casting process, it resulted in a refinement of the grains, which led to an equiaxed refined dendritic structure. Without undergoing any extra processing, the improved A242 + TiB alloy displayed grains that were homogenous, equiaxed, and refined to be three times smaller than those of the standard alloy. The standard A242 alloy was analyzed using ImageJ software, and the average grain size was observed as 470 ± 80 μm, while the modified A242 + TiB has a 35 ± 5 µm average grain size. The A242 alloy, much like the reference alloy, had big grains and a typical coarse dendritic structure. The grain size of the dendrites had a substantial impact on both the geometry of the dendrites and the structures that formed after solidification. Adding 0.50% titanium and 0.10% boron to alloy A242 modified the material’s grain shape and phase structure. These modifications included the transformation of coarse dendrites into homogenous and refined primary α-Al grains.

Scanning electron microscopy (SEM) was employed to analyze the chemical composition and the formed phase in the investigated alloys. Figure 3 shows the SEM micrographs and the EDS maps for different elements for reference and modified alloys. The EDS maps confirm the uniform distribution of the interdimeric phases along the dendritic arms. For the modified alloy with Ti-B, the Ti element appeared due to its suitable weight percentage in the alloy, while the B element did not appear due to its lower percentage of 0.1% and its light atomic weight, as shown in Figure 1C, D. Figure 4 shows the EDS point for the formed phases in the standard and modified alloys. The EDS analysis confirms the presence of the metallic phases of Al₂Cu, Mg₂Si, Al₃Ni, and Al₉FeNi based on the α-Al solid solution in the reference alloy. These results are in agreement with (Rodríguez et al., 2012). For the modified alloy, the presence of Ti-B resulted in an equiaxed dendritic structure surrounded by the eutectic phases (Figures 2B, 3B, 4B). The main reason for the refinement of the dendritic structure is the presence of the heterogenous enucleation centers TiB₂ and Al₃Ti which in agreement with (Mosleh et al., 2021).

Figure 5 presents a comparison of the compression strength and elongation findings in tested samples; thus, the standard and modified A242 alloy were evaluated at both room temperature (RT) and 250°C. At room temperature, the compressive yield strength of (σ_0.2) for the reference and the modified alloy was 213 ± 7 MPa and 300 ± 10 MPa, respectively. The improvement in the compressive yield in the modified alloy with Ti-B can be attributed to the refined structure, which increased the boundaries of the grains and acted as a barrier for the dislocation movements (Pattnaik et al., 2015). The ultimate compressive strength (CUS) value of the modified alloy reached 510 MPa at room temperature. These results are in agreement with (Kurt et al., 2015; Wang et al., 2021), which attributed to the refining effect of the TiB. As shown in Figures 13B, the CUS of this modified alloy had a greater value than that of the conventional A242 alloys used in manufacturing automobile component parts.

Compared with the other Al-Cu-Mg alloys, 242 alloys with 2% Ni and 1%Fe exhibit high compressive strength (Trejo Rivera et al., 2019). Due to iron’s limited solubility in the solid state (0.04%), most of the iron contained in aluminum exceeds this proportion as an intermetallic second phase with aluminum and possibly other elements such as Al₃Ni; moreover, with 1.0% iron, the strength of the alloy is increased. Solid nickel is 0.04% soluble in aluminum; hence, the two elements seldom interact. At this point, it forms an insoluble intermetallic complex with iron as Al₉FeNi. Nickel boosts hardness and strength at high temperatures and minimizes expansion in A242 alloy. Due to the grain refinement and redistribution of the alloying elements in the modified A242 alloy, the CUS increased by 1.45 times that of the standard A242 alloy at room temperature. As mentioned in the introduction section, the 242 aluminum alloy is widely used in cylinder heads for motorcycles and pistons; thus, it is important to maintain the mechanical properties of this alloy at an acceptable level. In this context, the compressive test at 250°C was performed for both standard and modified alloys. Increasing the temperature, both compressive yield and ultimate strengths decreased (Gündüz and Acarer, 2006). Increasing the temperature facilities the dislocation movements resulting in low characterizations. At 250°C, the σ_0.2 and σ_CUS decreased to 194 ± 6 MPa and 315 ± 10 MPa, respectively. There is a direct correlation between refined granules and mechanical qualities; hence, the findings revealed that granule reduction enhanced the distribution and size of the alloying components and generated phases during the casting process, which is reflected in mechanical behavior. Effects of grain refiners on certain properties Alloy A242 It has been found that adding an appropriate number of grain refiners improves compressive yield and ultimate strengths compared with the reference alloy at 250°C. The most important explanation for this improvement is probably the smaller equiaxed dendrites’ size, which leads to a finer distribution of the second phases (i.e., the intermetallic phases) surrounding the dendrites (Kobayashi, 2000). The grain refinement provides a more uniform distribution of elements in the microstructure by creating heterogeneous nucleation sites with Ti₂ and Al₃Ti (Murty et al., 2002; Fan et al., 2015). In general, most aluminum alloys’ compressive yield points fall below 100 MPa as the temperature is raised above 250°C, while in the standard and modified alloys, the compressive yield strength was 194–230 MPa at 250°C. The quality index Q is represented by the equation Eq. 5 and is defined as a semilogarithmic plot of ultimate tensile strength versus elongation to fracture. (Wang et al., 2010; Khorshidi et al., 2011): Q = U T S   M P a + Log   ε % (5)

Table 2 shows the quality index for each alloy studied. It can be seen that the refined alloy Ti-B achieves the highest quality index. The main reason for this improvement is that the grain size of the castings has decreased, resulting in a more uniform distribution of the finer second phases. The microhardness values of the investigated alloys revealed that the TiB modifier elements have a significant effect on the improvement of the overall hardness. The modified alloy increased by 20 HV more than the as-cast alloy, as shown in Figure 6. The main reason for increasing the hardness can be attributed to the refining effect of the TiB and the uniform distribution of the second phases in the modified alloy.

The study of dynamic properties is one of the most critical pillars used during the design stages of products that are exposed to high operational load and conditions. The natural frequency and damping capacity are among the most important features that determine the dynamic behavior of materials. This study experimentally performs a free vibration impact test to identify the dynamic behavior. The time and frequency domains are used to analyze the acquired data, as shown in Figure 7; the real amplitude with corresponding Fast Fourier Transform FFT is represented for each sample. Figure 7A shows the as-cast A242 alloy acoustic behavior through different frequencies, while Figures 7B shows the modified A242 + TiB alloy. The vibration signals revealed that as-cast alloy excites more resonant frequencies in the frequency range between 80 and 1,200 Hz. In contrast, the modified A242 + TiB alloy absorbed the resonant frequency in the same range. The resonant frequencies at 680 and 1,269 Hz are observed only in standard A242 alloy; thus, they can be attributed to the random distribution of the intermetallic phases and extra-large grains during the solidification process. The absence of these resonant frequencies in the modified alloy means that the homogeneity of the modified A242 + TiB alloy is well-created. Through the use of the time domain curve, dynamic properties were able to be determined, as seen in Figure 8. Calculations are performed on the data using the initial resonant frequency as the basis (first mode). In Figure 9, we see a comparison of the frequency response functions (FRFs) that were calculated using FFT analysis applied to the time domain for both of the studied alloys. The new resonance peaks developed in the 600–800 Hz region that corresponded to the standard alloy A242. This points to the fact that the structure of the coarse grain and the random production of the intermetallic phases create greater resonant frequencies in the cast than in the homogenized modified alloy. Table 3 provides a summary of the dynamic characteristics of the investigated alloys. The results revealed little change in the dynamic properties of the investigated samples; however, the damping ratio increased due to the refinement of the grain, which reflects the metal’s responsiveness to the external effect of mechanical vibrations. Table 3: Dynamic Characteristics of Investigated Alloys.

Finite Element Model:

A numerical two-dimensional thermomechanical FE model was constructed to measure the plastic behavior of the alloys at various temperatures. A strip of the alloys is being pressed by a press over a die, as shown in Figure 10. The strip is 100 mm long, 1 mm thick, and held in place by a pressure pad. The press, pressure pad, and die all have 5 mm fillets around the edges. The mechanical properties from Figure 5 for both A242 and A242 + TiB alloys were used in the FE simulation. The FE simulation was run multiple times while having the strip preheated to different temperatures. The temperatures used were RT, 50, 100, 150, 200, and 250^oC. The strip was modeled using 2,500 two-dimensional quadratic plain strain elements with a total of 8511 nodes. Figure 11 shows a close-up of the developed FE mesh of the strip. For the boundary conditions, the strip was prevented from moving along the horizontal x-direction along surfaces a-b. The press, pressure pad, and die were all treated as rigid bodies with reference points RP ₁ , RP _2, and RP ₃ for the boundary conditions. The press was prevented from movement and rotation in all directions at RP ₁ except the vertical y-direction, where it was forced to move 30 mm downward.

Similarly, the pressure pad was prevented from movement and rotation in all directions at RP ₂ except the vertical y-direction, where it was subjected to a force of value 440 KN acting downwards. As for the die, it was completely fixed in place at RP ₃ . A contact interaction was defined between the die and the strip’s bottom surface with a coefficient of friction of 0.3. A similar interaction was defined for the strip’s top surface and the pressure pad. As for the contact between the press and the strip’s top surface, frictionless contact was used.

Figure 12 shows the deformed shape of the strip from the FE simulation. A close-up view of the developed stresses on the strip after the forming process using both A242 and the modified A242 + TiB alloys at 250 °C is shown in Figure 13. Figures 13A shows the stresses developed on the top half of the A242 strip, while Figures 13B shows them on the bottom half of the alloy. Similarly, Figure 13C, D shows the developed stresses for the top and bottom halves of the A242 + TiB strip. Figures 14, 15 show the developed stresses at the contact region between the strip and other rigid bodies for both alloys when the press moves by 5, 10, and 15 mm, respectively. To have an overall view of the stresses developed on both A242 and A242 + TiB alloys, Figure 16 shows the stresses developed on the strip’s top and bottom surfaces at different locations for both alloys at room temperature (RT) and 250°C. At location A, the stresses are at their lowest points on the top surface but higher on the strip’s bottom surface. The temperature difference barely shows any effect on that location on both sides. At location B, the stresses on the top surface of the strip are still lower than on the bottom surface, but the effect of the temperature difference is visible. When at room temperature, the required stresses to deform both alloys are higher than at 250°C, which is expected due to the higher ductility of the material and strain softening at higher temperatures. At location C, the stress values reach their highest values. Furthermore, the stresses on both surfaces allow the top and bottom surfaces to be almost identical at higher temperatures. But, at room temperature, the top surface exhibits higher stresses due to the friction between the die and the strip’s bottom surface. This causes the top surface to strain more than the bottom surface of the strip. At location D, the stress values drop when the strip is relieved from the bending stresses present at location E as it slides over the die. The stress values at location F are very small since the strip is sliding over the frictionless surfaces of the die and pressure pad.

Further simulations were conducted at various temperatures for both alloys to determine the stress variations at temperatures between RT and 250°C. Figures 17A, B shows the developed stresses for A242 at these temperatures, while Figures 17C, D shows the developed stresses for A242 + TiB. As expected, the increase in temperature causes a reduction in stresses due to the reduced ductility of the materials. In all figures, the stresses at locations A, D, and F are almost unaffected by the increase in temperature. While locations B, C, and E are where the effect of raising the temperature is more prominent, This shows that the alloy A242 + TiB is more resistant to deformation than A242, even at high temperatures, making it suitable for highly elevated applications. At room temperature, the A242 + TiB alloy was more resistant to deformation than A242 by 9.1%. At 250 C, A242 + TiB was 7.6% more resistant.