3.1. Thermodynamic Simulations
The equilibrium solidification diagram for the A356RE-A alloy is shown below in
Figure 3. The simulation suggests that, at room temperature, the alloy should consist of Al, eutectic Si, Al
4Ce
3Si
6, Al
11La
3, Al
5Cu
2Mg
8Si
6, Al
9Fe
2Si
2, Al
3Ti_LT, Mg
2Si, and Al
15(Fe,Mn)
3Si
2. The mass fraction of each phase is listed in
Table 2. In addition to these phases, some metastable phases were also observed to form during the solidification process, specifically AlCeSi, AlCeSi
2, LaSi
2_A1, LaSi
2_A2, and AlSi
3Ti
2.
Due to the high melting temperature of the La-Si phases, a small mass fraction of the LaSi
2_A1 phase was observed above 1000 °C (i.e., ~0.6%) which transforms into LaSi
2_A2 at 822 °C. Subsequently, the nucleation of the AlCeSi phase begins at 811 °C. This phase reaches a maximum mass fraction of 2.23% at 612 °C before temporarily transforming into AlCeSi
2 and then, finally, into Al
4Ce
3Si
6 which reaches a maximum mass fraction of 3.03%. The Al
4Ce
3Si
6 phase has been reported to improve the thermal stability of Al-Si-RE alloys [
1] and is one of the reasons for the impressive increase in elevated temperature strength of the A356 system [
14]. In addition to the Al
4Ce
3Si
6 phase, the presence of La in the RE mischmetal led to the formation of Al
11La
3. This phase has nearly identical properties as the Al
11Ce
3 phase, and its presence further improves the alloy’s thermal stability [
37].
Combating the positive benefits from the thermally stable Al
11La
3 and Al
4Ce
3Si
6 phases, ~0.74% of the harmful Al
9Fe
2Si
2 phase was observed. The Al
9Fe
2Si
2 phase is also commonly reported as Al
5FeSi. The nucleation of this phase begins at ~570 °C, which is followed by a rapid increase in the mass fraction to 0.53%. Subsequently, the mass fraction of this phase slowly increases until about 140 °C, after which the mass fraction remains relatively unchanged. Congruently, the more desirable Al
15(Fe,Mn)
3Si
2 phase begins to nucleate just prior to the onset of the Al
9Fe
2Si
2 phase, reaching a maximum mass fraction of 0.43%. This suggests that due to the low amount of Mn in the alloy (i.e., 0.1 wt.%), the Al
9Fe
2Si
2 phase is the preferential Fe phase to form. This supports the work done by N. Below et al. [
17].
The nucleation and solidus temperature of FCC-Al was observed to be 617 and 562 °C, respectively. The moderate solidification range (i.e., 55 °C) of FCC-Al suggests that porosity, caused by slow solidification, should not be a considerable issue.
Due to moderate Mg concentration in the A356RE-A alloy, it was expected that the Al9FeMg3Si5 would be observed in the solidification diagram but, instead, part of the Mg forms a solid solution with the Al matrix and the rest is present in the Q_Al5Cu2Mg8Si6 phase. Fortunately, the Q_Al5Cu2Mg8Si6 phase improves the thermal stability of the alloy and is much more desirable as compared to the Al9FeMg3Si5 phase.
The Scheil diagram for the A356RE alloy is shown below in
Figure 4. Due to the assumptions associated with Scheil solidification, the AlCeSi phase could not transform in the AlCeSi
2 or Al
4Ce
3Si
6 phases. Similarly, the growth of the Q_Al
5Cu
2Mg
8Si
6 phase was also halted by the assumption that zero diffusion occurs in the solid state. The non-equilibrium mode of solidification appears to have resulted in the formation of the Al
9FeMg
3Si
5 phase (i.e., Al
18Fe
2Mg
6Si
10 phase). The presence of this phase suggests that, depending on the cooling rate, this phase may or may not be present in the microstructures. However, it is likely that a combination of equilibrium and non-equilibrium solidification will occur, and thus, it is expected that some of this phase will be present at room temperature. The Scheil diagram indicates that this phase begins to nucleate at 556 °C and reaches a maximum mass fraction of 0.26% by 532 °C. In addition to this phase, the Al
9Fe
2Si
2 phase was also observed, however in a considerably lower amount (~0.43%).
The equilibrium solidification diagram for A356RE-B (high Mn, low Mg alloy) is shown below in
Figure 5. The results from the equilibrium simulation suggest that the chemistry modification greatly increases the volume fraction of the Al
15(Fe,Mn)
3Si
2 phase (i.e., from 0.42 to 1.75%) while reducing the mass fraction of the Al
9Fe
2Si
2 by ~65% (i.e., from 0.74% to 0.48%). The chemistry modification had a negligible effect on the mass fraction or solidification kinetics of the RE-containing phases (see
Table 3). Moreover, it was observed that nucleation and solidus temperature of FCC-Al increased by only ~5 °C.
Interestingly, the chemistry modification greatly reduced the amount of the Q_Al5Cu2Mg8Si6 phase (i.e., 0.15 vs. 0.98%) and it was observed that this phase reaches a mass fraction of 0% by just ~240 °C as compared to 340 °C in A356RE-A. The reduction in mass fraction is likely caused by the lower amount of available Mg. The reduction in Mg did not significantly affect the maximum mass fraction of Mg2Si (i.e., 0.37% vs. 0.43%) but it did slightly alter the solidification kinetics. In A356RE-A, it was observed that at room temperature, the mass fraction of Mg2Si was just 0.069%; however, for A356RE-B the mass fraction was 0.32%. It appears that the Al5Cu2Mg8Si6 forms at the expense of Mg2Si, and thus the lower volume fraction of Al5Cu2Mg8Si6 in A356RE-B results in the greater volume fraction of Mg2Si. Fortunately, the mass fraction of Mg2Si is nearly at its maximum at 250 °C, suggesting that its hardening contribution will also be at its greatest.
The Scheil simulation of A356RE-B reveals similar observations as described for the equilibrium simulation (see
Figure 6). The mass fraction of the Al
9Fe
2Si
2 was further decreased, reaching only 0.08%. Similar levels of the Al
15(Fe,Mn)
3Si
2 phase were also observed, reaching a maximum of 1.32%. However, a small mass fraction of the undesirable Al
9FeMg
3Si
5 phase was also present at the end of the Scheil simulation. Similar to the discussion for A356RE-A, it is expected that a combination of both equilibrium and non-equilibrium solidification will occur during the actual casting process.
The thermodynamic simulations of A356RE-B suggest that the modified Mn and Mg concentration should result in a lower concentration of harmful Fe-bearing intermetallics, while having little negative effect on the RE-bearing intermetallics or solidification characteristics of the alloy. As a result, it is expected that A356RE-B will be superiorly during the fitness-for-service testing.
3.3. Fitness-for-Service
To determine how the microstructural differences between A356RE-A and A356RE-B affect the fitness-for-service performance of the alloy system, elevated temperature tensile and creep tests were performed on the two alloys. The tensile results of both alloys at 250 °C are shown below in
Figure 11. The YS and UTS of the A356RE-A alloy were determined to be 62 and 84.3 MPa, respectively. As compared to the tensile results for the original as-cast A356 + 3.5%RE alloy reported in [
14], the new cleaning process (i.e., mechanical mixing + argon-purging) has notably increased both the YS (from 56 to 62 MPa) and UTS (from 70 to 84 MPa). In addition, the modulus of elasticity increased from 56.3 to 64.2 GPa, respectively. Powertrain components that operate at elevated temperatures require sufficient stiffness to maintain tight dimensional tolerances; however, these applications also benefit from some degree of ductility which provides automotive manufacturers with another type of safety factor (i.e., the component will deform before cracking).
As compared to the A356RE-A alloy, the modified A356RE-B alloy performed markedly better. Specifically, the YS, UTS, and modulus of elasticity improved by ~14%, 9%, and 10%, respectively. Moreover, the elongation also increased from 8.9 to 9.6%. As compared to the original A356RE alloy [
14], the YS (i.e., 70.5 MPa), UTS (i.e., 91.5 MPa) and modulus of elasticity (i.e., 70.9 GPa) improved by ~25%, 31%, and 26%, respectively. The improved tensile performance of A356RE-B is believed to be attributed to three events, the first being the transformation of the harmful needle-like Fe phase to the more favourable Al
15(Fe,Mn)
2Si
3. This phase transformation leads to lower stress concentrations and interfacial energy while increasing the contribution from Orowan strengthening. Similarly, the partial transformation of the AlSiRE phase from elongated plates to smaller, faceted Chinese script improves the strength via the Hall-Petch relationship (i.e., the smaller the precipitate, the greater the resistance against dislocation motion). The last contribution to improved strength is related to the noticeable decrease in the interparticle spacing of the eutectic Si particles (compare
Figure 7 and
Figure 9).
The staircase creep results for A356RE-A are shown below in
Figure 12, and the steady-state creep rates for each stress are listed in
Table 1. At 22 MPa, the
was measured to be 9.02 × 10
−10 s
−1. Such a slow
suggests that diffusion-based creep is the rate-controlling mechanism. It should be noted, however, that to accurately measure such low creep rates, a much longer time interval is required (on the order of several days to weeks). One of the reasons for this is due to the limited precision of extensometers as well as the natural force and thermal fluctuations that occur during creep tests.
Increasing the applied stress from 22 to 30 and finally 35 MPa resulted in only a small increase in the
, reaching a maximum of 2.66 × 10
−9 s
−1. The calculated apparent stress exponent from 22 to 35 MPa was
na = 2.4 (see
Figure 13). This suggests that the diffusional creep at 22 MPa is transitioning to viscous drag and climb [
39,
40]. Increasing the stress to 40, 45, and finally 55 MPa quadruples the apparent stress exponent to
na = 10.0, indicating that creep deformation bypassed power-law-governed dislocation climb, and transitioned to power-law breakdown dislocation creep. Such a high apparent stress exponent suggests that a threshold stress is present, likely present due to the high volume fraction of hard intermetallics in the alloy which pins dislocations [
30]. Plotting
1/n vs.
σ, with
n = 4.4, 5 and 7 revealed that
n = 7 results in the greatest linear fit (R
2 = 0.998), suggesting that deformation is governed by lattice diffusion-controlled creep with a constant sub grain structure. Following the procedure described in [
41], the threshold stress was determined to be ~34.3 MPa (see red arrows in
Figure 13). The following increase in the applied stress (i.e., to 65 MPa) caused rapid elongation; the sample failed in ~2 h. A true steady-state creep was not established; however, the brief minimum creep rate was determined to be 7.99 × 10
−6 s
−1. Such a large increase in the creep rate caused a near-instantaneous transition through dislocation glide into plasticity (
na = 38.9) [
42].
The effects from the chemistry modification (increased Mn, decreased Mg) are clear from the clear increase in the material’s creep resistance (see
Figure 12 and
Figure 13). Specifically, at every stress, the
was lower for A356RE-B than for A356RE-A (see
Table 6). Moreover, A356RE-B failed at a stress ~10–15 MPa greater than A356RE-A, reaching 100% of A356RE-B’s measured YS. At 65 MPa (i.e., the fracture stress of A356RE-A), the
for A356RE-B was more than a full magnitude lower than observed for A356RE-A, reaching just 1.12 × 10
−7 s
−1 as compared to 7.99 × 10
−6 s
−1. In fact, the creep rate for the A356RE-B was a well-established steady-state creep rate, as opposed to the brief minimum creep rate for A356RE-A.
The steady-state creep rate at 22 MPa was determined to be 6.57 × 10
−10 s
−1. Similar to the discussion for A356RE-A, accurate measurement of such a low strain rate is difficult without subjecting the stress for several days to weeks. Thus, there is expected to be some minor error associated with the measurement at 22 MPa. Once the stress on the A356RE-B alloy was increased from 22 to 30 MPa, the resulting creep rate was half that of A356RE-A (i.e., 1.03 × 10
−9 s
−1 vs. 2.00 × 10
−9 s
−1, respectively). The calculated apparent stress exponent was also lower, reaching only
na = 1.1 (as compared to
na = 2.4). Such a low
na suggests that diffusional creep is the rate-controlling mechanism. Further increasing the stress to 40 and 45 MPa only slightly raised the creep rate, reaching a maximum of 4.7× 10
−9 s
−1. Albeit only a small increase, the calculated apparent stress exponent increased to
na = 4.2, suggesting that the dominant creep mechanism has transitioned from diffusion creep to power-law-governed dislocation climb [
39,
40]. This correlates well with the deformation mechanism map (see
Figure 2). Increasing the applied stress from 55 to 65 MPa raises the
by a magnitude, reaching 3.32 × 10
−8 s
−1. As a result, the power-law breaks down and dislocation creep is the now rate-controlling mechanism (
na = 11.0) [
42]. This is dissimilar from the deformation mechanism map for pure Al (see
Figure 2). According to the deformation map and the calculated shear modulus (26.6 GPa), the power-law should breakdown with an applied stress of ~133 MPa. This difference is likely attributed to the high-volume fraction of precipitates in A356RE-B as compared to pure Al. Further increasing the stress to 70 MPa leads to a complete breakdown of the power-law, and as a result, the dislocation creep transitions to dislocation glide (
na = 21.2) [
42]. It should be noted that the
at 70 MPa was more than a full magnitude lower for A356RE-B than at 65 MPa for A356RE-A. Finally, once the stress is increased to 80 MPa, the creep enters the plasticity regime where the sample begins to neck and, ultimately, leads to failure.
Similar to the A356RE-A alloy, the high apparent stress exponents suggest that there is a threshold stress in the A356RE-B alloy. Plotting 1/n vs. σ reveals good linear fits for all three stress exponents, with n = 7 having the greatest fit at R2 = 0.983, leading to an estimated threshold stress of 44.2 MPa.
The chemistry modification has clearly led to an increase in the creep resistance of the alloy. Not only has the threshold stress increased by ~10 MPa, the steady-state creep rates are lower at every stress. This is likely attributed to the higher volume fraction of intermetallics available for pinning dislocations, the transformation of the brittle Al5FeSi and Al9FeMg3Si5 into Al15(Fe,Mn)3Si2, as well as the size reduction and morphology modification of the AlSiRE intermetallic.
Similar to the results presented by A. Farkoosh and M. Pekguleryuz [
33] who studied the effects of Mn on the creep resistance of Al-Si-Cu-Mg-Ni alloys, the increased Mn addition in the A356RE alloy not only lowered the steady-state creep rates for all stresses but also raised the fracture stress. The driving factors behind the observed improved strength and creep resistance is believed to be the transformation of Al
5FeSi and Al
9FeSi
3Mg
5 into Al
15(Fe,Mn)
3Si
2 phase, as well as the partial transformation of the elongated AlSiRE plates into a smaller, faceted Chinese script morphology. The lower aspect ratio, a stronger bonding with the Al matrix, and the more favorable morphology of Al
15(Fe,Mn)
3Si
2 phase lead to a decrease in the precipitate-matrix interfacial energy, lower the development of stress concentrations and improve the alloy’s ductility. Moreover, this transformation increased the total volume fraction of intermetallics, leading to a greater number of obstacles available to pin dislocation motion (Orowan strengthening), and thereby, further improving the creep resistance of the alloy.