Evaluation of Microstructure and Tensile Properties of Al-12Si-4Cu-2Ni-0.5Mg Alloy Modified with Ca/P and TCB Complex

Abstract

The room-temperature and high-temperature microstructural characteristics and tensile properties of an Al-12Si-4Cu-2Ni-0.5Mg piston alloy modified with calcium (Ca; denoted as AC sample) or phosphorus (P; denoted as AP sample) under different heat treatment conditions were systematically analyzed. Under Ca modification, the second-phase network structure of the alloy was adjusted and strengthened by an Al-TCB master alloy. Eutectic silicon (Si) particles in the AC sample possessed a fibrous structure, whereas the AP sample contained elongated eutectic Si particles, and Ca modification was found to be a potential method for simultaneously enhancing the strength and plasticity of the alloy to a matching degree at high temperatures. The T6 treatment noticeably increased the density of nanoscale precipitates; however, it also disrupted the growth of the second-phase network structure. Micron and submicron C-TiB₂ and Al₄C₃ particles formed by the in-situ reaction of TCB particles acted as bridging phases within the second-phase network structure and enhanced the strength of the piston alloy. The ultimate tensile strength of the alloy at 350 °C increased from 74 to 101 MPa, representing a 36.5% enhancement. A comprehensive analysis revealed that Orowan strengthening was the main strengthening mechanism of the alloy at room temperature, whereas load transfer and network structure strengthening were the dominant strengthening mechanisms at high temperatures.

1. Introduction

In order to realize the lightweight design requirements for automobiles and ships, eutectic and near-eutectic Al-Si alloys are widely used in engine block and piston manufacturing due to their low density, high strength, outstanding castability, and superb wear resistance [1,2,3]. Near-eutectic Al-Si-Cu-Ni-Mg alloys have become the preferred materials for piston manufacturing because of their excellent strengthening phases, such as Al₂Cu, Mg₂Si, Al₅Cu₂Mg₈Si₆, Al₃Ni, Al₃CuNi, and Al₇Cu₄Ni [4,5,6]. The morphology of Si particles and intermetallic compounds in Al-Si alloys plays a crucial role in their mechanical properties [7,8]. In as-cast Al-Si alloys, the Si phase possesses a blocky and needle-like shape [9]. When Al-Si alloys are subjected to external forces, they become prone to splitting and cracking, resulting in the deterioration of their mechanical properties [10,11,12], and large Si particles are mainly responsible for high-temperature fatigue crack initiation [13,14,15]. Hence, certain elements are often added to Al-Si alloys to change the morphology of the Si phase.

Common modifying elements for gravity-cast Al-Si piston alloys include sodium, strontium, phosphorus, and calcium. The modification effect of sodium (Na) is prone to rapid fading, and its addition increases the gas absorption tendency of the melt, leading to the formation of microscopic porosity in castings [16]. Strontium (Sr) is considered a superior modifier to sodium, exhibiting relatively better modification effects and longer-lasting effectiveness, hence its widespread application in cast Al-Si alloys [16,17,18]. However, the addition of Sr still increases the air absorption of the Al-Si alloy melt, which increases the number of pores in the casting and decreases the density of the casting [19,20]. Phosphorus (P) can not only effectively refine the primary silicon but also play a key role in the nucleation of eutectic silicon. Therefore, it is widely used as a modifier in the production of Al-Si piston alloys [21,22,23]. Calcium (Ca), an element from the same group as Sr, has also been shown to effectively modify eutectic Si in many studies, and eutectic silicon particles can change from coarse plate-like structures to fine fibrous silicon particles [24,25]. Although there have been numerous studies on the modification potential of Ca in Al-Si alloys [24,25,26], the effects of Ca modification on the microstructure and mechanical properties of Al-Si-Cu-Ni-Mg piston alloys are rarely reported.

Sui et al. [16] found that after adding Sr to the Al-12Si-4Cu-2Ni-0.8Mg alloy, the morphology of eutectic silicon changed from needle-like to fibrous. When the amount of Sr added increased from 0% to 0.02%, the room-temperature ultimate tensile strength (UTS) of the alloy increased from 196 MPa to 249 MPa, and the elongation (EL) rate increased from 0.74% to 0.95%. Pistons usually work in high-temperature environments; therefore, in the design of piston alloys, in addition to room-temperature mechanical properties, their high-temperature performance should also be considered. Hu et al. [6] studied the tensile properties of the Al-12Si-1Cu-1Ni-1Mg alloy with different microstructures and found that alloys with a dispersed particle distribution of the second phase had higher strength at room temperature, while alloys with a semicontinuous network distribution of the second phase at high temperature had better strength. Zuo et al. [27] designed different volume fractions of the δ-Al₃CuNi phase in the Al-Si-Cu-Ni alloy through thermodynamic calculations and found that with an increase in the δ-Al₃CuNi phase, the ultimate tensile strength and yield strength of the alloy increased. Wang et al. [28] investigated the effect of microstructure on the tensile and low cycle fatigue behavior of the Al12Si4Cu3NiMg alloy using conventional and in situ testing techniques. For gravity-cast alloys, microcracks rapidly propagate through eutectic Si and intermetallic compounds. By using ultrasonic melt treatment to refine the microstructure, higher crack propagation resistance can be achieved.

The TCB complex consisting of B-doped TiC (B-TiC) and C-doped TiB₂ (C-TiB₂) in the Al-TCB master alloy exhibits excellent resistance to Si/Zr poisoning; thus, it is widely used in the strengthening and toughening of Al-Si alloys. A study by Li et al. found that adding 0.1% Al-TCB master alloy refined the average grain size of Al-7Si-0.4Mg and Al-5Cu-0.15Zr from 784 ± 58 μm and 402 ± 32 μm to 84 ± 7 μm and 63 ± 4 μm, respectively [29]. Research by Pan et al. showed that after the T6 treatment, the room-temperature UTS of a sand-cast ZL114A alloy modified with Sr and the Al-TCB master alloy increased to 290 ± 5.4 MPa, representing an improvement of approximately 71.6% compared to the Sr-modified ZL114A alloy [30]. Furthermore, Li et al. introduced the Al-TCB master alloy into an Al-5Cu alloy and discovered that adding 0.3% Al-TCB master alloy not only effectively refined the α-Al grains but also modulated the network structure of the second phase at grain boundaries, enabling it to withstand loads more effectively. The maximum increase in the high-temperature tensile strength of the alloy reached 28% [31].

The present study explored the relationship between the room-temperature and high-temperature (350 °C) microstructural characteristics and tensile properties of the Al-12Si-4Cu-2Ni-0.5Mg alloy modified with Ca or P under different heat treatment conditions. Under Ca modification, the second-phase network structure was tailored by the Al-TCB master alloy to further increase the high-temperature strength of the piston alloy. The strengthening mechanisms of the alloy were analyzed to provide a reference for the design of heat-resistant aluminum alloys.

2. Materials and Methods

The Al-12Si-4Cu-2Ni-0.5Mg alloy was used in the experiments. The raw materials included the following: commercial pure aluminum (99.7%; unless otherwise stated, all compositions are in wt.%); pure silicon (99.85%); and the Al-50Mg, Al-50Cu, Al-30Ni, Al-10Ca, Al-4P, and Al-TCB (Al-2Ti-0.3C-0.2B) master alloys. All the materials were supplied by Shandong Al & Mg Melt Technology Co., Ltd. (Jinan, China).

The alloy was melted in a graphite–clay crucible using a resistance furnace according to the following specific preparation procedure. Commercial pure aluminum was first melted at 780 °C, and then pure Si was added to the molten aluminum and allowed to dissolve completely. Subsequently, the Al-50Cu and Al-15Ni master alloys were added sequentially. At 730 °C, the Al-50Mg master alloy was added. After complete melting, the melt was stirred thoroughly and held at temperature for 15 min. Subsequently, 0.6% C₂Cl₆ was added for refining and slag removal. For Ca modification, 0.6% Al-10Ca (AC alloy) was added, while for P modification, 1% Al-4P (AP alloy) was added. In addition, for the ATCB alloy, 0.6% Al-10Ca and 20%Al-TCB were added. Finally, the melt was poured into a permanent metal mold preheated to 300 °C to obtain the as-cast samples, as shown in Figure 1a. The cooling rate was 2–4 °C/s.

The dimensions of the tensile specimens at room temperature (GB/T 228.1-2021 [32]) and high temperature (GB/T 228.2-2015 [33]) used in this study are shown in Figure 1b and Figure 1c, respectively, and photos of the tensile specimens are shown in Figure S1 (in the Supplementary Materials). ZC (water quenching at about 500 °C after solidification, aging at 180 °C for 6 h) and T6 (solid solution at 500 °C for 4 h, water quenching, aging at 180 °C for 6 h) treatments were carried out for the alloy, and a schematic diagram of the heat treatment process is shown in Figure S2 (in the Supplementary Materials). The AC, AP, and ATCB alloy samples treated by ZC heat treatment are marked as AC-ZC, AP-ZC, and ATCB-ZC. The AC, AP, and ATCB alloy samples treated by T6 heat treatment are AC-T6, AP-T6, and ATCB-T6.

The actual chemical composition of the samples was determined using Spark emission spectrometry (QSN750, OBLF, Dortmund, Germany), as presented in

Table 1. Chemical compositions of alloy samples (wt.%).

Samples	Si	Cu	Mg	Ni	Fe	Ca	P	Ti	B	C	Al
AC	11.41	3.67	0.52	1.92	0.17	0.052	-	-	-	-	Bal.
AP	11.47	3.61	0.58	1.91	0.16	-	0.026	-	-	-	Bal.
ATCB	12.03	3.64	0.68	2.04	0.21	0.059	-	0.404	0.031	0.039	Bal.

. A phase analysis of the alloys was performed using an X-ray diffractometer (XRD, DMAX-2500PC, Rigaku Corporation, Akishima, Japan) with a scanning speed of 10°/min over a 2θ range of 10° to 90°. The microstructure of the alloy samples and the tensile fracture surfaces were examined using a field emission scanning electron microscope (FE-SEM, SU-70, Hitachi Limited, Tokyo, Japan) equipped with energy-dispersive spectroscopy (EDS). The bright-field (BF), high-resolution TEM (HRTEM), and selective area electron diffraction (SAED) images of the precipitates were characterized by a transmission electron microscope with EDS (TEM, Talos F200X, Thermo Fisher Scientific, Waltham, MA, USA).

Tensile property tests at room temperature (25 °C) and high temperature (350 °C) were carried out on a WDW-100D universal testing machine. Tensile rates were 0.5 mm/min (room temperature) and 2 mm/min (high temperature), and at least three samples were tested for each alloy. During high-temperature tensile testing, a thermocouple was attached to the specimen to ensure it reached the test temperature. Testing commenced after the specimen temperature reached 350 °C and was held for 15 min.

3. Results and Discussion

3.1. Microstructures of AC and AP Samples

Figure 2a–d display the SEM morphologies of the as-cast Ca-modified and P-modified Al-12Si-4Cu-2Ni-0.5Mg alloy samples (denoted as AC and AP, respectively). In both the AC and AP samples, α-Al grain boundaries contained gray lamellar, gray fibrous, bright white blocky, and white fishbone-shaped phases. The microstructures of both samples possessed a semi-network structure [6]. The XRD results of the AC and AP samples are presented in Figure 3 and

Table 2. EDS compositions (at. %) and corresponding phases indicated in Figure 2.

Point	Al	Si	Cu	Ni	Mg	Fe	Phase
1	15.8	83.9	0.3	-	-	-	eutectic Si
2	24.8	32.5	9.1	0.4	33.3	-	Al5Cu2Mg8Si6
3	58.8	-	30.4	10.8	-	-	Al7Cu4Ni
4	66.7	1.2	31.5	0.7	-	-	Al2Cu
5	57.7	-	23.2	18.3	-	0.8	Al3CuNi
6	15.8	83.9	0.3	-	-	-	eutectic Si
7	2.4	97.6	-	-	-	-	primary Si
8	58.9	-	31.6	9.6	-	-	Al7Cu4Ni
9	66.7	1.4	31.4	0.6	-	-	Al2Cu
10	59.0	-	20.4	20.3	-	0.3	Al3CuNi

, and Figure 4 presents their EDS elemental mapping results. It is discernible from Figure 3 that the main phases in both alloy samples included Al, Si, Al₂Cu, Al₃CuNi, Al₇Cu₄Ni, and Al₅Cu₂Mg₈Si₆; thus, no significant difference in the compositions of their intermetallic compounds was observed.

In the AC sample, fibrous eutectic Si particles were primarily distributed along the grain boundaries of α-Al dendrites and formed eutectic colonies with the α-Al matrix. In the Al-Si eutectic system, the formation of eutectic colonies mainly results from the coupled growth of the α-Al matrix and eutectic Si particles [34]. In contrast, plate-like eutectic Si particles predominantly appeared in the AP sample, and some of these particles also exhibited intergranular penetration growth. It is noticeable from Figure 2 that the white fishbone-shaped Al₃CuNi phase in the AC sample was finer and more uniformly distributed. Furthermore, the Si-containing (eutectic Si and Al₅Cu₂Mg₈Si₆) and Ni-rich (Al₃CuNi and Al₇Cu₄Ni) phases in the AC samples were interconnected and uniformly distributed along the grain boundaries of α-Al dendrites, forming a relatively continuous network structure. In the AC alloy sample, the isothermal solid–liquid interface led to eutectic colony formation [16], and consequently, eutectic colony boundaries became enriched with Cu and Ni during solidification. Therefore, intermetallic compound phases were mainly distributed along these eutectic colony boundaries and formed a network structure. The EDS mapping results of the AC and AP samples presented in Figure 4 provide a clearer visualization of this phenomenon.

In the solidified microstructure of an alloy, the secondary dendrite arm spacing (SDAS) is an important parameter [16]. Intermetallic compounds are generally distributed along grain boundaries and form a network structure. The smaller the SDAS, the denser the network of secondary phases [24]. It is evident from Figure 2 that α-Al dendrites in the AC sample were finer than those in the AP sample. In order to quantitatively compare the dendrite sizes of the two samples, their SDAS values were measured using the linear intercept method [35,36,37].

SDAS = L/n

(1)

where L is the length of the straight line drawn from one secondary dendrite arm to another, and n represents the number of secondary dendrite arms passed through this line. The SDAS values for the AC and AP samples were determined as 19.8 and 27.4 μm, respectively.

Al-Si-Cu-Ni-Mg alloys are generally strengthened by heat treatments. Prior to tensile property tests, both the AC and AP alloy samples were subjected to ZC and T6 heat treatments. The microstructural morphologies and EDS results of the alloy samples after heat treatments are presented in Figure 5 and Figure 6, respectively. It is observable from Figure 5 that the microstructures of both samples after the ZC treatment remained consistent with their as-cast structures. In contrast, significant changes occurred in the eutectic Si phase after the T6 treatment, and a considerable amount of eutectic Si particles was dissolved into the α-Al matrix. In the AC-T6 sample, fibrous eutectic Si particles were transformed into granular ones, and the network configuration at grain boundaries was partially disrupted, whereas in the AP sample, plate-like eutectic Si particles changed to short rod-shaped ones. No distinct morphological difference in the bright white Ni-rich phases of the ZC-treated, T6-treated, and as-cast samples was observed, indicating that heat treatments had minimal impact on these Ni-rich phases. The EDS elemental maps of the AC and AP samples after the ZC heat treatment are presented in Figure 6a and Figure 6b, respectively. It is noticeable that the continuous bar-like structure of the Al₅Cu₂Mg₈Si₆ phase changed to a discontinuous bar shape, exhibiting partial dissolution in the α-Al matrix and fragmentation. This phenomenon was more pronounced in the T6-treated samples, where the Al₅Cu₂Mg₈Si₆ phase appeared as short rods and small blocks (Figure 6c,d).

Generally speaking, a large number of precipitated phases can appear in the alloy after heat treatment, which has a significant impact on the microstructure and properties of the alloy [38,39]. In order to gain deeper insights into the microstructural transformations of the AC-ZC and AC-T6 samples during heat treatments, TEM analysis was conducted (Figure 7). Figure 7a–c present the TEM results for the AC-ZC sample. Figure 7a reveals the presence of black blocky and rod-shaped nanophases in the alloy, and these nanophases (precipitated during solidification) were identified as Si from the EDS mapping results presented in Figure 7b. Figure 7a also indicates the existence of a small amount of black needle-shaped precipitates (identified as the β″-Mg₅Si₆ phase) in the AC-ZC sample, and the HRTEM image and corresponding FFT pattern of these precipitates are exhibited in Figure 7c. Figure 7d displays the TEM morphology of the AC-T6 sample, and the presence of a high density of black needle-shaped precipitates (identified as the β″-Mg₅Si₆ phase) was detected in this sample. Figure 7e and Figure 7f present the HRTEM images and FFT patterns of these needle-shaped precipitates along the [110]_Al and [100]_Al axes, respectively. It is noticeable from Figure 7d that nano-Si phases similar to those observed in Figure 7a were not detected in the AC-T6 sample, and the number of β″ precipitates also significantly increased. It is well proven that nanoscale Mg-Si phases precipitate during the aging process in Al-Si-Mg alloys and act as strengthening phases [40,41,42]. During deformation, the movement of dislocations is hindered by β″ phases, resulting in coherency strain, order strengthening, and modulus mismatch strengthening, which contribute to the deformation resistance of the alloy [43].

3.2. Tensile Properties of AC and AP Samples

Figure 8 presents the engineering stress–strain curves of the alloy samples, and the corresponding tensile property data are listed in

Table 3. Room- and high-temperature tensile properties of AC and AP samples.

Temp/°C	Samples	UTS/MPa	YS/MPa	EL/%	FR/%	QI/MPa
25	AC-ZC	293.2 ± 3.1	270.5 ± 5.1	0.71 ± 0.07	-	306
	AP-ZC	327.9 ± 2.3	310.0 ± 4.0	0.63 ± 0.04	-	341.5
	AC-T6	353.9 ± 6.2	317.5 ± 3.4	0.87 ± 0.02	-	361
	AP-T6	370.3 ± 4.2	330.5 ± 3.3	0.75 ± 0.05	-	368
350	AC-ZC	74.0 ± 0.5	59.2 ± 1.1	7.16 ± 0.24	74.76	417.2
	AP-ZC	75.5 ± 1.5	61.3 ± 2.0	6.84 ± 0.31	76.97	403.3
	AC-T6	78.0 ± 1.0	69.0 ± 1.0	8.95 ± 0.22	77.96	516.5
	AP-T6	80.1 ± 1.0	71.0 ± 1.0	8.31 ± 0.28	78.38	486.5

. In addition, the true stress–strain curve of the alloy samples before the ultimate tensile strength was plotted, as shown in Figure S3 (in the Supplementary Materials). It is clear from Figure 8a,b that under room-temperature tensile test conditions, the strength of the AP-ZC sample was significantly higher than that of the AC-ZC sample, whereas the opposite trend was noticed for plasticity. Moreover, the elongations of both alloy samples at room temperature were less than 1%, indicating the occurrence of a pronounced brittle fracture. The UTS and yield strength (YS) of the AP-ZC sample are 327.9 and 310 MPa, respectively, which were 11.8 and 14.6% higher than those of the AC-ZC sample, respectively. The elongation of the AP-ZC sample was determined as 0.63%, which was 12.6% lower than that of the AC-ZC sample. The high-temperature tensile properties of both samples were consistent with those observed at room temperature; however, at high temperatures, both samples exhibited higher elongation and experienced a ductile fracture. The increase in temperature caused a significant decrease in strength due to the gradual weakening of various strengthening mechanisms (softening phenomenon) at high temperatures [44].

The room-temperature and high-temperature tensile properties of the T6-treated AC and AP alloy samples were consistent with those of the ZC-treated samples. The AP-T6 sample had higher strength and lower elongation than AC-T6. Moreover, under room-temperature tensile test conditions, the slopes of the stress–strain curves for the T6-treated samples were lower than those for the ZC-treated sample, and the opposite trend was observed under high-temperature test conditions (Figure 8).

In order to further analyze the deformation behaviors of the AC and AP samples, the alloy strength data (UTS, YS, EL, failure rate (FR), and quality index (QI)) obtained from the engineering stress–strain curves are summarized in Table 3. The failure rate is used to quantify the rate at which the tensile strength of a material decreases with the rise in temperature [6].

$$FR={\displaystyle \frac{{\left(UTS\right)}_{HT}-{\left(UTS\right)}_{RT}}{{\left(UTS\right)}_{RT}}}\times 100\%$$

(2)

where FR is the failure rate, and (UTS)_RT and (UTS)_HT are the room-temperature and high-temperature tensile strengths. The AP-T6 sample exhibited the highest FR value of 78.37%, whereas the AC-ZC sample had the lowest FR value of 74.76%, indicating that the strength of the AP-T6 sample was more sensitive to the temperature change. The FR values of the AC samples were consistently lower than those of the AP samples, regardless of whether the ZC or T6 treatment was applied. Furthermore, for the same sample, the FR values after the ZC treatment were lower than those after the T6 treatment. These results demonstrate that the AC samples exhibited strength stability superior to that of the AP samples under rising temperatures.

The overall performance of a material can be demonstrated by QI [45].

$$QI=YS+50\times EL$$

(3)

It is observable from Table 3 that the QI values under high-temperature test conditions were higher than those under room-temperature test conditions, and the QI values of the T6-treated alloy samples were found to be higher than those of the ZC-treated samples. Furthermore, at room temperature, the QI values of the AP samples were higher than those of the AC samples, with the AP-T6 sample exhibiting the maximum QI value of 368 MPa. However, under high-temperature test conditions, the opposite trend was observed. The QI values of the AC alloy samples exceeded those of the AP samples, and the AC-T6 sample achieved the highest QI value of 516.5 MPa. Hence, the synergistic application of Ca modification and the T6 heat treatment was found to be more effective to simultaneously enhance the strength and elongation of the Al-12Si-4Cu-2Ni-0.5Mg piston alloy to a matching degree at high temperatures.

3.3. Fracture Morphology of AC and AP Samples

To further analyze the fracture mechanisms of the alloys, the tensile samples after fracture were analyzed, as shown in Figure S4 (in the Supplementary Materials). Figure 9 presents the room-temperature tensile fracture surfaces and the corresponding cross-sectional microstructures near the fractures. Numerous cleavage surfaces of Si phases are present on the fracture surfaces of all samples at room temperature. Furthermore, the dimples are almost absent, indicating poor plasticity and characteristic brittle fracture surfaces, which is consistent with the results obtained from the tensile curves. As seen in Figure 9a,b, in the fracture surface of the AC-ZC sample, the cleavage surfaces of Si phases appear as small, block-shaped particles with a more uniform distribution. In contrast, the cleavage surfaces of Si phases in the AP-ZC sample exhibit large, elongated shapes. This phenomenon is more pronounced in the AC-T6 and AP-T6 samples, as shown in Figure 9c,d.

Figure 9(a1–d1) present the cross-sectional microstructures near the room temperature fracture surfaces of the samples. It is observed that micro-voids in the fractured microstructure almost invariably originate at the interfaces between the eutectic Si particles and the Al matrix (indicated by orange arrows). Under sustained tensile stress, these micro-voids grow, coalesce, and become initiation sites for cracks, as shown by the red arrows. The subsequent propagation of these cracks leads to material fracture. Furthermore, Figure 9(b1) reveals that cracks not only propagate along the eutectic Si particles but also cause a significant fracture of the elongated eutectic Si particles themselves. This indicates that the elongated eutectic Si particles are subjected to higher stress during tensile deformation, corresponding to the significantly higher strength of the AP sample compared to the AC sample at room temperature.

As temperature increases, the α-Al matrix undergoes further softening, enabling it to sustain greater damage during tensile deformation and thereby enhancing alloy elongation [46]. Figure 10 presents the high-temperature tensile fracture surfaces and corresponding cross-sectional microstructures near the fractures. Compared to room-temperature fracture surfaces, the high-temperature fracture surfaces exhibit numerous dimples (indicated by red arrows) and tear ridges, characteristic of ductile fracture. Notably, fractured Si particles are present within these dimples, serving as primary initiation sites for microcracks. Figure 10a,c indicate that AC samples display a higher density of relatively smaller, more uniformly distributed dimples. The dimples in the AC-T6 sample are the deepest, signifying the extensive plastic deformation of the Al matrix during loading, which correlates with its highest elongation. Furthermore, Figure 10b,d clearly show that the dimples in both ZC-treated and T6-treated AP samples contain larger fractured Si particles at their bases. Figure 10(a1–d1) display the cross-sectional microstructures near the tensile fractures at 350 °C. Similarly to room-temperature fractures, numerous micro-voids are observed at eutectic Si particles. However, many larger cavities are also distributed throughout the microstructure. The softened α-Al matrix at high temperature facilitates the coalescence and growth of micro-voids into larger cavities under stress.

3.4. Microstructure and Tensile Properties of ATCB Alloy Samples

Studies on AC and AP alloys revealed that Ca modification is a potential method for improving the strength–ductility balance of Al-Si-Cu-Ni-Mg piston alloys at high temperatures. As shown in Table 3, the AP-T6 alloy with the highest tensile strength reached only 80.1 MPa, indicating that the simple modification of the Si phase and heat treatment had limited effectiveness in enhancing the high-temperature strength of the alloy. Previous research demonstrated that the Al-TCB master alloy can effectively refine the grains of the Al-5Cu alloy and modulate the second-phase structure at grain boundaries, significantly improving the high-temperature tensile properties [31]. Based on this, a 10% Al-TCB master alloy was introduced into the Al-12Si-4Cu-2Ni-0.5Mg alloy, and the room-temperature and high-temperature tensile properties of the alloy under the ZC and T6 heat treatments were tested.

Due to the addition of the Al-TCB master alloy, intermetallic compounds and coral-like eutectic Si are distributed in a network pattern along the α-Al grain boundaries, as shown in Figure 11a,b. Between these networks, micron and submicron C-TiB₂ (C-doped TiB₂, yellow arrows) and Al₄C₃ (red arrows) particles, formed in situ from TCB complex particles, are present. These particles are located between intermetallic compounds and eutectic Si phases, which can enhance the cohesion between such phases. The reaction process is shown in below [29]:

Al(l) + TCB complex(s)→Al₄C₃(s) + [Ti] + C-TiB₂(s)

(4)

Furthermore, Figure 11c,d show that after the T6 treatment, the interconnected coral-like eutectic Si transforms into isolated granular particles, and the network structure is disrupted, consistent with the phenomena observed in T6-treated AC and AP alloys.

Figure 12 presents the engineering stress–strain curves of the ATCB and AC-ZC alloy samples, and Figure S5 (in the Supplementary Materials) presents the true stress–strain curve of the alloy samples before the UTS. In addition, the corresponding tensile property data are listed in

Table 4. Room- and high-temperature tensile properties of ATCB and AC-ZC samples.

Temp/°C	Samples	UTS/MPa	YS/MPa	EL/%	FR/%	QI/MPa
25	ATCB-ZC	257.8 ± 5.8	255.4 ± 3.4	0.40 ± 0.01	-	277.8
	ATCB-T6	280.9 ± 6.3	275.8 ± 2.8	0.42 ± 0.02	-	301.9
	AC-ZC	293.2 ± 3.1	270.5 ± 5.1	0.71 ± 0.07	-	306
350	ATCB-ZC	101 ± 1.0	78 ± 2.0	3.8 ± 0.11	60.82	291.0
	ATCB-T6	91.5 ± 0.5	75 ± 1.0	4.75 ± 0.25	67.42	329.0
	AC-ZC	74.0 ± 0.5	59.2 ± 1.1	7.16 ± 0.24	74.76	417.2

, and it is evident that the high-temperature strength of the ATCB alloy samples is significantly improved compared to the AC alloy. The high-temperature tensile strength increases from 74 MPa (AC-ZC) to 101 MPa (ATCB-ZC) with a 36.5% enhancement. Based on the microstructure of the alloy in Figure 11, the micron and submicron C-TiB₂ and Al₄C₃ particles formed in situ from TCB complex particles provide flexible connectivity to the second-phase network structure. During high-temperature deformation, these particles not only hinder dislocation motion but also stabilize the second-phase network structure, thereby enhancing the high-temperature strength of the alloy. Additionally, it is noteworthy that the T6 treatment disrupts the second-phase network structure. Consequently, the improvement in high-temperature tensile strength for the T6-treated ATCB alloy compared to the AC alloy is 17.3%, which is lower than that of the ZC-treated alloy (36.5%).

Both the ATCB-ZC and ATCB-T6 samples exhibit low elongation at room temperature, characteristic of brittle fracture. The room-temperature fracture morphology in Figure 13a,c reveals numerous cleavage planes with almost no dimples. In contrast, the high-temperature fracture surfaces show more dimples, indicative of ductile fracture. Figure 13d further shows that the T6-treated ATCB alloy sample exhibits deeper and more uniformly distributed dimples, corresponding to its higher elongation.

3.5. Strengthening Mechanism

To gain deeper insight into strengthening mechanisms at both room and high temperatures, a more thorough analysis of their tensile deformation behavior is necessary. Generally, load transfer strengthening refers to the phenomenon where a soft matrix transfers the applied load to reinforcing phases, thereby increasing the strength of the material. Evsevleev et al. [47] employed computed tomography to determine the volume fractions of eutectic Si and intermetallic compound phases in an AlSi12CuMgNi alloy and assessed internal damage after in situ compression tests. They identified a load transfer effect between the Si phase and the intermetallic compound phases, with significant stress concentrations existing around the intermetallic compounds.

Under room-temperature testing conditions, the deformation of the α-Al matrix is difficult to attain, and the stress concentrates at the Si phases, generating micro-voids that evolve into microcracks. Subsequently, the cracks propagate along the Si particles and intermetallic compounds, leading to the cleavage fracture of the Si phase and the ultimately failure of the alloy. According to the data in Table 3, the strength of the AP-ZC sample is significantly higher than that of the AC-ZC sample at room temperature. It is evident that in the AC sample, the eutectic Si phase is fibrous and distributed along grain boundaries. In contrast, the eutectic Si in the AP sample appears as elongated plates, and a small amount of blocky primary Si phase is also present. Under load, significant stress concentration occurs at the interfaces between the α-Al matrix and the brittle eutectic Si particles. This explains the substantially higher room-temperature tensile strength of the AP-ZC sample compared to the AC-ZC sample. An examination of Table 3 and Table 4 further reveals that the room-temperature strength of the T6-treated samples is significantly higher than that of the ZC-treated samples. The primary reason for this lies in the fact that Orowan strengthening resulting from the precipitates hinders dislocation motion. Generally, the effectiveness of Orowan strengthening is proportional to the volume fraction of such precipitates [43]. Referring to Figure 7, it is evident that a high density of precipitates is present in the T6-treated sample, whereas only a small amount exists in the ZC-treated sample. Consequently, the room-temperature strength of the T6-treated sample is substantially greater than that of the ZC-treated sample.

However, under high-temperature testing conditions, the precipitation strengthening effect is diminished due to the softening of the α-Al matrix and Ostwald ripening of the precipitates [48], and the softened α-Al matrix exhibits improved deformation compatibility and can absorb fracture energy at high temperatures. Microcracks are less prone to propagation under external load, and load can be effectively transferred to the Si and intermetallic compound phases. Under these conditions, the networked second phases at the grain boundaries play a critical role [6,9].

Comparing Table 3 and Table 4, it can be observed that the high-temperature strength of the AC-T6 sample is higher than that of AC-ZC. The primary reason for this is that the second-phase network in the AC alloy is relatively simple, consisting only of Si-containing phases (eutectic Si and Al₅Cu₂Mg₈Si₆ phase) and Ni-rich phases (Al₃CuNi and Al₇Cu₄Ni phase), and the connectivity between the second phases is weak, resulting in a limited strengthening effect at high temperatures. After the T6 treatment, although the network structure of the alloy is disrupted, the precipitates formed during T6 still exert some strengthening effect on the high-temperature strength of the alloy. However, this strengthening effect is relatively poor, resulting in the overall low strength of the alloy. For the ATCB alloy, the in situ formed C-TiB₂ and Al₄C₃ particles act as bridging agents within the second-phase network, enhancing the high-temperature strengthening effect of the network structure, as shown in Figure 14. Consequently, the high-temperature tensile strength of the ATCB alloy is significantly improved compared to the AC alloy. However, after the T6 treatment, the network structure is damaged, leading to lower high-temperature strength in the ATCB-T6 sample compared to ATCB-ZC. This further demonstrates that the second-phase network structure plays a dominant role in strengthening the alloy at high temperatures.

4. Conclusions

(1)

Eutectic Si particles in the AC sample had a fibrous structure, whereas the AP sample contained elongated eutectic Si particles. The AC sample exhibited higher elongation because of its better plasticity under the effect of fibrous eutectic Si particles; thus, it had a higher QI value at 350 °C. Ca modification is a potential method for enhancing the matching degree of strength and plastic at high temperature for Al-Si-Cu-Ni-Mg piston alloys.

(2)

The micron and submicron C-TiB₂ and Al₄C₃ particles formed by the in situ reaction of TCB particles acted as bridging phases within the second-phase network structure. During high-temperature deformation, these particles improved the stability of the second-phase network structure, causing a significant strength improvement in the alloy. The ultimate tensile strength at 350 °C increased from 74 MPa for the AC-ZC sample to 101 MPa for the ATCB-ZC sample, representing a 36.5% enhancement.

(3)

The T6-treated samples displayed higher FR than the ZC-treated samples due to the disruption of the second-phase network structure and the Ostwald ripening of nanoscale precipitates. The comprehensive analysis revealed that Orowan strengthening was the dominant strengthening mechanism at room temperature, whereas load transfer and network structure strengthening were the key strengthening mechanisms at high temperatures.