Effect of Nickel Addition on Precipitate Microstructures and Strength at an Elevated Temperature in Wrought Al-Si-Cu-Mg Alloys

Abstract

Al-Si-Cu-Mg-(Ni) alloys are used for engine pistons or compressor scrolls due to their excellent heat and wear resistance. Although these components are typically produced using mold-casting or die-casting techniques, forging techniques have recently attracted attention because they allow for a lightweight design. However, wrought Al-Si-Cu-Mg-(Ni) alloys have rarely been investigated, and the effect of nickel on precipitate microstructures and strength remains unclear. Here, we investigated the precipitate microstructure and strength at 473 K of wrought Al-Si-Cu-Mg alloys with different nickel concentrations. The 0.2% proof stress and tensile strength decreased with increasing nickel concentration. The TEM observations showed that the decline was attributed to the decrease in the θ′ phase. The increase in the Al₇Cu₄Ni and Al-Fe-Cu-Si-Ni phases suggested that the nickel decreased the θ′ phase by consuming the copper in the matrix. In addition, the HRTEM observation revealed the difference in the orientation relationships of the Q phase between with and without nickel alloys. The Q phase observed in the 2%Ni alloy had a smaller lattice misfit with the matrix than that observed in the 0%Ni alloy, indicating excellent thermal stability.

1. Introduction

Al-Si-Cu-Mg-(Ni) alloys are suitable for sliding parts used at elevated temperatures, such as engine pistons or compressor scrolls, due to their excellent heat and wear resistance [1, 2]. Typically, these parts are produced using mold-casting or die-casting techniques. However, castings have been limited in further weight reduction because these casting methods cause internal defects, such as pores and coarse constituent particles [3–5]. Therefore, lightweight design requires a technique that can produce highly reliable parts without internal defects.

Forging techniques have recently attracted attention. Forgings have no internal defects, enabling lightweight design [5, 6]. In addition, there are two more reasons that forging techniques have attracted attention. The first is that new casting methods have been developed for producing small-diameter billets for forgings. Forgings used to require extrusion bars. Thus, the manufacturing cost of forgings had been higher than that of castings. However, using small-diameter billets as forging materials has reduced the manufacturing cost of forgings [5, 7–11]. The second is that the reduction of billet diameter has allowed the improvement of the strength of forgings. The extrusion bars are produced from large-diameter billets. The large-diameter billets have slow cooling rates and the possibility to form coarse constituent particles at the center of billets. Thus, amounts of alloying elements have been limited in the large-diameter billets. However, the improvement of the cooling rate by reducing the billet diameter has allowed for the increase of alloying elements, such as nickel, copper, and magnesium, enabling the improvement of the strength of forgings [12–16].

In our previous work, we fabricated the wrought Al-Si-Mg-Ni alloy billets with different copper concentrations up to 6 mass%, and we investigated their microstructures and mechanical properties at elevated temperatures after forging. It revealed that the increase in the volume fraction of the Al₇Cu₄Ni and Al-Fe-Cu-Si-Ni phases increased the mechanical properties at 623 K [15]. However, our results imply that the nickel reduces the strength below 473 K by consuming the copper, which contributes to the precipitate hardening, as reported in cast Al-Si-Cu-Mg-Ni alloys [17]. In contrast, Medrano-Prieto et al. reported that nickel increased the peak hardness and made the hardness decrease slower during the overaging [18]. Thus, the effect of nickel on precipitate microstructures and strength at elevated temperatures remains controversial. An advantage of forged products is that they contain no gas defects, allowing for solution heat treatment at high temperatures and maximizing the effects of precipitate hardening. Therefore, improving the strength below 473 K requires an understanding of the effect of nickel on precipitate microstructures.

Here, we report the effect of nickel addition on the precipitate microstructures and the strength at 473 K of the wrought Al-Si-Cu-Mg alloys.

2. Experimental Procedure

Three Al-Si-Cu-Mg alloy billets (φ 95) with different nickel concentrations were fabricated. Table 1 shows their chemical compositions analyzed by an optical emission spectrometer (PDA-7000, Shimadzu Corp.). These alloys are referred to as the 0%Ni alloy, 1%Ni alloy, and 2%Ni alloy. The billets were homogenized at 753 K for 28.8 ks and then cooled to room temperature. Materials for forging with a diameter of 50 mm and a height of 80 mm were cut from the center of the billets and forged to 12 mm thickness at 703 K.

Table 1 Chemical compositions of the experimental alloys (mass%).

Forged materials were solution-treated at 773 K for 3.6 ks, quenched into hot water at 333 K for 30 s, and aged at 443 K for 24 ks in an oil bath. Afterward, these materials were exposed at 473 K for 360 ks and machined into the specimen for the tensile test shown in Fig. 1.

Fig. 1

Specimen for the tensile test at the elevated temperature.

The tensile tests were performed at 473 K using the tensile tester with an electric furnace (AG-100kNXplus, Shimadzu Corp.). The tests were conducted at an initial strain rate of 5 × 10⁻⁵ s⁻¹ until 2% strain and at a strain rate of 1.67 × 10⁻³ s⁻¹ after 2% strain following the ISO 6892-2:2011: Method of test at elevated temperature.

The microstructural observations were performed using an SEM equipped with an EDS (JSM-6010PLUS/LA, JEOL Ltd.) and a TEM (JEM-2100, Ltd.). The specimens for SEM observation were prepared by grinding with SiC papers, polishing with colloidal silica, and etching with 0.2% HF. The thin foils for TEM observation were prepared by an electrolytic polishing technique in a solution of 33% HNO₃ and 67% methanol at 253 K. A Low-energy ion milling (TEM Mill Model1051, E.A. Fischione Instruments. Inc.) was also conducted for the TEM specimens to minimize the oxide layer. The ion milling was performed for 600 s at an acceleration voltage of 0.5 kV and a milling angle of ±6 degrees. An image analysis software (WinROOF 2018, Mitani Corp.) was used to quantify the observed structures.

3. Results

3.1 Mechanical properties at 473 K

Figure 2 shows the mechanical properties of the alloys tested at 473 K. The 0.2% proof stress and tensile strength decreased with the increase in nickel concentrations. Compared with the 0%Ni alloy, the 0.2% proof stress and tensile strength of the 2%Ni alloy decreased by 4% and 5%, respectively. These results reveal that a nickel addition to wrought Al-Si-Cu-Mg alloys reduces the strength at 473 K. In contrast, the elongations showed almost 10% independent of the nickel concentration.

Fig. 2

Mechanical properties of the 0%Ni, 1%Ni, and 2%Ni alloy tested at 473 K.

3.2 Microstructures after exposure at 473 K for 360 ks

3.2.1 Constituent particles and volume fractions

Figure 3 shows the SEM images of the 0%Ni alloy, 1%Ni alloy, and 2%Ni alloy after exposure at 473 K for 360 ks and an example of EDS mapping analysis obtained from the 2%Ni alloy. Eutectic Si, Al-Cu, Al-Cu-Fe, and Al-Cu-Mg-Si phases were determined in the 0%Ni alloy. In the 1%Ni ally and 2%Ni alloy, the Al-Cu-Ni and Al-Si-Fe-Ni-Cu phases were determined instead of the Al-Cu and Al-Cu-Fe phases. According to the previous studies [2, 15], these phases were identified as Al₂Cu, Al₇Cu₂Fe, Al₄Cu₂Mg₈Si₇, and Al₇Cu₄Ni. The constituent particles, which include copper or nickel, showed a bright contrast in SEM images, and their volume fractions tended to increase with the increase in nickel concentration. In the cast Al-Si-Cu-Mg-Ni alloys, it is considered that a rigid 3D network consisting of eutectic Si and these constituent particles contribute to the strength [19–21]. However, no network-like structure was observed in the forged materials, and their particles were distributed in the matrix. In addition, no dendrite-like α-Al phase was observed, and the α-Al phase exhibited recrystallized grain structures. These observation results show that the microstructures of forgings and castings are different and that the strengthening mechanism caused by the rigid 3D network structure does not work after forgings.

Fig. 3

SEM images of (a) 0%Ni alloy, (b) 1%Ni alloy, and (c) 2%Ni alloy, and (d) EDS elemental maps of 2%Ni alloy. (online color)

Figure 4 shows the volume fractions of the eutectic Si and the constituent particles in the 0%Ni alloy, 1%Ni alloy, and 2%Ni alloy obtained from image analyses of SEM images. The volume fractions were classified into the eutectic Si, Al₄Cu₂Mg₈Si₇, and the others, according to the difference of their contrast in SEM images. The average area fractions obtained from the image analyses were substituted for the volume fractions because the following relationship is established [22]:

  

\begin{equation} \overline{A_{A}} = V_{f}, \end{equation} (1)

where $\overline{A_{A}}$ is the average area fraction obtained from several image analyses. The increase in nickel concentration increased the total volume fraction by forming the Al₇Cu₄Ni and Al-Si-Fe-Ni-Cu phases. This result implies that the increase in nickel concentration decreases the copper concentration in the matrix by forming the Al₇Cu₄Ni and Al-Si-Fe-Ni-Cu phases.

Fig. 4

Volume fractions of the eutectic Si and constituent particles of 0%Ni, 1%Ni, and 2%Ni alloy.

3.2.2 Precipitate microstructures

Figure 5 shows TEM bright field images and SAED patterns observed from [001]_Al directions. Two types of precipitates were observed in every alloy. The first was θ′ phases ($\text{I}\bar{4}\text{m}2$, a = b = 0.404 nm, c = 0.580 nm), indicated by the white arrows [23]. The θ′ phases were identified by analyzing the SAED patterns. The length of the θ′ phases was 50 to 100 nm and elongated to the [100]_Al or [010]_Al directions. The second was particle-like precipitates with 10 nm, shown by the black arrows. These precipitates were not identified because no clear diffraction patterns from the particle-like precipitates were observed in the SAED patterns. HRTEM observations were carried out for the 0%Ni alloy and 2%Ni alloy to identify the particle-like precipitates.

Fig. 5

TEM bright field images and their SAED patterns: (a), (b) 0%Ni alloy, (c), (d) 1%Ni alloy, and (e), (f) 2%Ni alloy.

Figure 6 shows typical HRTEM images of the particle-like precipitates and their FFT patterns observed in the 0%Ni alloy. The particle-like precipitates were distinguished into two types, as shown in Fig. 6(a) and (c). The first was the rod-shaped Q′ phase shown in Fig. 6(a). Its cross-section showed periodic hexagonal bright dots with a spacing of 1.04 nm, as reported previously [24–27]. The FFT pattern displayed a hexagonal spot corresponding to $\{ \bar{1}100\}_{\text{Q}' }$ and a strong hexagonal spot corresponding to the interplanar spacing of 1.964 Å marked by the orange circles. This strong hexagonal spot corresponded to the interplanar spacing of $\{ \bar{5}140\}_{\text{Q}' }$, that is, the periodicity of the atomic positions of the Q′ phase marked by the orange circles in the HRTEM image. These results show that the positions of Al, Cu, Mg, and Si atoms are located on the hexagonal lattice with a = 0.226 nm. The upper right image in the HRTEM image shows the IFFT image of the Q′ phase obtained from the spots marked by yellow and orange circles in the FFT pattern. The IFFT image indicates that the Q′ phase consists of the blue and red triangles. These two properties agree with the modified structure of the Q phase suggested by Cayron et al., that is, the hexagonal lattice with a = 0.226 nm and triangles correspond to the qh-lattice (a = 0.226 nm) and the QP sub-unit clusters, respectively [28, 29]. From now on, the Q′ phase is referred to as the Q phase following their reports. According to their reports, two pairs of the orientation relationships (ORs) between the Q phase, the qh-lattice, and the matrix are deduced as follows:

  

\begin{align*} &[\bar{1}\bar{1}20]_{\text{Q}}\ \angle\ [010]_{\text{Al}} = 10.9{{}^{\circ}},\ [\bar{1}2\bar{1}0]_{\text{Q}}\ \angle\ [100]_{\text{Al}} = 19.1{{}^{\circ}},\\ &\text{and}\ (\bar{1}100)_{\text{qh}} \mathrel{/\!/} (100)_{\text{Al}}\ (\text{OR}1) \end{align*}

  

\begin{align*} &[\bar{1}\bar{1}20]_{\text{Q}}\ \angle\ [010]_{\text{Al}} = 34.1{{}^{\circ}},\ [\bar{1}2\bar{1}0]_{\text{Q}}\ \angle\ [100]_{\text{Al}} = 64.1{{}^{\circ}},\\ & \text{and}\ (\bar{1}100)_{\text{qh}}\ \angle\ (100)_{\text{Al}} = 45{{}^{\circ}}\ (\text{OR}2) \end{align*}

However, the ORs between the Q phase, the qh-lattice, and the matrix differed, as follows:

  

\begin{align*} &[\bar{1}\bar{1}20]_{\text{Q}}\ \angle\ [010]_{\text{Al}} = 1.4{{}^{\circ}},\ [\bar{1}2\bar{1}0]_{\text{Q}}\ \angle\ [100]_{\text{Al}} = 28.6{{}^{\circ}},\\ & \text{and}\ (\bar{1}100)_{\text{qh}}\ \angle\ (100)_{\text{Al}} = 9.5{{}^{\circ}}\ (\text{OR}3) \end{align*}

Fig. 6

Crystal lattice images of the precipitates and their FFT patterns observed in 0%Ni alloy. (online color)

The second was the C phase-like precipitate shown in Fig. 6(c). Its cross-section showed a monoclinic unit cell with parameters of a = 1.03 nm, b = 0.81 nm, and γ = 101°, and was elongated to the [100]_Al direction as reported by Marioara et al. [30]. In addition, its interface with the matrix was coherent. Their report indicated that the C phase had a periodic hexagonal atomic arrangement named the Si network (a = b = 4.05 Å). However, no atomic arrangement corresponding to the Si network was observed in the HRTEM image (Fig. 6(c)), and the atomic arrangement was the same as the matrix. The FFT pattern of the HRTEM image displayed the spots corresponding to (100)_C and (010)_C. In addition, the FFT pattern displayed strong spots pointed by the yellow arrows. These spots corresponded to (420)_C, (020)_C, and $(\bar{1}20)_{\text{C}}$, respectively, but no hexagonal spot corresponding to the Si network was also identified in the FFT pattern. Therefore, this precipitate differed from the C phase reported by Marioara et al. [30]. This phase is referred to as the C phase-like precipitate.

The Q phase and the C phase-like precipitate were observed in addition to the θ′ phase in the 0%Ni alloy, but the Q phase was more frequent than the C phase-like precipitate. These observations indicate that the Q phase and θ′ phase contribute to the strength at 473 K.

In the 2%Ni alloy, both Q phases and the C phase-like precipitates were also observed, and the Q phase accounted for a significant fraction as well as the 0%Ni alloy. However, the ORs with the matrix were different from the OR1, OR2, and those observed in the 0%Ni alloy. Figure 7 shows an HRTEM image of the Q phase and its FFT pattern. The Q phase cross-section showed periodic hexagonal bright dots with a spacing of 1.04 nm, and the lower right enlarged image showed the QP sub-unit clusters consisting of blue and red triangles. In addition, the FFT pattern showed the strong spots corresponding to the $\{ \bar{1}100\}_{\text{qh}}$. These properties also agree with the modified structure of the Q phase suggested by Cayron et al. [28, 29]. In contrast, the ORs with the matrix were different, as follows:

  

\begin{align*} &[\bar{1}\bar{1}20]_{\text{Q}}\ \angle\ [010]_{\text{Al}} = 17.4{{}^{\circ}},\ [\bar{1}2\bar{1}0]_{\text{Q}}\ \angle\ [100]_{\text{Al}} = 12.6{{}^{\circ}},\\ & \text{and}\ (\bar{1}100)_{\text{qh}}\ \angle\ (100)_{\text{Al}} = 6.5{{}^{\circ}}\ (\text{OR}4) \end{align*}

Fig. 7

Crystal lattice image of the Q phase and its FFT pattern observed in the 2%Ni alloy. (online color)

4. Discussion

4.1 Effect of nickel on the mechanical properties at 473 K

Although the volume fractions of the constituent particles increased with the nickel concentration, the 0.2% proof stress and tensile strength at 473 K decreased. These results indicate that the volume fraction of the constituent particles does not affect the mechanical properties at 473 K. Figure 3 and Fig. 5 showed that the fine precipitates were in the matrix, and their interparticle spacings were shorter than those of the eutectic Si and constituent particles. In addition, the eutectic Si and constituent particles were on grain boundaries, not in the matrix. Therefore, the obstacles of dislocation slips should be fine precipitates in the matrix, and the θ′ and Q phases are likely to contribute to the strength at 473 K.

Figure 8 shows the number densities of θ′ and Q phase of the 0%Ni alloy, 1%Ni alloy, and 2%Ni alloy. The number density of the θ′ phase decreased with the increase in nickel concentration, whereas that of the Q phase did not change. These results indicate that the decrease in the mechanical properties at 473 K with increasing nickel concentration is caused by the decline in the number density of the θ′ phase. Additionally, the decline can be explained by the decrease in copper concentration in the matrix by forming the Al₇Cu₄Ni and Al-Fe-Cu-Si-Ni phases, as demonstrated in Fig. 4. However, although the number density of the θ′ phase in the 2%Ni alloy decreased by 25% compared to the 0%Ni alloy, the strength reduction in the 2%Ni alloy was limited to only 4 to 5%.

Fig. 8

Number density of the θ′ phase and the Q phase.

4.2 Orientation relationships and misfits between the Q phase, the qh-lattice, and the matrix

Figure 9 shows the schematic illustrations of the ORs between the Q phase, the qh-lattice, and the matrix reported by Cayron (OR1 and OR2) and observed in the 0%Ni alloy and 2%Ni alloy (OR3 and OR4). The ORs with the matrix are determined by rotating the $(\bar{1}100)_{\text{qh}}$ of qh-lattice with respect to the origin O. For the OR1 and OR2, the rotation angles were 0 and 45 degrees clockwise, respectively. In contrast, for the OR3 and OR4, the rotation angles were 9.5 degrees clockwise and 6.5 degrees counterclockwise, respectively. These comparisons highlight that the ORs between the Q phases and the matrix observed in the 0%Ni alloy and 2%Ni alloy differ from each other and from those reported previously. The reason for their differences is unclear, but the differences in the ORs with the matrix suggest that the interface structures with the matrix are different; that is, the thermal stabilities are different. Here, we discuss the thermal stabilities of the Q phases observed in the 0%Ni alloy and 2%Ni alloy by evaluating the lattice misfits between the atomic rows of the Q phase and the matrix.

Fig. 9

Schematic illustrations of the orientation relationships between the Q phase and the matrix: (a) OR1, (b) OR2, (c) OR3, and (d) OR4. (online color)

Figure 10 shows the interface structures of the Q phase with the matrix observed in the 0%Ni alloy and the 2%Ni alloy (Fig. 6(a) and Fig. 7(a)) and their schematic illustrations. The schematic illustrations shown in Fig. 10(b) and (c) were represented by rotating the $(\bar{1}100)_{\text{qh}}$ of qh-lattice 9.5 degrees clockwise and 6.5 degrees counterclockwise with respect to the origin O, respectively. As shown in Fig. 10(a), the Q phase observed in the 0%Ni alloy had a transition layer with a disordered crystal structure around the interface with the matrix. Figure 10(c) shows that the atomic rows qh-lattice are incoherent with the (100)_Al. The lattice misfit of the Q phase with the matrix calculated from the Q lattice and 4d_200Al was 9.90%. In contrast, the atomic rows of the Q phase observed in the 2%Ni alloy were fully coherent with the (100)_Al, as shown in Fig. 10(b) and (d). The lattice misfit with (100)_Al was calculated to be 0.26%, and it was smaller than that observed in the 0%Ni alloy. The smaller the misfit between precipitates and matrix, the lower the interfacial energy, resulting in a slower growth rate of the precipitates [31, 32]. Therefore, these results indicate that the Q phase with OR4, which has a smaller misfit, coarsens more slowly; that is, more stable at elevated temperatures compared to the Q phase with OR3. In contrast, the smaller the misfit between precipitates and matrix, the greater the elastic strain energy and the higher the stress required to shear the precipitates [33]. This suggests that the Q phase with OR4, which has a smaller misfit, could be the reason why the strength of 2%Ni alloy was not reduced to an extreme degree. However, to what extent the orientation difference contributes to the strength is currently unclear; further detailed research is needed.

Fig. 10

Interface structures between the Q phase and the matrix and their schematic illustrations: (a), (c) the Q phase observed in 0%Ni alloy and (b), (d) the Q phase observed in 2%Ni alloy. (online color)

5. Conclusion

The mechanical properties at 473 K and the precipitate microstructures were investigated for the wrought Al-Si-Cu-Mg alloys with different Ni concentrations. The findings are summarized as follows:

1. (1)    The 0.2% proof stress and tensile strength of the 0%Ni alloy were the highest, and their strength decreased with the increase in nickel concentration.
2. (2)    The increase in nickel concentration increased the total volume fraction of the eutectic Si and constituent particles due to the forming of Al₇Cu₄Ni and Al-Si-Fe-Ni-Cu phases.
3. (3)    The θ′ phase was observed in each alloy. In the 0%Ni alloy and 2%Ni alloy, the Q phase and C phase-like precipitate were observed. The θ′ phase and Q phase were more frequent than the C phase-like precipitate.
4. (4)    The number density of the θ′ phase decreased with the increase in nickel concentration, whereas that of the Q phase did not change, suggesting that the decline in the mechanical properties at 473 K was attributed to the decrease in the θ′ phase. The decrease in the number density of the θ′ phase was explained by the decrease in copper concentration in the matrix by forming the Al₇Cu₄Ni and Al-Fe-Cu-Si-Ni phases.
5. (5)    The Q phases observed in the 0%Ni alloy and 2%Ni alloy agreed with the modified structure of the Q phase suggested by Cayron et al. However, their ORs with the matrix were different from each other and from those reported by them. The Q phase observed in the 2%Ni alloy had a smaller lattice misfit with the matrix than that observed in the 0%Ni alloy, indicating excellent thermal stability.

Acknowledgments

This work was supported by the Division of Instrumental Analysis at the University of Toyama and the Toyama Industrial Technology Research & Development Center.

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