Enhanced mechanical and damping properties of Cu-Modified Al-Ni eutectic alloys: Solid-solution and precipitation hardening effects

1. Introduction

Next-generation materials for various applications must exhibit excellent structural properties (strength, density, ductility, thermal stability, etc) together with diverse functional properties such as electromagnetic shielding, damping, conductivity, magnetic properties, self-healing, and so on [[1], [2], [3], [4], [5], [6]]. Especially, Al-based alloys and composites have attracted the most attention for lightweight characteristics, which enhance fuel efficiency, as well as their controllable mechanical properties and various functional properties, including high damping properties, self-healing behavior, and so on [[7], [8], [9], [10], [11], [12], [13], [14], [15], [16], [17], [18], [19], [20], [21], [22], [23], [24], [25]]. Furthermore, certain Al alloys and composites even demonstrate good thermal stability [26,27], enabling them to maintain their strength at elevated temperatures and thereby expanding the potential applications of the Al-based materials [23,28].

For the high-temperature application, there are two main factors; first, the reinforcement (intermetallic compounds, precipitates, etc) should possess high thermal stability; second, the additional elements should exhibit low diffusivity within the Al matrix [28]. For instance, element such as Fe, Mg, Si, and Zn have a relatively higher diffusion coefficient in the range of 10⁻¹²∼10⁻¹⁴ m²/s, whereas elements like Ni, Ti, Cr, Mn, and Sc have lower diffusion coefficients in the range of 10⁻¹⁵∼10⁻²⁰ m²/s [29]. Due to the lower diffusivity of the Ni in the face-centered cubic (fcc) Al matrix, which resists coarsening, the Al-Al₃Ni eutectic alloy exhibits excellent thermal properties up to 773 K [23,28]. However, the binary Al-Al₃Ni eutectic alloy has a quite low yield strength at room temperature, and the further strengthening effects induced by solid-solution hardening are obscure due to the limited solubility of Ni in Al [18], rendering it unsuitable for structural applications. To enhance the structural properties of this alloy system, the addition of third elements/solutes that promote additional strengthening is necessary. It has been reported that the addition of transition metals like Sc and Zr can lead to precipitation hardening through the formation of coherent intermetallic phases of L1₂ in the fcc Al matrix [23,25,28]. While the addition of Sc or Zr to Al-Al₃Ni eutectic alloy enhances mechanical strength via precipitation strengthening by coherent Al₃Zr or Al₃Sc, the volume fraction of the precipitates is low due to their limited solubility in fcc Al phase and solid-solution hardening effects are not significant [25]. Therefore, further strengthening mechanisms should be activated for the further improvement of the mechanical strength of the Al-Al₃Ni alloy/composite.

In view of the above discussion, Cu can be a promising candidate for facilitating further strengthening by higher solubility in the fcc Al matrix inducing the solid-solution hardening as well as precipitation [30]. Cu is the well-known element for enhancing the strength of Al alloys and some Al-based eutectic alloys containing Cu have been reported that they exhibit very high strength of GPa level and reasonable compressive strains [10,11,17,31].

The study of damping behavior in Al alloys is essential due to their widespread use in structural applications, particularly in aerospace, automotive, and construction industries. These alloys are favored for their lightweight properties, high mechanical strength, and corrosion resistance. In structural applications, components are often exposed to dynamic loads, vibrations, and impacts, making effective vibration control crucial. The damping characteristics of Al alloys play a vital role in dissipating vibrational energy, which is key to enhancing the stability and reliability of structures. Understanding the damping behavior of Al alloys is therefore critical for improving the overall performance and durability of such materials in dynamic environments [[32], [33], [34], [35]].

We reported in previous studies on the change of characteristics and damping behavior in a compressive condition according to the addition of Cu to AlNi process alloys [14]. In our previous study, the possibility of improving the characteristics of AlNiCu alloys was confirmed through the analysis of compressive strength and compressive strain according to the addition of Cu and the study of damping behavior in a compressed environment. However, in order to be applied to the actual industrial environment, the characteristics in the tensile environment are important, and the frequency depletion in the damping characteristics is selected as an important factor.

Thus, the present study focuses on the effect of the Cu addition on the microstructural evolutions, mechanical properties, and hardening effects of Al-Al₃Ni alloy. Additionally, in order to explore new application fields, the damping properties which are important for automotive applications, were investigated under the different temperatures and frequencies.

2. Materials and methods

The (Al₉₇Ni₃)_100-xCu_x with x = 0, 5, and 10 at.% (AN973(Al₉₇Ni₃), ANC5(Al_92.15Ni_2.85Cu₅), ANC10(Al_87.3Ni_2.7Cu₁₀) for short) were fabricated by arc-melting with high-purity of elements at least 99.998 % in a titanium-gettered Ar atmosphere. The ingots were re-melted 5 times. Rod-type samples with a diameter of 3 mm and a length of 150 mm were cast into the cylindrical water-cooled Cu mold by suction pressure (Edmund Bühler GmbH). The phase analysis was performed by X-ray diffraction with Co K_α radiation (λ = 1.79026 Å) with a Fe filter for Co K_β filter employing a Bruker D2 Phaser. The cross-sectional microstructure of the as-cast alloys was investigated by field emission scanning electron microscope (FE-SEM, Zeiss Leo 1525) with an energy-dispersive X-ray spectroscopy system (EDX, XFlash 6/60, Bruker) and the precipitation behavior was investigated by in-situ heating transmission electron microscopy (TEM, JEOL 2100F). Thin films for TEM analysis were prepared by ion milling. The tensile properties of the as-cast alloys were examined by a universal testing machine (UTM, Zwick/Roell, Z100) at a strain rate of 1 x 10⁻⁴s⁻¹. For the tensile test, the round cross-section specimens were machined into a gauge section of 6 mm in length and 2 mm in diameter, and the surface was finely polished by 2500 grit SiC paper. The visualization of the deformation behavior and the strain correction was conducted with a video extensometer system. In order to investigate the precipitation strengthening behavior, the aging process was carried out at 473 K for 30 min, 1h, 3h, 5h, 7h, and 10h in the vacuum chamber. The precipitation strengthening was investigated by the Vickers hardness test using 0.98 N, and 10 tests were performed under each aging condition. The test specimens had a diameter of 3 mm and a length of 5 mm. The Vickers hardness measurements were conducted at the radial midpoint of each specimen. The dynamic mechanical analysis was performed using a Discovery Hybrid Rheometer HR3 (TA Instruments) by the three points bending module in an air atmosphere. The storage modulus, E′, and the loss modulus, E″, were measured at a heating rate of 5 K/min from 300 to 845 K. In order to investigate the frequency dependence damping behavior, the frequency ranged between 0.5 Hz and 15 Hz with an applied load of 5 N.

3. Results

Fig. 1 shows the XRD patterns and backscattered SEM images of as-cast AN973, ANC5, and ANC10. As depicted in Fig. 1 (a), the AN973 alloy displays the strong diffraction patterns of the fcc Al phase along with orthorhombic Al₃Ni intermetallic compounds. The ANC5 exhibits diffraction peaks corresponding to rhombohedral Al₇Cu₄Ni and tetragonal Al₂Cu intermetallic compounds in addition to the diffraction peaks observed in AN973 alloys. With the further addition of Cu up to 10 at.%, the diffraction peaks of Al₇Cu₄Ni and Al₂Cu intermetallic compound become more pronounced, whereas the peaks of Al₃Ni disappear. In addition, the diffraction peaks corresponding to hexagonal Al₃(CuNi)₂ intermetallic compounds were detected. The calculated lattice parameters of the constituent phases (See Supplementary Table S1) suggest that the fcc Al phases are getting smaller with increasing Cu contents, and Al₃Ni intermetallic compound slightly stretches along the c-axis. The inset of Fig. 1 (a) exhibits an enlargement of a comparison of the (111) peaks fitted with the Pseudo-Voigt line shape function, indicating that the peak shifts toward higher angles side with Cu content increase, suggesting a smaller lattice parameter of Al phase.

The SEM micrograph of AN973 alloy clearly shows the typical rod-type eutectic structure, which is in good agreement with the literature [16,18,[24], [25], [26],28,36,37]. The bright contrast nanofibers are the Al₃Ni intermetallic compound, with diameters of 30∼40 nm, and the dark contrast is the Al matrix. With 5 at.% of Cu added (ANC5), the microstructure changes to a dendrite-nanostructured composite, including three different eutectic structures together with two different dendritic phases. From the EDX and contrast-based analysis, the dark contrast phases were identified as the α-Al solid-solution (Al_SS) phase, while the bright contrast phases were the Al₃Ni intermetallic compounds. In addition, there are three different eutectic structures: one is a coarse eutectic colony, another is a fine cellular eutectic, and the other is the ultrafine eutectic matrix. From the EDX results, the atomic ratio of those areas was detected as (Al: 92.2, Ni: 2.9, Cu: 4.9), (Al: 79.2, Ni: 2.3, Cu: 18.5), and (Al: 80.3, Ni: 1.1, Cu: 18.7) with an error range below 3 at.%. This suggests that the coarse eutectic colony is α-Al + Al₃Ni, the fine cellular eutectic is α-Al + Al₇Cu₄Ni, and ultrafine eutectic matrix is α-Al + Al₂Cu.

The ANC10 alloy displays a similar microstructure to ANC5 but with different volume fractions of each area, e.g., the coarse eutectic colony disappears, the dark dendrite phase decreases, and the bright contrast dendrite increases. EDX results identified the bright contrast phase as the Al₃(NiCu)₂ phase with 60.5 at.% of Al, 18.7 at.% of Ni, and 20.8 at.% of Cu. The dark contrast phase was verified as α-Al with a higher amount of Cu contents than that confirmed from ANC5 alloy, i.e., 88.6 at.% Al, 1.3 at.% Ni, and 10.1 at.% Cu, indicating supersaturated Cu in Al primary phases, which can lead to further hardening with the aging process. Moreover, the composite structure of fine cellular eutectic and ultrafine eutectic areas was also analyzed with the EDX area scan, indicating that the eutectic structures observed in ANC5 and ANC10 alloy consist of the same phases such as the cellular eutectic with α-Al + Al₃(NiCu)₂ and the eutectic matrix with α-Al + Al₂Cu. The detailed EDX results and high magnification SEM images can be found in Supplementary Fig. S1 and Fig. S2.

The typical engineering tensile stress-strain curves measured at room temperature for as-cast AN973, ANC5, and ANC10 alloys are shown in Fig. 2. The tensile properties of these alloys are given in Table 1. The results of the tensile test indicated that although the tensile elongation decreased from 14 % to 2 % with increasing Cu content, the tensile strength increased from 172 to 420 MPa. The results indicated that the addition of Cu to the AN973 alloy led to significant increases in the tensile strength with various strengthening mechanisms such as solid-solution hardening, second phase hardening, and so on. The fracture surfaces after the tensile test are shown in Fig. 3. Fig. 3(a) and (b) exhibit the fracture surface obtained from the AN973 alloy. As shown in Fig. 3(a) and (b), the highly dense and small dimples were observed, indicating ductile deformation behavior. The high magnification image of the fractured AN973 alloy as shown in the inset of Fig. 3(b) reveals the fractured Al₃Ni nanofibers as marked by red arrows. Small cracks were only observed at the fibers. In addition, necking was discovered, as shown in Supplementary Fig. S3. With a Cu content increase to 5 at.% (ANC5), the fraction of the dimple patterns is decreased and the mirror-like cleavage fracture patterns observed as marked by the yellow arrows as shown in Fig. 3(c) and (d) corresponding to the Al₃Ni intermetallic compounds, indicating that the Al3Ni intermetallic compounds are intrinsically brittle. Moreover, the mixture patterns composed of dimples and cleavage were presented at the eutectic matrix area as marked by red-dot circles. With the further addition of Cu to 10 at.% (ANC10), this tendency continues; the dimple patterns further decreased, while the alloy exhibited a larger area of cleavage planes compared to the ANC5 alloy. In addition, some cracks are found on the cleavage patterns as marked by blue arrows, suggesting that the dendritic intermetallic compound phases deformed in a brittle manner.

Fig. 4 illustrates TEM results obtained from ANC5 and ANC10 together with in-situ heating for ANC10 alloys in order to verify the possibility of further strengthening induced by supersaturated Cu in the Al phases as mentioned above. Fig. 4(a)and(b) show the STEM images with SAED patterns obtained from ANC5. They clearly display the Al-Al₃Ni eutectic colony, Al-Al₂Cu eutectic matrix, and Al-Al₇Cu₄Ni cellular eutectic with soft/hard dendrite phases composed of dark contrast Al solid-solution and bright contrasted intermetallic compounds. As shown in Fig. 4(c)and(d), ANC10 exhibits the cellular eutectic consisting of Al-Al₇Cu₄Ni and Al-Al₂Cu ultrafine eutectic matrix as well as micron dendrite. Overall, the microstructure is almost similar to that of ANC5 except for the difference in volume fraction, but one difference is that the phases corresponding to Al₃Ni₂ are found as shown in the SAED pattern in Fig. 4(e). These findings are consistent with the results observed in Fig. 1 and Supplementary Fig. S1.

Fig. 4(f)–(h) shows STEM images for ANC10 before and after heating obtained from in-situ TEM. The in-situ STEM images were obtained for ANC10 maintained at 553 K for 10 min. In the intermetallic compound, there was no change during heat treatment, while the precipitation behavior was observed in the micron-scale α-Al phases. As shown in Fig. 4(h), the precipitates generated by the heat treatment appeared as Al₂Cu, which is a phenomenon that generally occurs frequently in the supersaturated α-Al phase and is known to contribute to an additional increase in the strength of the Al-based alloys [1,10,14,20].

Fig. 5 shows the results of measuring the Vickers hardness of specimens heat-treated at 473 K for 30 min to 10 h for AN973, ANC5, and ANC10 to confirm the effect of heat-treatment-induced precipitation on the strength. Fig. 5(a) displays the change in hardness according to the heat treatment time based on the normalized hardness values before and after the heat treatment of each alloy. As shown in Fig. 5(a), it can be seen that the hardness of AN973 slightly increased at the beginning of heat treatment and then continuously decreased as the aging time increased. On the other hand, for the Cu-added ANC5 and ANC10, it was confirmed that the hardness continuously increased with me. Fig. 5(b) shows the measured hardness values for each alloy. In the case of AN973, it can be confirmed that the error range of hardness obtained from each point is wider than others. This variability can be attributed to the eutectic structure of Al-Al₃Ni, containing the Al₃Ni intermetallic compound in the form of a long rod (or wire type), where the measured hardness is affected by the orientation of the microstructure. For that reason, it can be judged that the increase in hardness at the initial stage of aging shown in Fig. 5(a) has no scientific significance. Moreover, it was confirmed that ANC5 and ANC10 showed an increase in hardness of about 10 % through heat-treatment as shown in Fig. 5(b).

The frequency dependence of the loss tangent (damping capacity), Tan δ, for the present alloys is presented in Fig. 6 and Table 2, over a temperature range from 300 to 845 K and an applied frequency range from 0.1 to 15 Hz. From our previous work [14] and phase diagram, the melting point of AN973 is about 910 K, but the melting point gradually decreases as the Cu content increases. Therefore, as the experimental temperature approaches the melting point, the uncertainty of the DMA results increases, so we set the region of interest to the maximum up to 750 K, which is the vicinity where the melting point of ANC10 begins to be affected as shown in Supplementary Fig. S4, and compared the loss tangent values. In general, in complex notation, the damping capacity, Tan δ, is also linked to the ratio between the storage modulus and the loss modulus (Loss modulus divided by the storage modulus). These two factors are related to the stored energy and the dissipated energy by heat, sound, vibration, and so on. In other words, a higher Tan δ value means that the ratio of dissipated energy is greater than stored energy, indicating excellent damping capacity. As shown in Fig. 6, it can be seen that as frequency increases, the damping capacity is lower in almost all temperature ranges in all alloys. It was found that the present alloys showed similar trends in damping capacity with temperature and frequency, which is that the damping capacity increased exponentially as the temperature increased and the applied frequency decreased. Moreover, the slope of Tan δ, which increases exponentially at higher temperatures according to the addition of Cu, tended to become steeper at higher temperatures. In addition, these present alloys similarly showed little dependence on a frequency below 400 K but above 400 K confirmed that the damping capacity tended to decrease as the applied frequency increased.

4. Discussion

Based on the results of Fig. 1 and the ternary phase diagram calculated using Thermo-Calc software in the literature [20,38], the solidification sequences of the alloys can be characterized as follows.

As indicated by the EDX results, the primary Al phases exhibit increased Cu content due to the rapid cooling rate of this casting setup, estimated at approximately 5200 K/s (see appendix A), indicating a state of supersaturated Cu in the Al phase. In addition, the simple binary Al₃Ni₂ intermetallic compound demonstrates a wide range of compositional variations, expandable to the ternary system such as the Al₃(NiCu)₂ intermetallic compound. The Al₃(NiCu)₂ phase shows that the ratio between Cu and Ni is in the range of 0.6–1.4, with total amount of Cu and Ni content around 40 at.%. On the other hand, the Al₇Cu₄Ni intermetallic compound exhibits higher Cu + Ni contents than Al₃(NiCu)₂ with the ratio between Cu and Ni ranging in from 3 to 4 [39]. Based on these findings and the results obtained from this study, including EDX and XRD analyses, the phases comprising each alloy can be accurately indetified.

Furthermore, the lattice strain generated by compositional change-induced lattice distortion. Analysis of XRD peaks, such as their shifting and broadening, enables calculation of lattice strain resulting from lattice distortion [40,41]. The compositional change in solid-solution and the solution-induced lattice distortion can be classified into homogeneously strained (uniformly) and inhomogeneously strained (non-uniformly), resulting in the shifting of peaks and an increase in full width at half maximum (FWHM) by broadening, respectively. Parameters such as crystallite size, dislocation density, and strain can be estimated using Scherer's formula and the Williamson-Hall method. Scherer's equation addresses XRD peak broadening by crystallite size as:

where λ is the wavelength of radiation (Å) used in the test, β is the instrumental fitted FWHM in radians, the θ is the Bragg's angle of diffraction peaks.

The lattice strain can be estimated by the Williamson-Hall method as follows:

Simply, the dislocation density can be expressed as δ = FWHM⁻².

The estimated crystallite size (t), dislocation density (δ), and lattice strain ($\mathrm{\epsilon })$ using the above equations are presented in Table 3. The estimation used the results calculated in (111)_hkl and (200)_hkl of α-Al from the Voigt-fitted XRD pattern (the inset of Fig. 1) because the Voigt-fitted curve has superior accuracy for a single peak compared to the commonly used Gaussian fitting [41].

As shown in Fig. 1 and Table 3, the diffraction peaks corresponding to α-Al show a higher angle shift and the increase of FWHM by Cu addition. This phenomenon is attributed to the substitutional incorporation of Cu into α-Al, causing the lattice distortion. The increase in FWHM was remarkable at (200)_hkl and increased by about 30 % according to the increase in Cu. As shown in Table 3, the lattice strain calculated via the Williamson-Hall method (Eq. (2)) increased from 2.0616 to 2.6718 based on (200)_hkl with increasing Cu, indicating significant lattice distortion of α-Al due to Cu addition. This lattice distortion creates a stress field that impedes dislocation movement, thereby enhancing strength, known as the solid-solution strengthening effect. Furthermore, these results predominantly stem from the changes in composition, as the specimens were processed using the same casting process which have the consistent cooling rate.

As shown in Fig. 3(a) and (b), the fracture surface of AN973 alloy exhibits predominant dimples, indicating a ductile mode of fracture. Despite the presence of Al3Ni intermetallic regions, crack propagation at these interfaces, which could potentially lead to alloy failure, was not observed. In general, a reduction in the brittleness of alloy containing intermetallic compounds can be achieved by maximizing their interface length scales, e.g. nanoscale [1,15]. In this case, the dislocation density at the interface can be reduced, thereby enhancing plasticity. Even though the Al₃Ni intermetallic compound in AN973 has a rather higher volume fraction, it can suppress an increase in brittleness and maintain excellent plasticity without any evidence of strain localization at the interface.

On the other hand, in ANC alloys containing Cu, strength significantly increases due to the relatively higher the volume fraction of eutectic structure (lamellar) and the formation of new intermetallic compounds such as Al₇Cu₄Ni, Al₃NiCu, and Al₂Cu. In general, intermetallic compounds exhibit much higher hardness than the solid-solution, and the eutectic structure, i.e. fine lamellar structure, with numerous interfaces contributes to high strength [14,22,25,42]. In other words, as shown in Fig. 5(b) the hardness in the as-cast state increases from 80 HV → 107HV → 157 HV with increasing Cu contents, and the tensile yield strength of present alloys also increases in a similar trend from 170 MPa → 270 MPa → 440 MPa. This underscores the solid-solution strengthening effect of α-Al combined with strengthening effects from second phases and dispersion of new phases, crucial for enhancing alloy strength. As shown in Fig. 3, α-Al phase in the solid-solution and eutectic structure mainly shows dimples and evidence of plastic deformation, whereas mirror-like cleavage fracture surfaces are observed elsewhere. This suggests stress concentration at interfaces of brittle intermetallic compounds after slip-band deformation of soft α-Al, contributing to crack propagation and eventual failure of the alloy. This trend is notably pronounced in ANC10, indicating reduced plasticity due to increased volume fraction and length scale of brittle intermetallic compounds with associated strengthening effects [10,12].

Furthermore, Cu-containing alloys such as ANC5 and ANC10 can achieve further precipitation hardening with appropriate aging as evidenced in Fig. 4, Fig. 5. In general, Al alloys containing supersaturated α-Al with Cu form Al₂Cu through artificial aging. The Al₂Cu formed through the artificial aging exhibits superior resistance to dislocation movement compared to α-Al and thus shows a precipitation hardening effect. The increase in hardness over the aging time (Fig. 5) is strong evidence to support this mechanism. Hence, ANC5 and ANC10 benefit from both precipitation hardening and the previously discussed solid-solution hardening and second-phase/dispersion hardening effects.

The precipitation induced by artificial aging is inevitably accompanied by the movement of atoms. Especially, it is notable due to the high diffusivity of Cu in the Al matrix [21,43]. This phenomenon can potentially influence their damping properties, as indicated in Table 2, which presents Tan δ values (nominal values) measured every 100 K from 400 K to 700 K. It is a well-established trend that damping capacity increases with increasing temperature, which is similar to the many crystalline metallic materials [4,14,[44], [45], [46], [47]]. Damping capacity, defined as the ratio of storage modulus to loss modulus, reflects the energy dissipation (“loss”) mechanisms associated with various structural relaxations (e.g., Snoek and Zener types) and defect relaxations, including dislocations, interfaces, and vacancies [[48], [49], [50]]. The Snoek-type is found to explain the carbon in α-Fe induced damping behavior. This relaxation originated from the movement of interstitial atoms such as C, H, N, etc. The Zener-type means the stress-induced reorientation of atomic pairs or clusters. In other words, many explanations and descriptions for the damping behavior of conventional metallic materials are mostly based on the relaxation of defects. Along this line, as temperature increases, the enhanced mobility of atoms involved in these relaxation processes facilitates easier dissipation of externally applied loads, thereby increasing damping capacity.

AN973 shows higher mechanical stability at relatively high temperatures, attributed to the enhanced phase stability of Al₃Ni intermetallic compound and the limited solubility of Ni in Al, preventing the formation of supersaturated solid-solution phases. Therefore, in the case of AN973, slight decreases in hardness with aging time (Fig. 5) are attributed to grain growth during aging, while the damping capacity increases according to the thermally increased atomic mobility.

In contrast, the steep increase in Tan δ with Cu addition at the similar temperature is influenced by the presence of supersaturated α-Al with Cu, which precipitates as Al₂Cu around 500 K. The precipitation behavior necessarily accompanies the atomic movement, consuming external energy and contributes to a high damping capacity. Furthermore, the deviation of Tan δ with frequency change at a specific temperature enlarges as the Cu contents increase, indicating that the frequency dependence rises as the Cu addition. The rapid increase in the Tan δ post-precipitation temperature suggests that the low melting temperature and the phase selection may change the nature of damping behavior (relaxation behavior) by Cu addition.

5. Conclusion

In this study, we investigated the enhancement of tensile strength through solid-solution hardening and dispersion-strengthening effects of Cu addition in Al-Ni eutectic alloys, along with additional strengthening via precipitation hardening induced by heat-treatment. The deformation behavior of the alloys was also analyzed by examining the fracture surfaces post-tensile testing. Furthermore, we examined the temperature and frequency dependence of damping capacity. The main points of this study are summarized as follows:

The addition of Cu to Al-Ni eutectic alloys significantly improves the tensile strength through solid-solution hardening and dispersion (second phase) strengthening effects. The tensile strength increases from 172 MPa in the base Al-Ni alloy to 420 MPa in the 10 at.% Cu alloy, demonstrating a clear relationship between Cu content and mechanical strength.

The heat-treatment of Cu-containing alloys resulted in the precipitation of Al₂Cu phase within the α-Al solid solution, contributing to an additional strength increase of approximately 10 %. This precipitation hardening effect induced by the artificial aging process is a key mechanism for further improving the mechanical properties of these alloys.

The fracture surface analysis results show that plastic deformation mainly occurs within the primary α-Al and eutectic Al phases, which is evidenced by the presence of dimples and necking on the fracture surface. The ultimate fracture was associated with crack formation and propagation in Cu-containing intermetallic compounds including Al₂Cu, Al₃(NiCu)₂, and Al₇Cu₄Ni.

The study found that the addition of Cu enhanced the damping capacity of the alloy at high temperatures. This is due to the precipitation of Al₂Cu and the higher diffusivity of Cu in the Al matrix, which facilitates atomic movement and contributes to better energy dissipation at high temperatures.

These conclusions highlight that Cu plays a dual role in enhancing both the mechanical strength and damping capacity of Al-Ni eutectic alloys, making them suitable for advanced applications that require high strength and effective energy dissipation at high temperatures.

Moreover, the study identified that Cu addition facilitates atomic movement at elevated temperatures, notably through the precipitation of Al₂Cu and the higher diffusivity of Cu in Al, thereby enhancing damping capacity at high temperatures.

Data availability statement

The data included in article/supp. Material/referenced in article.

CRediT authorship contribution statement

Jeong Tae Kim: Writing – original draft, Supervision, Resources, Project administration, Methodology, Investigation, Formal analysis, Data curation, Conceptualization. Sung Min Park: Writing – review & editing, Visualization, Validation.

Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

There is no financial interests.

Appendices.

Appendix B. Supplementary data

The following is the Supplementary data to this article.