1 Introduction

The use of aluminum alloys is of great interest in automotive and industrial applications that depend on the weight-to-strength ratio ^{[1, 2, 3]}. Al-Si-Cu systems ^{[4, 5]}, among them the AA319, can further improve their mechanical properties by standard heat treatments ^{[6, 7]}, by the Al₂Cu (θ') nano-precipitates and Al₂Si ( ϵ-Si) micro-precipitates phases ^{[8, 9]}. Although improvements in the mechanical performance of Al-Si-Cu alloys due to phase-related changes of θ' after various heat treatments, their engineering applications remain limited to temperatures below 150 [¹⁰]. A necessary additional characteristic to modern applications is lightweight to work at higher temperatures. However, Al-Si-Cu systems can be modified in their chemical composition through additional alloying elements. In this regard, Ni, which generates Al-Ni intermetallic compounds that exhibit high thermal stability and high melting point, is an excellent option for aluminum alloys. It has been reported that additions of 1 to 2% Ni to 2xx and 3xx series alloys enhance hardness and tensile properties at elevated temperatures [^11,12], due to precipitation of Al₃Ni, Al₃ (CuNi)₂, and Al₇Cu₄Ni intermetallics [^13,14]. Specifically, to AA319, Ni improved UTS by forming Fe-rich intermetallics, reducing this impurity in the alloy. Fe-Ni intermetallics, as Al₉FeNi, have a relatively high thermal stability at typical temperatures used in AA319 heat treatments, so they would not be affected by subsequent heat treatments.

This research studied the Nickel effect as an additional alloying element in aluminum alloy 319 alloys regarding microstructure and mechanical properties in as-cast, T6 heat treatment, and overheating conditions.

2 Experimental procedure

The raw materials were commercial AA319and Al-20Ni (wt.%) master alloys. The AA319 alloy was melted at 740 in a graphite crucible using a Lindberg Blue electric furnace. Al-20Ni master alloy was added in different proportions to obtain +1, +1.5, and + 2%Ni modified alloys. Synthesized alloys were degassed for 5 minutes using Ar gas and a graphite propeller at 490 rpm, then 0.33 (wt.%) of Al-5Ti-B grain refiner was added. Table I contains elementary quantifications obtained by inductively coupled plasma-atomic emission spectroscopy.

Mechanical properties were measured by tensile and hardness tests using as-cast and heat-treated samples; this last case considered room temperature and 250 conditions. Regarding room temperature tests, the thermal history of samples includes four categories:

A Panalytical X’Pert PRO diffractometer was used for the initial characterization by X-ray diffraction (XRD) technique with Cu radiation at 40 kV and 35 mA. Samples prepared by conventional metallographic techniques and chemically attacked using 20 _HCl04 /80 _MeOH were analyzed by scanning electron microscopy (SEM), JEOL 5800-LV operated at 20 kV. In the transmission electron microscopy (TEM), we used PHILIPS CM-200, and JEOL JEM2200FS operated at 200 kV. Metallographic preparation was based on electropolishing at -30 using 30 _HCl03 /70 _MeOH electrolyte solution and 20 V. Samples to tensile tests were manufactured according to the E8-ASTM standard and tested in an Instron-337 universal machine at 0.1 mm/s. Evaluation hardness in a Rockwell-Wilson tester was done and reported on the Rockwell-B scale.

3 Results

In Fig. 1, XRD patterns show the peaks corresponding to the significant constitutive elements Al, Si, and Cu. Quantification of metallic-Cu was slightly higher in T6 than as-cast. The Ni addition does not change this behavior. Cu clusters suffered significant changes due to heat treatment, where no significant variations were observed compared to the AA319 and AA319 + 2%Ni. The Al₂Cu phase was observed in AA319 after T6 heat treatment. An unexpected AA319 + 2%Ni A₂lCu phase was obtained in a lower concentration.

Figure 1 X-ray diffraction patterns for as-cast and T6 conditions: a) Al-319, b) +2%Ni. Crystal structures were determined using the Rietveld method.  

To AA319 + 2%Ni, additional phases were indexed corresponding to expect AlNi and AlCu. The T6 heat treatment does not modify Ni intermetallic compounds because of thermal stability, which is related to an expected increment in the hardness performance.

SEM micrographs (Fig. 2a)) of the as-cast AA319 show that α-Al precipitates do not present a clear dendritic arrangement morphology, obtaining a similar microstructural arrangement that the observed for as-cast AA319 + 2%Ni (Fig. 2b)). ϵ-Si precipitates were rounder in shape and smaller than the case of as-cast AA319 + 2%Ni than in commercial AA319, while ϵ-Si precipitates appear in block-shape morphology after Ni addition.

Figure 2 SEM micrographs of as-cast and T6 heat treatment conditions; a) as-cast AA319, b) AA319_T6, c) as-cast AA319 + 2%Ni _T6 , d) AA319 + 2%Ni _T6 . The EDS results in Figs. c), d), and g) correspond to the representative phases found in more significant quantities in each thermal state and chemical composition.  

Al₂Cu clusters were identified as coarsening particles in the as-cast stage. Intermetallic particles were identified along Al-matrix containing as main elements: Cu, Fe, Mn, and Ni [¹⁵]. In the case of AA319 + 2%Ni (Fig. 2e)), Al₃Ni in ribbon-like morphology was the most notable effect after Ni addition [¹⁶]. Ni intermetallics precipitated with Fe and Cu elements [¹⁷]. Microstructural changes due to T6 heat treatment were observed following the Oswald ripening process, a notable effect of coarsening and rounded ϵ-Si in AA319 _T6 until remained some sharpest in AA319 + 2%Ni _T6 (Fig. 2f)). Elemental analysis is included for phases commonly present in alloys in Fig. 2c), d) and g).

Regarding Cu clusters, there was a similar amount of Cu dissolved in α-Al after Ni addition, which remained without significant changes in both AA319 _T6 and AA319 + 2%Ni _T6 . Additionally, in both solution temperatures schemes (495 and 520), Cu clusters were maintained in similar amounts and sizes. The morphology of Al₉FeNi was modified in Ni samples, differing from undesirable Fe phases with sharp edges.

Figure 3 characterized by TEM the evolution of micro- and nano- precipitates in AA319 and AA319 + 1%Ni at T6 and over-heating thermal stages. T6 heat-treated samples in Figs. 3a) and 3c) show dark phases of irregular morphology, corresponding to high Si content precipitates with a similar composition range. AA319 _T6 (Fig. 3a)) contains a high density of Al₂Cu-θ' precipitates in needle-like morphology. However, in AA319 + 1%Ni _T6 , the presence of needles-pairs with a relative angle of 90 was increased, as seen in Fig. 3c). The average thickness of θ' in AA319 + 1%Ni _T6 was close to 60%, a smaller amount than those needles contained in AA319 _T6 . Interestingly, even Ni and Cu’s affinity at the nanoscale level is not observed as the dominant interaction to form Ni-Cu phases in competition with Al₂Cu needles.

Figure 3 TEM micrographs of a) AA319 _T6 , b) AA319 _OH1 , c) AA319 + 1%N _T6 , d) AA319 + 1%N _OH . In all cases, elemental analyses are listed and labeled as Spectrum and the studio number.  

The TEM micrograph presented in Fig. 3b) reveals that high temperatures, described in the experimental section, produce thermal degradation of the AA319. AA319 _OH sample shown in Fig. 3b) a representative area where θ' needles have been coarsened, retaining a minimal quantity in the matrix. In the evolution of the nanostructure of AA319 + 1%N _OH , Fig. 3d) differs from AA319 _OH due to the higher amount of θ' phases. Even though θ' needles in AA319 + 1%N _OH retain the morphology and high aspect ratio, compared to T6 heat treated thermal state, they are coarser and retain a close and similar spatial distribution.

Figure 4 concentrates hardness and tensile test measures. The presence of hard Ni phases should be enough to increase hardness value regardless of thermal history; thus, this was observed since the as-cast condition, as seen in Fig. 4a). It was observed a gradual increment in hardness with Ni content additions in modified alloys. In the T6 condition, hardness variation is not significant since changes were just a few units. It was observed mechanical degradation by over-heated condition after high-temperature condition, with increment in hardness as a function of Ni content. The hardness of AA319 + 2%N _OH increased significantly from its previous thermal condition, T6. AA319 + 1%N _OH was enhanced around 10% relative to the reference alloy AA319 _OH and 5% compared to the maximum of AA319 _T6 .

Figure 4 Rockwell B Hardness and UTS values for all chemical compositions; AA319, AA319 + 1%Ni, AA319 + 1.5%Ni, and AA319 + 2%Ni. Used thermal stages were; as-cast, T6, over-aging, and high temperature for two different solubilization temperatures HT1 (495) and HT2 (520).  

In Fig. 4b, tensile tests are shown: in the as-cast condition, AA319 + 1%Ni sample obtained the highest UTS values of 218 MPa, followed by AA319, +1.5%Ni, and +2%Ni (178 MPa). After T6 thermal treatment, AA319 + 1.5%Ni _T6 and AA319 + 1%Ni _T6 samples present the highest UTS values of 278 and 267 MPa, respectively. In comparison, the AA319 value was 246 MPa; over-heating conditions, the deteriorating effect in AA319 _OH and AA319 + 1%N _OH samples were evident. For the AA319 + 1.5%N _OH and AA319 + 2%N _OH samples, UTS values were 260 and 264 MPa, respectively, evidencing the effect of the addition of Ni. UTS reaches maximum values at the T6 stage for AA319 _T6 , AA319 + 1%%N _T6 , and AA319 + 1.5%N _T6 ; however, a pronounced drop of every maximum peak was observed in the over-heating condition. The maximum UTS was obtained at over-heating for AA319 + 2%N _OH . UTS values until high-temperature conditions were maintained in the range of 136 MPa (AA319 + 1%N _HT1 ) to 182 MPa (AA319 + 2%Ni _HT1 ). Regarding the higher solubilized temperature, HT2, the range was 112 (AA319 + 1%N _HT2 ) to 162 MPa (AA319 + 2%N _HT2 ).

4 Discussion

Commercial aluminum alloy 319 (AA319) after T6 heat treatment reduced the amount and size of Cu clusters due to the solubilization process activated by high temperature [?]. In Ni-modified samples, a significant reduction in the size and quantity of mentioned Cu-clusters was presented. Regarding indexed crystal phases, XRD showed a quantity decrease in θ'-Al₂Cu due to Ni addition. Such results contrasted with TEM observations showing a higher density of θ' to all modified samples. This behavior can be attributed to XRD differences related to reducing the coarseness of θ' phases that precipitated in Ni samples. Ni phases formed in modified samples consisted of Al₃Ni, identified at the microscale dimension and indexed by XRD. Ni precipitates must consist of a range of other reported phases. However, just Al₃Ni was in the range of XRD sensitivity. θ'-Al₂Cu dimensional reduction with the highest number of needles could improve UTS due to the highest number of obstacles to nanofractures diffusion. An essential difference between AA319 and Ni modified samples showed θ' needles in AA319 with higher density. In the case of Ni modified samples, there were precipitated in two different directions creating a high amount of θ' pairs. θ' mainly drove the hardness behavior until needle pairs were not influenced positively because of hardness depending on the global value of α-Al deformation. It can be rationalized that T6 results in θ' needles, which are the principal hardness enhancement, and Ni precipitates were screening contribution. The 250 overheating conditions did not significantly modify the hardness values, and therefore, the presence of the Al₃Ni intermetallic was responsible for thermal stability. Due to that, a dynamic transformation was not observed; hardness decreased due to the coarsening of ϵ-Si and θ'-Al₂Cu phases. Hardness increased at the as-cast state was expected due to hard Ni precipitates in modified alloys. However, the T6 heat treatment condition and hardness behavior depended on θ' formation and presence as the principal hardness phase. Hardness decrement was observed at over-aging due to θ' phase coarsening as observed by TEM micrograph (Fig. 3b).

On the other hand, the retaining of θ' in Ni modified samples (Fig. 3d) was responsible for the favorable mechanical response. Copper-rich nanostructures were responsible for hardness enhancement. Thus Ni precipitates attenuated detrimental effect on UTS due to the increasing number of θ'-Al₂Cu needles. After the high-temperature event, the maximum hardness of AA319 + 2%N _OH could be reasonable to reduce θ'-Al₂Cu kinetics of precipitation by Ni presence.

The mechanism, such as Ni interacting with Cu to reduce the size of the needles, was not clear after using TEM micrographs at solubilized samples after 1 to 4 hrs. However, it was supposed a reduction of kinetics transformation from GP to θ'. Thin needles precipitated due to Ni additions under the same heat treatment as AA319 _T6 , however, coarsening of θ' at 250 occurred at all samples. Improvement of Ni samples was because the size of needles was smaller before starting the high temperature, thereby delaying the alloy’s deterioration.

Al-Ni precipitates are hard and brittle particles, which can decrease UTS in the way that occurred in AA319 + 2%N _AC . However, the increase of Al₂Cu phases positively counterbalanced the presence of Ni phases in a significant way. UTS value increased by Ni addition due to the higher density of Al₂Cu with a smaller dimension. The observed drop in UTS values at high temperatures and long times is due to the coarsening of Al-Cu and Al-Si precipitates, which quickly screen the improvements done by Ni precipitates. It is a known fact that the coarsening of Al₂Cu phases is presented as a function of the temperature. However, the detrimental effect of transformation kinetics of the θ' phase in this condition was attenuated mainly in the AA319 + 1%N _OH sample.

5 Conclusion

Commercial aluminum alloy 319 and a series of alloys modified with +1.0%, +1.5%, and +2.0% Nickel were used to investigate their micro and nanostructural evolution. The synthesized conditions were as-cast, T6, overheating, and high temperature. By TEM micrographs, it was concluded that the modified alloys reduce θ'-Al₂Cu precipitation kinetics in comparison to the commercial alloy. The direct effect of Ni additions on mechanical properties was an increase in both hardness and UTS. In the latter case, the maximum UTS values at T6 were for AA319 + 1.5%Ni addition, followed by AA319 + 1.0%Ni. After being subjected to 250 for 30 minutes, the AA319 + 1.5%Ni alloy was the only one that preserved the mechanical conditions for engineering applications. Compared to the tests performed on the material at 250, the alloy AA319 + 2%Ni was the least affected, confirming the reduction of θ'-Al₂Cu phase kinetics.