# Change of Tensile Properties with Aging Time and Temperature in Al-Si-Cu-Mg 354 Cast Alloys with/without Minor Addition of Ni and/or Zr

Authors: J. Hernandez-Sandoval; M. H. Abdelaziz; E. A. Elsharkawi; A. M. Samuel; F. H. Samuel
Journal: Advances in Materials Science and Engineering
 (2021)

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## Abstract

The principal aim of the present research work was to investigate the effects of minor additions of nickel and zirconium on the strength of cast aluminum alloy 354 at room temperature (25°C). A decrease in tensile properties of ∼10% with the addition of 0.4 wt.% nickel is attributed to a nickel-copper reaction which interferes with the formation of Al2Cu precipitates. The platelet-like phases (Al,Si)3(Zr,Ni,Fe) and (Al,Si)3(Zr,Ti) are the main features observed in the microstructures of the tensile samples of alloys with Zr additions. The reduction in mechanical properties is due to the increase in the percentage of intermetallic phases formed during solidification; such particles would act as stress concentrators, decreasing the alloy ductility. The main effect of Zr addition lies in a significant reduction in the alloy grain size ∼40%, rather than an increase in the mechanical properties. Quality index charts could be used in assessing the effects of the Ni and Zr additions to the base alloy, as well as evaluating the heat treatment relationships to the alloy tensile properties, in particular when the system shows multiple precipitation reactions. Due to the high liquidus temperature of the Al-Zr binary phase diagram, addition of Zr beyond 0.2% is not recommended to avoid undissolved Zr.


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## Body

## 1. Introduction

The aging treatment of an aluminum alloy containing copper as the alloying element results in the formation of various forms of Al-Cu-containing precipitates. The precipitation sequence of an Al-Cu alloy during the aging process has been proposed as follows [1–3]:(1)αSSS⟶GP zones⟶θ″⟶θ′⟶θ.The coherent and semicoherent phases,θ” and θ′, respectively, contribute to increasing the strength of the alloys. On the other hand, the incoherent equilibrium precipitate θ (Al2Cu) results in diminishing the hardening level of the alloys because of the loss of coherency between the stable phases and the metal matrix. With regard to an aluminum alloy containing Cu and Mg as the hardening elements, the aging treatment results in the formation of a range of precipitates. The generally accepted sequence of precipitation in Al-Cu-Mg alloys is [4,5] θ-Al2Cu precipitate and its precursors, along with other hardening phases/precipitates such as the S-Al2CuMg phase and its precursors, which are observed to form during the aging treatment of an aluminum alloy containing Cu and Mg. The formation of the S phases during aging treatment occurs via the following precipitation sequence [6–10]:(2)αSSS⟶GPB zones⟶S″⟶S′⟶S,starting with the decomposition of the supersaturated solid solution (SSS), followed firstly by the formation of GPB (Guinier–Preston–Bagaryatsky) zones and then by the precipitation of the coherent S″ and the semicoherent S′ and finally the incoherent equilibrium S (Al2CuMg) phase. The coherent S″ and the semicoherent S′ precipitates are the main source of hardening in aluminum alloys containing Cu and Mg, and they are also responsible for increasing the strengthening level of this type of alloy, whereas over aging conditions result in the precipitation of the incoherent equilibrium S phase. Moreover, there is another phase containing Si which may form in the Al-Cu-Mg-Si system, called the Q phase or quaternary Al4Mg8Cu2Si6 phase which can form upon solidification or during aging [11–14]. This Q phase can also coexist with Al2Cu, Mg2Si, and Si depending on the ratio between Cu, Mg, and Si. One of the suggested mechanisms for the formation of the Q phase is that the Cu atoms dissolve in the β” which then evolves either to β (Mg2Si) or to the Q phase based on the chemical composition of the alloy and precipitation status [15, 16].The Mg/Si ratio, the Cu/Mg ratio, as well as the Cu and Si contents are all principal controlling factors which determine the main precipitation-hardening phases in Al-Cu-Mg-Si alloys. The strengthening mechanism results from an age-hardening treatment which depends on the interaction between the moving dislocations and the precipitated phases [17–19]. The addition of copper to Al-Si alloys leads to the formation of the copper intermetallic phase, which increases the alloy strength both at room and at high temperatures. This intermetallic phase is present in three forms: (i) eutectic-like Al-Al2Cu, (ii) block-like Al2Cu, and (iii) Q-Al4Mg8Cu2Si6 phases. Silicon alone contributes very little to the strength of aluminum casting alloys yet provides a very effective level of strengthening when combined with magnesium to form Mg2Si. The Mg2Si phase is soluble in the alloy to a limit of ∼0.7% Mg and provides the precipitation strengthening basis for heat-treatable alloys.According to Refs. [20, 21], it would be important to produce a microstructure containing coarsening-resistant dispersoids in order to enhance the alloy mechanical properties at high temperatures in an aluminum alloy. Zirconium has the smallest diffusion flux in aluminum of all the transition metals [22], leading to the formation of the Al3Zr phase, which precipitates out during the initial solution heat treatment in the form of metastable L12 Al3Zr particles. These particles are resistant to dissolution and coarsening; they can also control the evolution of the grain and subgrain structure, thereby making it possible to increase strength and ductility in the precipitation-hardened T6 condition [23] or during subsequent processing operations, such as the hot rolling of wrought alloys.The capacity of nickel for improving the strength of wrought aluminum alloys at high temperatures has already been established; nickel is thus used in combination with copper to enhance high temperature properties [24]. The solid solubility of nickel in aluminum cannot exceed 0.04%. If this amount is exceeded, then it is present as an insoluble intermetallic, usually in combination with iron. Nickel content of up to 2% increases the strength of high-purity aluminum but reduces its ductility. Binary aluminum-nickel alloys are no longer in use, but nickel is added to aluminum-copper and aluminum-silicon alloys to improve both hardness and strength parameters at elevated temperatures as well as to reduce the coefficient of thermal expansion [25].Intermetallic phases in an Al-14Si-3Cu-4.5Ni casting alloy with trace additions of Zr were investigated by Jo et al. [26] who reported on a new Zr-rich intermetallic (Al,Si)3(Zr,Ni,Fe) phase with a tetragonal structure with a = 3.275 Å and c = 15.475 Å. Strengthening mechanisms at ambient and elevated temperatures were studied by Michi et al. [27] in a cast Al-0.11Zr-0.02Si-0.005Er (at.%) alloy containing 2.86 at.% Ni. The strengthening process is composed of (i) incoherent Al3Ni microfibers formed during eutectic solidification and (ii) coherent, equiaxed Al3Zr (L12-structure) nanoprecipitates created on subsequent aging. Tunçay et al. [28] investigated the effects of Cr and Zr additives on the microstructure and mechanical properties of A356 alloy. The authors reported on the formation of Cr/Zr-based intermetallics in precipitates with fish bone-like morphologies.Quality index is a pivotal concept originally developed by Drouzy, Jacob, and Richard [29] who introduced an empirical parameter Q or quality index, to characterize the mechanical performance of Al-7%Si-Mg casting alloys. They related the quality index Q to the tensile strength, UTS, and the plastic strain of the material to fracture, Sf, as represented by the following equation:(3)Q=UTS+dlogSf.Although the concept was developed specifically for alloys 356 and 357, it has occasionally been applied to other alloy systems as well [30–32]. A number of studies on Al-Cu-Mg-Ag alloys, however, showed that in contrast with the linear behavior of alloy 356, a plot of UTS vs. Sf describes a curvilinear contour if the material has undergone aging, as was the case for alloy 201. The Cáceres model assumes that the material can be described applying Holloman’s equation:(4)σ=Kεn.Theq values used in the Cáceres model for the quality index charts make use of equation (5), in view of the fact that necking or localized deformation begins at maximum load, where the increase in stress due to a decrease in the cross-sectional area of the tensile specimen becomes greater than the increase in the load-carrying ability of the metal as a result of strain-hardening. Equation (6) makes it possible to calculate the quality index from the tensile test results based solely on the knowledge of the value of K [30].(5)q=εplasticn,(6)Q=UTS+0.4Klogsf.The principal aim of the present research work is to investigate the effects of minor additions of nickel and zirconium on the strength of the cast aluminum alloy 354 at high temperatures.In order to meet this goal, it is necessary to attain the following objectives:(i)
To determine the tensile properties of the alloys studied at room temperature, using the different aging conditions of temperature and time(ii)
To correlate the results obtained from the room temperature tests with the principal microstructural features observed in the corresponding alloy samples in order to analyze and understand the major parameters involved in the strengthening of alloy 354 at high temperatures


## 2. Experimental Procedure

The chemical composition in wt.% of the as-received 354 alloy is listed in Table1. In order to attain the objectives outlined previously, an experimental work plan was developed as follows:(i)
Castings to be made so as to obtain tension test bars for the 5 alloys investigated;(ii)
Heat treatment of all test bars produced is as follows:(a)
Solution heat treatment for 8 h at 495°C;(b)
Quenching in warm water at 60°C;(c)
Artificial aging at 6 temperatures covering a wide range of aging behavior on-going from under-, peak-, and over aging conditions for times ranging from 2 hours to 100 hours, as shown in Table2;(iii)
Tensile testing of the heat-treated test bars.Table 1
Chemical composition of the as-received 354 alloy.
Element(wt.%)Si9.1Fe0.12Cu1.8Mn0.0085Mg0.6Al87.6Table 2
Artificial aging conditions used for room temperature tension tests.
Temperature (°C)Aging time (h) and aging condition codes24681012162024364872100155123456789101112131701415161718192021222324252619027282930313233343536373839240404142434445464748495051523005354555657585960616263646535066676869707172737475767778The as-received alloy was melted in a 120-kg capacity SiC crucible, using an electrical resistance furnace. The metal temperature was maintained at 780°C (to dissolve maximum Zr [33]), while the melt was degassed using pure, dry argon injected into the melt for 30 min by means of a rotating graphite impeller at 135 rpm. Grain refining and modification of the melt were, respectively, carried out using Al-5% Ti-1% B and Al-10% Sr master alloys to obtain levels of 0.25% Ti and 200 ppm Sr in the melt. It should be mentioned here that to minimize the oxidation of Sr, the Al-Sr master alloy was added 5 minutes before the end of degassing.Three samples for chemical analysis were also taken at the time of the casting; this was done at the beginning, in the middle, and at the end of the casting process to ascertain the exact chemical composition of each alloy. Table3 lists the chemical analysis of the various alloys studied together with their respective codes, as obtained from samplings taken for this purpose from the corresponding melts used in this work.Table 3
Chemical composition of the alloys used in this work (wt.%).
Alloy/elementSiFeCuMnMgTiSrNiZrAlA9.430.081.850.010.490.220.0150∼∼87.5D9.160.081.840.010.490.220.01490.46∼87.7E9.100.071.830.000.450.210.0145∼0.3987.7F9.100.081.860.000.460.220.01220.400.3987.5G9.010.081.850.000.450.210.01270.210.19087.8The melt was poured into a preheated ASTM B-108 permanent mold (preheated to 460ºC) to prepare test bars for the tensile tests. Each mold casting provides two tensile bars, each with a gauge length of 70 mm and a cross-sectional diameter of 12.7 mm, as shown in Figure1.Figure 1
Measurements of the tensile tests.As will be mentioned in a later subsection, heat treatment of the test bars used for tensile testing involves solution heat treating them at 495ºC for 8 h (to avoid incipient melting [34]), followed by quenching in warm water at 60ºC, after which artificial aging is applied according to the plan established in Table 2. After aging, the test bars are allowed to cool naturally at room temperature (25°C). All heat treatments are carried out in a Lindberg Blue M air forced electric resistance furnace.All of the samples, whether as-cast, solution heat-treated, or aged, were tested at 25 ºC to the point of fracture using an MTS servo-hydraulic mechanical testing machine at a strain rate of 4 × 10−4 s−1. The yield strength (YS) was calculated according to the standard 0.2% offset strain, and the fracture elongation was calculated as the percent elongation (%E1) over the 50 mm gauge length, as recorded by the extensometer. The ultimate tensile strength (UTS) was also obtained from the data acquisition system of the MTS machine. This machine was calibrated anew each time any testing was carried out. The average %E1, YS, or UTS values obtained from the 10 samples tested (using bars with standard deviation ± 5) where at least 7 bars were considered to determine the properties representing a specific condition.Porosity percentage was determined using an image analyzer whereas determination of percentage of intermetallics as well as phase identification was carried out using electron probe microanalysis (EPMA) in conjunction with wavelength dispersive spectroscopic (WDS) analysis, using a JEOL combined microanalyzer operating at 20 kV and 30 nA, where the electron beam size was ∼2μm. Mapping of certain specific areas of the polished sample surfaces was also carried out where required, to show the distribution of different elements within the phases.


## 3. Results and Discussion

### 3.1. Microstructure Characterization

Figure2(a) presents the microstructure of alloy A in the as-cast condition showing a large amount of Al2Cu phase as confirmed by the EDS spectrum displayed in Figure 2(b) corresponding to the white circled area in Figure 2(a). Figure 2(c) is a high magnification image of Figure 2(a) revealing the coexistence of the quaternary Al4Mg8Cu2Si6 phase with the Al2Cu phase in addition to precipitation of fine ß-Al5FeSi phase platelets. Solutionizing at 495°C for 8h was proven to be effective in dissolving most of the Cu-containing phase with traces of insoluble Al-Fe-Cu and ß-Fe as shown in Figure 2(d).Figure 2
Microstructure of the base alloy (A): (a) backscattered electron image of alloy A in the as-cast condition, (b) EDS spectrum corresponding to (a), (c) high magnification image of (a), and (d) alloy A following solutionizing at 495°C/8 h.
(a)(b)(c)(d)The microstructure of alloyF (containing 0.4% Ni + 0.4% Zr) is shown in Figure 3(a). As can be seen from Table 3, the amount of precipitated intermetallics has noticeably increased. Figure 3(b) is a high magnification image of Figure 3(a) revealing precipitation of a Ni-containing phase in the form of short platelets. Figures 3(c) and 3(d) are X-ray images of the distribution of Ni and Zr within the platelets indicating that these platelets are the (Al,Si)3 (Ni,Zr,Fe) phase reported by Jo et al. [26]. The composition of these platelets is presented in Table 4. Figure 3(e) illustrates the fracture surface of alloy E showing precipitation of a large Zr-Ti intermetallic during solidification of the alloy.Figure 3
Microstructure of alloyF in the as-cast condition: (a) backscattered electron micrograph, (b) a high magnification of (a), (c) X-ray image of Ni distribution, (d) X-ray image of Zr distribution, and (e) fracture surface of alloy E in the as-cast condition.
(a)(b)(c)(d)(e)Table 4
Chemical composition corresponding to Figure3(e).
ElementWeight %Atomic %Net int.Error %K ratioZRAFMgK0.460.6318.6117.630.001.10.960.691.02AlK58.2971.882,505.573.620.491.060.970.81SiK10.4712.40300.616.930.061.080.970.551.01ZrL14.745.38199.376.300.090.811.150.741.01ScK0.310.235.3759.340.000.941.020.921.04TiK5.583.8888.025.550.050.921.020.951.04NiK6.873.9035.9610.530.070.911.030.991.09CuK3.281.7212.6525.820.030.861.0311.1Table5 summarizes secondary dendrite arm spacing SDAS, grain size, porosity %, and volume fraction of intermetallics for alloys A through G (samples were prepared from the tensile test bars, 10 mm below the fracture surface). Since all alloys were grain refined using Al-5%Ti-1%B master alloy, the grain size revealed noticeable variation from one alloy to another reaching about 40% reduction in the case of alloy E. Similarly, the secondary dendrite arm spacing values did not vary much as the solidification rate was about 8°C/s in all cases. The difference noted in the amount of intermetallics after solution heat treatment is mainly due to dissolution of Al2Cu phase as illustrated in Figures 4(a) and 4(d), where Figure 4 demonstrates the variation in the size and density of the precipitated hardening phase (Al2Cu phase in the present case) with the increase in aging temperature from ultra-fine particles at 170°C (Figure 4(b)-peak aging) to rod-shaped particles at 350°C (Figure 4(d)-over aging) indicating the possibility of the commencement of incoherency.Table 5
SDAS, grain size, porosity %, level of modification, and volume fraction of intermetallics for alloysA, D, E, F, and G.
Alloy code conditionSDAS (μm)Grain size (μm)Porosity (%)Volume fraction of intermetallics (%)AvAvAvSDAvSDA-AC∗19.32010.1350.063.080.32A-SHT23.11920.120.051.270.13D-AC18.21860.1210.074.220.35D-SHT18.61840.1220.062.760.47E-AC18.91220.110.054.230.40E-SHT19.91160.1080.042.860.43F-AC19940.110.055.020.43F-SHT21960.1110.063.740.31G-AC25.31250.120.033.570.23G-SHT21.21110.130.042.030.34∗AC : as cast.Figure 4
Development of precipitation-hardening Al2Cu phase in alloy A during 2h aging at (a) 155°C, (b) 170°C, (c) 240°C, and (d) 350°C.
(a)(b)(c)(d)


### 3.2. Tensile Properties

In this section, the tensile properties of the present alloys following aging treatment at 155°C (under aging), 190°C (peak aging), and 350°C (over aging) will be discussed. The results are presented in Figures5–7. In spite of the relatively low solution heat treatment temperature of 495ºC used for the purposes of this work, increased strength after the treatment is normally to be observed in similar 354 alloys which use a solution heat treatment temperature of 525ºC for 2-3 h period [32–38]. This selected temperature for the solution heat treatment might appear to be overconservative, but it was deemed necessary to avoid the risk of any incipient melting of the copper phases, which could have the potential for deteriorating the mechanical properties of the alloys to a great extent.Figure 5
Variation of the tensile properties as a function of aging at 155°C for times up to 100 h: (a) UTS, (b) YS, and (c) %El. Same legends as in Figure5(a).
(a)(b)(c)Figure 6
Variation of the tensile properties as a function of aging at 190°C for times up to 100 h: (a) UTS, (b) YS, and (c) % El. Same legends as in Figure5(a).
(a)(b)(c)Figure 7
Variation of the tensile properties as a function of aging at 350°C for times up to 100 h: (a) UTS, (b) YS, and (c) % El. Same legends as in Figure5(a).
(a)(b)(c)The strength values, particularly the UTS, increased by as much as 30 pct between the as-cast (AC) and the solution heat-treated (SHT) conditions. This increase in the mechanical properties is related to the changes occurring in the form of the silicon particles which become more rounded, decreasing their aspect ratio and density. Through solution heat treatment and artificial aging, the strength of alloyA was increased by ∼64% over its as-cast strength. From the tensile properties as shown in Figures 5–7, the maximum value for UTS of ∼386 MPa may be obtained with two different sets of aging conditions, i.e., 155°C/100 hrs and 190°C/2 hrs. The greatest mechanical stability for the aging treatments was observed at 155°C. Even though a significant decrease in the alloy strength is not in evidence from 2 to 100 hours at 350°C, nonetheless, the greatest mechanical deterioration will be observed after less than 2 hours of aging treatment. This previous assumption is based on a consideration of the tensile properties of the material upon solution heat treatment, namely, UTS = 305 MPa and YS = 161 MPa, implying that the over aged samples tended to lose between 15% and 40% of their original strength within the first 2 hours of being subjected to aging at 350°C.Considering the yield strength in the solution heat-treated condition, in conjunction with the artificial aging conditions used, it was possible to increase the yield strength value by more than 100 pct. The maximum yield strength was reached after 100 hours at the aging temperature of 155°C and after 2 hours at 190ºC. The yield strength behavior appears to be most stable at 190°C aging temperature, showing a variation of 50 MPa in yield strength values for aging times between 2 and 100 hours. These observations were taken into account for the selection of this specific temperature for experiments designed to investigate how the mechanical properties may be affected by long exposure times.The general behavior of alloyA with regard to ductility displays a decrease from 155°C to 190°C. A lower ductility value is observed for the condition involving 190°C for 10 hours. The greatest ductility values may be observed for aging at 350°C although none of the aging conditions reaches the higher ductility values shown in the solution heat-treated condition. This observation may be considered as evidence that the mechanical behavior displayed by alloy A is common to that of the Al-Si-Cu-Mg alloys whose strength is obtained at the expense of ductility [39, 40].The YS values for alloysA and D at aging temperatures of 155°C and 350°C represent the extreme ends of the range of aging temperatures used for these alloys. A decrease in yield strength as a result of the addition of nickel is noted at the aging temperature of 155ºC, of the order of 20 MPa. In contrast, no appreciable difference in the YS values is observed between alloys D and A at the 350°C aging temperature. This decrease in strength through the addition of nickel is related to the interaction between copper and nickel to form precipitates of Al3CuNi in the microstructure. The Al3CuNi phase was detected in alloys containing nickel, namely, alloys D, F, and G. Since copper, as well as magnesium, determines the precipitation strengthening of the Al-Si-Cu-Mg alloys, the available copper would form Al2Cu precipitates through the solution heat treatment/quenching/aging process, thereby strengthening the cast alloys. The formation of Al-Cu-Ni precipitates would thus subtract part of the copper content available for strengthening; this is the principal reason for which a slight decrease in YS is observed with the addition of 0.4 wt.% nickel.With regard to the base alloyA and alloy D, the maximum strength is reached at low aging temperatures ranging from 155ºC to 190ºC, showing a decrease in ductility with the increase in strength; it should be noted that this is the expected behavior for Al-Si-Cu-Mg alloys [41–45]. It will also be observed that, at low aging temperatures, the effect of intermetallics in the reduction of ductility is more pronounced than it is at high aging temperatures of 350 ºC. The increase in ductility at an aging temperature of 190 ºC, or more, becomes evident since the alloy softens and the effect of the intermetallics on ductility seems attenuated.From Figure6, it is observed that alloy E, containing 0.4 wt.% Zr, does not show any significant change in YS, when compared with the base alloy A. The decrease in the ductility of alloy E with respect to alloy A may be assimilated by observing Table 5 where the level of insoluble intermetallics increases from 1.27 to 2.86 %, between the base alloy and alloy E, respectively. This behavior may also be seen for the long-exposure-time at the aging temperature of 190ºC. In general terms, there is no appreciable increase in strength (UTS and YS) to be observed in alloy E containing 0.4 wt.% Zr although there is, however, a visible decrease in ductility which is greater than 20% for nearly all of the tests performed.To understand the latter observation, it is necessary to consider three characteristics observed from the microstructural analysis of alloyE. Firstly, the amount of insoluble intermetallic phases increased from 1.27% in the base alloy to 2. 86% in alloy E. Secondly, the grain size was reduced from 200 μm in alloy A to 120 μm in alloy E. Both arguments could indicate that Zr available for strengthening was used to form large intermetallics in the form of (Al,Si)3(Zr,Ti) phases, as well as to reduce the grain size. Finally, a third aspect which should be taken into consideration is that during the casting procedures, the temperature of the melt was maintained at 780°C, implying that Zr added in the form of Al-20% Zr master alloy may not be capable of dissolving completely; this is due to the fact that, with 0.4 wt.% Zr, the Al-Zr phase diagram indicates the liquidus temperature at 780°C [41].As a result of examining the negative effects of nickel on the strength of alloyE, it becomes relevant to discuss why the addition of nickel did not reduce the yield strength in alloy F which contains 0.4 wt.% Ni plus 0.4 wt.% Zr. With regard to intermetallics, our earlier studies [46–48] show the precipitation of (Al,Si)3(Zr,Ti) phase with the Ni-rich phase Al9NiFe adhering to it. The Al9NiFe phase was observed in the microstructure of tensile samples, as well as phases which appear similar to the Ni-rich phase were observed to adhere to the Zr-rich phases in the microstructures of the tensile samples of alloy F as demonstrated in Figure 3. These observations may suggest that the copper did not significantly interact with the nickel to form Al3CuNi phases, and as a result, the copper was free to strengthen alloy F, in the same way as the copper was able to strengthen alloy A.Ductility seems to be affected principally at low aging temperatures of 155°C through 190°C. This decrease in ductility may be interpreted by noting that alloyF had the greatest amount of insoluble phases as listed in Table 5; these would then act as stress concentrators during flow stress, thereby limiting the ductility in alloy F [49, 50].AlloyG, containing 0.2 wt.% Zr and 0.2 wt.% Ni, displays greater strength than alloy E, which contains 0.4 wt.% Zr, or greater even than alloy F, which contains 0.4 wt.%. Zr plus 0.4 wt.% Ni. According to the Al-Zr phase diagram [41], with 0.2 wt.% Zr content, the liquidus temperature is approximately 720°C. Taking into consideration that, during casting procedures, the temperature of the melt was maintained at 780°C, it is fairly likely that the added Zr dissolved completely which is essential for strengthening the alloys by means of Zr-rich precipitates [51, 52]. As may be seen from Figure 6, alloy G is the only alloy which shows an increase in strength in comparison with alloy A. In fact, even at the aging temperature of 350°C, the strength values are observed to be greater than those obtained with the base alloy.Figure8 depicts an example of the tensile properties obtained for the five studied alloys at 170°C relative to the values obtained for the A base alloy in the as-cast condition, i.e., subtracting the values obtained for the base alloy A in each case and plotted as △P values on the Y-axis (P = property = YS, UTS or %El), with the X-axis representing the aging time (see Supplementary Materials for 155°C and 350°C aging temperatures). This approach helps to better visualize the effects and interactions of the various additions used and the different heat treatment conditions, so as to arrive at a better understanding of the effects of Ni and Zr on the tensile properties of 354 casting alloys. As may be seen from Figure 8, additions of these elements improve the alloy performance, i.e., the tensile properties at 170°C. Figure 8(a) shows that alloy A exhibits the best contribution to UTS at all aging times. After 10 h aging time, the comparative UTS (or △PUTS) values start to decrease as the aging time increases, indicating that these elements provide little or no improvement in mechanical properties. On the other hand, the comparative yield strength (or △PYS) values of all alloys are improved, as shown in Figure 8(b) and are stabilized after 8 h aging time, while Figure 8(c) shows that the △P%El or percent elongation to fracture values for all alloys is much lower than that of the A base alloy, which is to be expected, in light of the results shown in Figure 8(a).Figure 8
Comparison of tensile properties ofA, D, E, F, G alloys relative to those of A alloy in the as-cast condition: (a) UTS, (b) YS, and (c) % El at 170°C. Same legends as in Figure 8(a).
(a)(b)(c)Figure9 compiles the strength-ductility values obtained for all the 354 alloys investigated. These values fall into regions corresponding to peak aging, under aging, and over aging. Using regression analysis, the UTS versus %El relationships may be expressed as follows:UTS = –38.26x + 408.1 with a fit of R2 = 0.82 (alloy A)UTS = –45.0x + 373.0 with a fit of R2 = 0.87 (alloy F)Figure 9
The UTS vs. % El relationship for (a) alloyA and (b) alloy F.
(a)(b)It is evident that the slope of alloyA is higher than that for alloy F. Also, the UTS levels are higher than those for alloy F due to the higher volume fraction of intermetallics in the latter as noted from Table 5.


### 3.3. Q Charts

This section will be limited to alloyA. The Cáceres method which is used here for constructing the quality index charts involves the use of a single value of the strength coefficient (K) for all the conditions appearing in the charts. In order to obtain the values of n and K described in equations (5) and (6), a series of stress-strain curves were used as shown in Figure 10, whereas their aging conditions are listed in Table 6. Based on the analysis of these curves, the values of n = 0.193 and K = 530 were obtained. The reason for taking only 15 conditions out of a total of 80 was principally to avoid an overcrowded diagram. The curves were divided into three groups. The criterion applied for using these 15 conditions (including the as-cast and SHT conditions) was to take into consideration the extreme values, particularly those regarding the % elongation and yield strength values.Figure 10
A series of stress-strain diagrams used to calculate the n and K values (see Table6 for details).
(a)(b)(c)Table 6
Aging condition for curves shown in Figure10(a)–10(c).
Figure #Curve #Aging condition10(a)1As cast2SHT10(b)12 h/155°C22 h/170°C310 h/155°C4100 h/155°C52 h/190°C612 h/170°C7100 h/155°C10(c)12 h/350°C2100 h/240°C310 h/240°C42 h/240°C5100 h/190°C610 h/190°CFigure11(a) shows that the plastic strain and the quality index (Q) both display a significant increase upon solution heat treatment in spite of the fact that the relative ductility (q) attains a value of 0.31 implying that alloy A in its solution-treated condition has reached 31% of its maximum quality index value (Q). When the ductility increases sharply from the as-cast to the solution heat-treated condition, such changes can be related to the spheroidization of silicon particles and to the uniformity of the microstructure in the solution heat-treated condition as illustrated in Figure 2(d). Figure 11(b) demonstrates the Q-chart for alloy A aged at 155°C for times up to 100 h. As can be seen, most of the points are falling in a narrow range of Q (416-366 MPa). The maximum value of Q was achieved after aging for 2 h with minimum UTS value of 308 MPa, and a minimum Q value was obtained with a maximum UTS value of about 387 MPa after 100 h. Apparently, alloy ductility plays a major role in determining the Q level.Figure 11
Quality index charts of alloyA at different aging temperatures: (a) as-cast and SHT, (b) 155°C, (c) 170°C,(d) 190°C, and (e) 350°C.
(a)(b)(c)(d)(e)TheQ charts following aging at 170°C and 190°C are shown in Figures 11(c) and 11(d), respectively. Apparently, there is more than one C-curve due to multiple reactions that took place during the aging process. From Figure 11(c), the values of Q decreased from 421 MPa (2 h) to a value of about 375 MPa (100 h). The stress-strain curve corresponding to minimum Q value is displayed in Figure 10(b)-curve #6. With aging at the 190°C/2 hr condition, GP zones and θ′ precipitates develop such that they are coherent with the matrix; YS is thus increased because of the strain fields associated with these precipitates, which are sheared by the dislocations and subsequently lead to a low strain-hardening rate as shown in Figure 10(b)-curve #5.Aging at 350°C resulted in a significant improvement in the alloy ductility. The chart in Figure11(e) could be divided into two curves (solid line) or simple straight line 2h-100 h (broken line). During the course of aging, the values of UTS level were reduced from 221 MPa (2 h) to 200 MPa (100 h) with significant increase in % elongation from about 4.5% (2 h) (curve #1 in Figure 10(c)) to 6.5% (100 h) resulting in Q values of 350-MPa (2 h)–345 MPa (100 h). In other words, practically there is no change in the Q levels for the entire aging time (±5 MPa).It may be deduced from Figure4 that the θ-CuAl2 precipitates are formed at temperatures of 190°C and higher, i.e., 350°C [53–55]. These particles are incoherent with the matrix and cannot be cut by dislocations, thereby promoting lower strength values and a high strain-hardening rate resulting from the accumulation of Orowan loops around the copper phase particles. As the strain is increased, however, the gradual development of primary shear loops generates intense stress fields around the strengthening precipitates, which are in themselves limited [45, 56, 57] by the activation of a secondary dislocation cross-slip process, thus reducing the strain-hardening ability of the material.


## 3.1. Microstructure Characterization

Figure2(a) presents the microstructure of alloy A in the as-cast condition showing a large amount of Al2Cu phase as confirmed by the EDS spectrum displayed in Figure 2(b) corresponding to the white circled area in Figure 2(a). Figure 2(c) is a high magnification image of Figure 2(a) revealing the coexistence of the quaternary Al4Mg8Cu2Si6 phase with the Al2Cu phase in addition to precipitation of fine ß-Al5FeSi phase platelets. Solutionizing at 495°C for 8h was proven to be effective in dissolving most of the Cu-containing phase with traces of insoluble Al-Fe-Cu and ß-Fe as shown in Figure 2(d).Figure 2
Microstructure of the base alloy (A): (a) backscattered electron image of alloy A in the as-cast condition, (b) EDS spectrum corresponding to (a), (c) high magnification image of (a), and (d) alloy A following solutionizing at 495°C/8 h.
(a)(b)(c)(d)The microstructure of alloyF (containing 0.4% Ni + 0.4% Zr) is shown in Figure 3(a). As can be seen from Table 3, the amount of precipitated intermetallics has noticeably increased. Figure 3(b) is a high magnification image of Figure 3(a) revealing precipitation of a Ni-containing phase in the form of short platelets. Figures 3(c) and 3(d) are X-ray images of the distribution of Ni and Zr within the platelets indicating that these platelets are the (Al,Si)3 (Ni,Zr,Fe) phase reported by Jo et al. [26]. The composition of these platelets is presented in Table 4. Figure 3(e) illustrates the fracture surface of alloy E showing precipitation of a large Zr-Ti intermetallic during solidification of the alloy.Figure 3
Microstructure of alloyF in the as-cast condition: (a) backscattered electron micrograph, (b) a high magnification of (a), (c) X-ray image of Ni distribution, (d) X-ray image of Zr distribution, and (e) fracture surface of alloy E in the as-cast condition.
(a)(b)(c)(d)(e)Table 4
Chemical composition corresponding to Figure3(e).
ElementWeight %Atomic %Net int.Error %K ratioZRAFMgK0.460.6318.6117.630.001.10.960.691.02AlK58.2971.882,505.573.620.491.060.970.81SiK10.4712.40300.616.930.061.080.970.551.01ZrL14.745.38199.376.300.090.811.150.741.01ScK0.310.235.3759.340.000.941.020.921.04TiK5.583.8888.025.550.050.921.020.951.04NiK6.873.9035.9610.530.070.911.030.991.09CuK3.281.7212.6525.820.030.861.0311.1Table5 summarizes secondary dendrite arm spacing SDAS, grain size, porosity %, and volume fraction of intermetallics for alloys A through G (samples were prepared from the tensile test bars, 10 mm below the fracture surface). Since all alloys were grain refined using Al-5%Ti-1%B master alloy, the grain size revealed noticeable variation from one alloy to another reaching about 40% reduction in the case of alloy E. Similarly, the secondary dendrite arm spacing values did not vary much as the solidification rate was about 8°C/s in all cases. The difference noted in the amount of intermetallics after solution heat treatment is mainly due to dissolution of Al2Cu phase as illustrated in Figures 4(a) and 4(d), where Figure 4 demonstrates the variation in the size and density of the precipitated hardening phase (Al2Cu phase in the present case) with the increase in aging temperature from ultra-fine particles at 170°C (Figure 4(b)-peak aging) to rod-shaped particles at 350°C (Figure 4(d)-over aging) indicating the possibility of the commencement of incoherency.Table 5
SDAS, grain size, porosity %, level of modification, and volume fraction of intermetallics for alloysA, D, E, F, and G.
Alloy code conditionSDAS (μm)Grain size (μm)Porosity (%)Volume fraction of intermetallics (%)AvAvAvSDAvSDA-AC∗19.32010.1350.063.080.32A-SHT23.11920.120.051.270.13D-AC18.21860.1210.074.220.35D-SHT18.61840.1220.062.760.47E-AC18.91220.110.054.230.40E-SHT19.91160.1080.042.860.43F-AC19940.110.055.020.43F-SHT21960.1110.063.740.31G-AC25.31250.120.033.570.23G-SHT21.21110.130.042.030.34∗AC : as cast.Figure 4
Development of precipitation-hardening Al2Cu phase in alloy A during 2h aging at (a) 155°C, (b) 170°C, (c) 240°C, and (d) 350°C.
(a)(b)(c)(d)


## 3.2. Tensile Properties

In this section, the tensile properties of the present alloys following aging treatment at 155°C (under aging), 190°C (peak aging), and 350°C (over aging) will be discussed. The results are presented in Figures5–7. In spite of the relatively low solution heat treatment temperature of 495ºC used for the purposes of this work, increased strength after the treatment is normally to be observed in similar 354 alloys which use a solution heat treatment temperature of 525ºC for 2-3 h period [32–38]. This selected temperature for the solution heat treatment might appear to be overconservative, but it was deemed necessary to avoid the risk of any incipient melting of the copper phases, which could have the potential for deteriorating the mechanical properties of the alloys to a great extent.Figure 5
Variation of the tensile properties as a function of aging at 155°C for times up to 100 h: (a) UTS, (b) YS, and (c) %El. Same legends as in Figure5(a).
(a)(b)(c)Figure 6
Variation of the tensile properties as a function of aging at 190°C for times up to 100 h: (a) UTS, (b) YS, and (c) % El. Same legends as in Figure5(a).
(a)(b)(c)Figure 7
Variation of the tensile properties as a function of aging at 350°C for times up to 100 h: (a) UTS, (b) YS, and (c) % El. Same legends as in Figure5(a).
(a)(b)(c)The strength values, particularly the UTS, increased by as much as 30 pct between the as-cast (AC) and the solution heat-treated (SHT) conditions. This increase in the mechanical properties is related to the changes occurring in the form of the silicon particles which become more rounded, decreasing their aspect ratio and density. Through solution heat treatment and artificial aging, the strength of alloyA was increased by ∼64% over its as-cast strength. From the tensile properties as shown in Figures 5–7, the maximum value for UTS of ∼386 MPa may be obtained with two different sets of aging conditions, i.e., 155°C/100 hrs and 190°C/2 hrs. The greatest mechanical stability for the aging treatments was observed at 155°C. Even though a significant decrease in the alloy strength is not in evidence from 2 to 100 hours at 350°C, nonetheless, the greatest mechanical deterioration will be observed after less than 2 hours of aging treatment. This previous assumption is based on a consideration of the tensile properties of the material upon solution heat treatment, namely, UTS = 305 MPa and YS = 161 MPa, implying that the over aged samples tended to lose between 15% and 40% of their original strength within the first 2 hours of being subjected to aging at 350°C.Considering the yield strength in the solution heat-treated condition, in conjunction with the artificial aging conditions used, it was possible to increase the yield strength value by more than 100 pct. The maximum yield strength was reached after 100 hours at the aging temperature of 155°C and after 2 hours at 190ºC. The yield strength behavior appears to be most stable at 190°C aging temperature, showing a variation of 50 MPa in yield strength values for aging times between 2 and 100 hours. These observations were taken into account for the selection of this specific temperature for experiments designed to investigate how the mechanical properties may be affected by long exposure times.The general behavior of alloyA with regard to ductility displays a decrease from 155°C to 190°C. A lower ductility value is observed for the condition involving 190°C for 10 hours. The greatest ductility values may be observed for aging at 350°C although none of the aging conditions reaches the higher ductility values shown in the solution heat-treated condition. This observation may be considered as evidence that the mechanical behavior displayed by alloy A is common to that of the Al-Si-Cu-Mg alloys whose strength is obtained at the expense of ductility [39, 40].The YS values for alloysA and D at aging temperatures of 155°C and 350°C represent the extreme ends of the range of aging temperatures used for these alloys. A decrease in yield strength as a result of the addition of nickel is noted at the aging temperature of 155ºC, of the order of 20 MPa. In contrast, no appreciable difference in the YS values is observed between alloys D and A at the 350°C aging temperature. This decrease in strength through the addition of nickel is related to the interaction between copper and nickel to form precipitates of Al3CuNi in the microstructure. The Al3CuNi phase was detected in alloys containing nickel, namely, alloys D, F, and G. Since copper, as well as magnesium, determines the precipitation strengthening of the Al-Si-Cu-Mg alloys, the available copper would form Al2Cu precipitates through the solution heat treatment/quenching/aging process, thereby strengthening the cast alloys. The formation of Al-Cu-Ni precipitates would thus subtract part of the copper content available for strengthening; this is the principal reason for which a slight decrease in YS is observed with the addition of 0.4 wt.% nickel.With regard to the base alloyA and alloy D, the maximum strength is reached at low aging temperatures ranging from 155ºC to 190ºC, showing a decrease in ductility with the increase in strength; it should be noted that this is the expected behavior for Al-Si-Cu-Mg alloys [41–45]. It will also be observed that, at low aging temperatures, the effect of intermetallics in the reduction of ductility is more pronounced than it is at high aging temperatures of 350 ºC. The increase in ductility at an aging temperature of 190 ºC, or more, becomes evident since the alloy softens and the effect of the intermetallics on ductility seems attenuated.From Figure6, it is observed that alloy E, containing 0.4 wt.% Zr, does not show any significant change in YS, when compared with the base alloy A. The decrease in the ductility of alloy E with respect to alloy A may be assimilated by observing Table 5 where the level of insoluble intermetallics increases from 1.27 to 2.86 %, between the base alloy and alloy E, respectively. This behavior may also be seen for the long-exposure-time at the aging temperature of 190ºC. In general terms, there is no appreciable increase in strength (UTS and YS) to be observed in alloy E containing 0.4 wt.% Zr although there is, however, a visible decrease in ductility which is greater than 20% for nearly all of the tests performed.To understand the latter observation, it is necessary to consider three characteristics observed from the microstructural analysis of alloyE. Firstly, the amount of insoluble intermetallic phases increased from 1.27% in the base alloy to 2. 86% in alloy E. Secondly, the grain size was reduced from 200 μm in alloy A to 120 μm in alloy E. Both arguments could indicate that Zr available for strengthening was used to form large intermetallics in the form of (Al,Si)3(Zr,Ti) phases, as well as to reduce the grain size. Finally, a third aspect which should be taken into consideration is that during the casting procedures, the temperature of the melt was maintained at 780°C, implying that Zr added in the form of Al-20% Zr master alloy may not be capable of dissolving completely; this is due to the fact that, with 0.4 wt.% Zr, the Al-Zr phase diagram indicates the liquidus temperature at 780°C [41].As a result of examining the negative effects of nickel on the strength of alloyE, it becomes relevant to discuss why the addition of nickel did not reduce the yield strength in alloy F which contains 0.4 wt.% Ni plus 0.4 wt.% Zr. With regard to intermetallics, our earlier studies [46–48] show the precipitation of (Al,Si)3(Zr,Ti) phase with the Ni-rich phase Al9NiFe adhering to it. The Al9NiFe phase was observed in the microstructure of tensile samples, as well as phases which appear similar to the Ni-rich phase were observed to adhere to the Zr-rich phases in the microstructures of the tensile samples of alloy F as demonstrated in Figure 3. These observations may suggest that the copper did not significantly interact with the nickel to form Al3CuNi phases, and as a result, the copper was free to strengthen alloy F, in the same way as the copper was able to strengthen alloy A.Ductility seems to be affected principally at low aging temperatures of 155°C through 190°C. This decrease in ductility may be interpreted by noting that alloyF had the greatest amount of insoluble phases as listed in Table 5; these would then act as stress concentrators during flow stress, thereby limiting the ductility in alloy F [49, 50].AlloyG, containing 0.2 wt.% Zr and 0.2 wt.% Ni, displays greater strength than alloy E, which contains 0.4 wt.% Zr, or greater even than alloy F, which contains 0.4 wt.%. Zr plus 0.4 wt.% Ni. According to the Al-Zr phase diagram [41], with 0.2 wt.% Zr content, the liquidus temperature is approximately 720°C. Taking into consideration that, during casting procedures, the temperature of the melt was maintained at 780°C, it is fairly likely that the added Zr dissolved completely which is essential for strengthening the alloys by means of Zr-rich precipitates [51, 52]. As may be seen from Figure 6, alloy G is the only alloy which shows an increase in strength in comparison with alloy A. In fact, even at the aging temperature of 350°C, the strength values are observed to be greater than those obtained with the base alloy.Figure8 depicts an example of the tensile properties obtained for the five studied alloys at 170°C relative to the values obtained for the A base alloy in the as-cast condition, i.e., subtracting the values obtained for the base alloy A in each case and plotted as △P values on the Y-axis (P = property = YS, UTS or %El), with the X-axis representing the aging time (see Supplementary Materials for 155°C and 350°C aging temperatures). This approach helps to better visualize the effects and interactions of the various additions used and the different heat treatment conditions, so as to arrive at a better understanding of the effects of Ni and Zr on the tensile properties of 354 casting alloys. As may be seen from Figure 8, additions of these elements improve the alloy performance, i.e., the tensile properties at 170°C. Figure 8(a) shows that alloy A exhibits the best contribution to UTS at all aging times. After 10 h aging time, the comparative UTS (or △PUTS) values start to decrease as the aging time increases, indicating that these elements provide little or no improvement in mechanical properties. On the other hand, the comparative yield strength (or △PYS) values of all alloys are improved, as shown in Figure 8(b) and are stabilized after 8 h aging time, while Figure 8(c) shows that the △P%El or percent elongation to fracture values for all alloys is much lower than that of the A base alloy, which is to be expected, in light of the results shown in Figure 8(a).Figure 8
Comparison of tensile properties ofA, D, E, F, G alloys relative to those of A alloy in the as-cast condition: (a) UTS, (b) YS, and (c) % El at 170°C. Same legends as in Figure 8(a).
(a)(b)(c)Figure9 compiles the strength-ductility values obtained for all the 354 alloys investigated. These values fall into regions corresponding to peak aging, under aging, and over aging. Using regression analysis, the UTS versus %El relationships may be expressed as follows:UTS = –38.26x + 408.1 with a fit of R2 = 0.82 (alloy A)UTS = –45.0x + 373.0 with a fit of R2 = 0.87 (alloy F)Figure 9
The UTS vs. % El relationship for (a) alloyA and (b) alloy F.
(a)(b)It is evident that the slope of alloyA is higher than that for alloy F. Also, the UTS levels are higher than those for alloy F due to the higher volume fraction of intermetallics in the latter as noted from Table 5.


## 3.3. Q Charts

This section will be limited to alloyA. The Cáceres method which is used here for constructing the quality index charts involves the use of a single value of the strength coefficient (K) for all the conditions appearing in the charts. In order to obtain the values of n and K described in equations (5) and (6), a series of stress-strain curves were used as shown in Figure 10, whereas their aging conditions are listed in Table 6. Based on the analysis of these curves, the values of n = 0.193 and K = 530 were obtained. The reason for taking only 15 conditions out of a total of 80 was principally to avoid an overcrowded diagram. The curves were divided into three groups. The criterion applied for using these 15 conditions (including the as-cast and SHT conditions) was to take into consideration the extreme values, particularly those regarding the % elongation and yield strength values.Figure 10
A series of stress-strain diagrams used to calculate the n and K values (see Table6 for details).
(a)(b)(c)Table 6
Aging condition for curves shown in Figure10(a)–10(c).
Figure #Curve #Aging condition10(a)1As cast2SHT10(b)12 h/155°C22 h/170°C310 h/155°C4100 h/155°C52 h/190°C612 h/170°C7100 h/155°C10(c)12 h/350°C2100 h/240°C310 h/240°C42 h/240°C5100 h/190°C610 h/190°CFigure11(a) shows that the plastic strain and the quality index (Q) both display a significant increase upon solution heat treatment in spite of the fact that the relative ductility (q) attains a value of 0.31 implying that alloy A in its solution-treated condition has reached 31% of its maximum quality index value (Q). When the ductility increases sharply from the as-cast to the solution heat-treated condition, such changes can be related to the spheroidization of silicon particles and to the uniformity of the microstructure in the solution heat-treated condition as illustrated in Figure 2(d). Figure 11(b) demonstrates the Q-chart for alloy A aged at 155°C for times up to 100 h. As can be seen, most of the points are falling in a narrow range of Q (416-366 MPa). The maximum value of Q was achieved after aging for 2 h with minimum UTS value of 308 MPa, and a minimum Q value was obtained with a maximum UTS value of about 387 MPa after 100 h. Apparently, alloy ductility plays a major role in determining the Q level.Figure 11
Quality index charts of alloyA at different aging temperatures: (a) as-cast and SHT, (b) 155°C, (c) 170°C,(d) 190°C, and (e) 350°C.
(a)(b)(c)(d)(e)TheQ charts following aging at 170°C and 190°C are shown in Figures 11(c) and 11(d), respectively. Apparently, there is more than one C-curve due to multiple reactions that took place during the aging process. From Figure 11(c), the values of Q decreased from 421 MPa (2 h) to a value of about 375 MPa (100 h). The stress-strain curve corresponding to minimum Q value is displayed in Figure 10(b)-curve #6. With aging at the 190°C/2 hr condition, GP zones and θ′ precipitates develop such that they are coherent with the matrix; YS is thus increased because of the strain fields associated with these precipitates, which are sheared by the dislocations and subsequently lead to a low strain-hardening rate as shown in Figure 10(b)-curve #5.Aging at 350°C resulted in a significant improvement in the alloy ductility. The chart in Figure11(e) could be divided into two curves (solid line) or simple straight line 2h-100 h (broken line). During the course of aging, the values of UTS level were reduced from 221 MPa (2 h) to 200 MPa (100 h) with significant increase in % elongation from about 4.5% (2 h) (curve #1 in Figure 10(c)) to 6.5% (100 h) resulting in Q values of 350-MPa (2 h)–345 MPa (100 h). In other words, practically there is no change in the Q levels for the entire aging time (±5 MPa).It may be deduced from Figure4 that the θ-CuAl2 precipitates are formed at temperatures of 190°C and higher, i.e., 350°C [53–55]. These particles are incoherent with the matrix and cannot be cut by dislocations, thereby promoting lower strength values and a high strain-hardening rate resulting from the accumulation of Orowan loops around the copper phase particles. As the strain is increased, however, the gradual development of primary shear loops generates intense stress fields around the strengthening precipitates, which are in themselves limited [45, 56, 57] by the activation of a secondary dislocation cross-slip process, thus reducing the strain-hardening ability of the material.


## 4. Conclusions

From an analysis of the results obtained, the following conclusions may be formulated:(1)
Zirconium reacts only with Ni, Ti, Si, and Al in the alloys which are examined to form two phases: (Al,Si)3(Zr,Ni,Fe) and (Al,Si)3(Zr,Ti).(2)
There is no negative effect in the refinement of grain size with the addition of Zr to the alloys containing Ti. Addition of 0.2% Zr resulted in about 40% reduction in the alloy grain size.(3)
Addition of Zr or Zr + Ni has a minor effect on the alloy strength at aging temperatures as high as 240°C–350°C (over aging).(4)
For aging at 350ºC for times up to 100 h, the YS values for alloyG (with 0.2 wt.% Zr and 0.2 wt.% Ni) increased by about 15% with respect to the rest of the alloys studied.(5)
Addition of Zr beyond 0.2%, i.e., 0.4% is not recommended due to the high liquidus temperature (∼780°C) of the Al-Zr binary phase diagram.(6)
Nickel reacts with Cu resulting in a decrease in the alloy tensile properties by about 10% (∼30 MPa).(7)
Quality index charts could represent more than a single precipitation reaction, as inferred from the presence of multiple C curves. In such category of alloys, ductility plays a decisive role in determining theQ levels.



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*Source: 5551633-2021-04-07.xml*