Effects of high Zn content on the microstructure and mechanical properties of Al–Zn–Cu gravity-cast alloys

Author’s Accepted Manuscript Effects of high Zn content on the microstructure and mechanical properties of Al–Zn–Cu gravitycast alloys Sang-Soo Shin, Kyung-Mook Lim, Ik-Min Park www.elsevier.com/locate/msea

PII: DOI: Reference:

S0921-5093(16)31086-3 http://dx.doi.org/10.1016/j.msea.2016.09.022 MSA34107

To appear in: Materials Science & Engineering A Received date: 28 June 2016 Revised date: 3 September 2016 Accepted date: 3 September 2016 Cite this article as: Sang-Soo Shin, Kyung-Mook Lim and Ik-Min Park, Effects of high Zn content on the microstructure and mechanical properties of Al–Zn–Cu gravity-cast alloys, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2016.09.022 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

Effects of high Zn content on the microstructure and mechanical properties of Al–Zn–Cu gravity-cast alloys Sang-Soo Shina, Kyung-Mook Limb*, Ik-Min Parka* a

Department of Materials Science and Engineering, Pusan National University, Busan, 46241, Korea

b

Korea Institute of Rare Metal, Korea Institute of Industrial Technology, Incheon, 21999, Korea

[email protected] [email protected] *

Corresponding author. Doctor Kyoung-Mook Lim. Tel.: +82 32 458 5113; fax: +82 32

458 5120 *

Corresponding author. Professor Ik-Min Pakr. Tel.: +82 51 510 2393; fax: +82 51 514

4457 Abstract In this paper, the effects of Zn on the microstructural evolution, hardness, and various characteristics of Al–20~45Zn–3Cu alloys were investigated. High-strength Al–Zn–Cu alloys (> 450 MPa) were fabricated by gravity casting without melt modification or heat treatment (i.e. T4~T6). Al-based alloys containing more than 20% Zn were designed for the gravity-casting process. In terms of the microstructure, as the amount of Zn addition in the alloys increased, the α-phase size decreased and the α + η non-equilibrium solidification-phase fraction increased. A complex network of eutectoid  + η, supersaturated η, and Cu-related intermetallic particles formed in the grain boundary regions. In addition, increasing the Zn content improved the mechanical properties of the alloys but reduced their damping capacity and toughness. The fractographic examination of the fracture surfaces indicated that the Al-Zn-Cu alloys with high Zn addition had fewer ductile dimple surfaces and more brittle cleavage surfaces. Keywords: Al–Zn–Cu alloys; Gravity casting; High strength; Damping capacity; Impact test

1. Introduction Recently, the number of aluminum automobile components and electronic devices produced by the die-casting method has increased and now accounts for approximately 35% of the total production of aluminum parts. This increasing number of Al-cast alloys has been used for automotive parts, such as transmission cases, converter housings, and cylinder blocks. Furthermore, aluminum part production is expected to increase further due to the high-pressure die-casting (HPDC) process, as the HPDC method is more suitable for mass production due to its higher productive efficiency, and its capacity to produce near net shapes that are thin-walled with complex aluminum components,

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displaying good dimensional accuracy with a good surface finish, and excellent mechanical properties [1, 2]. Moreover, the commercial ADC12 (Al–Si–Cu) alloy is a cast aluminum alloy that is widely used for transmission cases and power train parts due to its high castability, low density, high productivity, and low shrinkage rate [3, 4]. However, ADC12 alloys have not always been suitable for automotive parts, as these alloys do not have superior characteristics in terms of their mechanical properties due to the brittle nature of the Si component and their relatively high melting points when compared with Zn-based die-cast alloys. Meanwhile, various studies have been undertaken to produce casting alloys with excellent mechanical properties by using Al-wrought alloys, thus replacing the relatively low-strength Al–Si casting alloys. Various Al–Zn-based alloys have been examined in an attempt to overcome the issues inherent in the brittle Si component and the high melting points of Al– Si-based alloys [5–10]. Compared to other cast alloys, Al-wrought alloys are characterized by high melting temperatures, low fluidity, and long freezing regions, which tend to create various defects such as hot tears, and result in shrinkage or porosity issues. However, the Al–Zn-based alloys developed by Lim and Shin et al. [9, 10] not only exhibited very good mechanical properties and fluidity, but also resulted in a lower melting point. Thus, the addition of large amounts of Zn can effectively lower the melting point of the resulting alloys and improve their fluidity [9]. Furthermore, the developed alloys exhibited an ultimate tensile strength of 470 MPa and an elongation of 3.5% without melt modification and post-casting heat treatment. The developed alloys exhibit improved mechanical properties with the addition of a large amount of Zn, which also leads to the presence of an α + η phase or η-rich phase in the grain boundary region due to the monotectic reaction (L1 = α + L2) [11]. Consequently, α-Al is enclosed in the Zn-rich phase to form a network structure. In particular, the grain boundary area is mostly composed of an α + η phase lamellar structure. This lamellar structure is greatly affected by the cooling rate, and the microstructure and mechanical properties of the Al–Zn– Cu alloy can also be changed depending on the manufacturing method or conditions. In this study, we use gravity-cast Al-based alloys containing more than 20Zn. In order to design conventional cast alloys with superior mechanical properties, Al-based alloys with various Zn contents were prepared by gravity casting, and their microstructure and mechanical properties were analyzed. Furthermore, we conducted thermal analysis of the alloys. Based on the results, we discuss the findings with respect to the flow behavior of the Al–Zn–Cu alloys. Additionally, we estimate the damping capacity and impact characteristics of the Al–Zn–Cu alloys as a function of their high Zn content. Therefore, in this paper, we investigate the effects of high Zn content on the microstructure and certain characteristic variations in the Al–Zn–Cu alloys produced by the gravity-casting method.

2. Experimental procedure

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2. 1. Materials and sample preparation The chemical compositions of the Al–Zn–Cu alloys studied in this work are shown in Table 1. The alloys, with a nominal composition of Al–20~45Zn–3Cu, were prepared from high purity Al and Zn, Al–10Cu, Al–10Fe, and Al–8Si master alloys in an electrical resistance furnace, and then cast into a metallic mold at a pouring temperature of 700~720 °C with a mold temperature of 180~200 °C. Before pouring, the molten alloys were stabilized by holding the temperature at 720 °C for 10 min to ensure homogeneity and dissolution of the present intermetallic compounds. The alloys were poured into a pre-heated permanent mold with the dimensions of w 75×t 60×h 240 mm. Table 2 shows the chemical compositions of the prepared ingots analyzed by inductively coupled plasma–atomic emission spectroscopy (Icap 6500 Duo, Thermo-Fisher, Switzerland). 2. 2. Microstructural analysis, fractographic analysis, and thermal analysis of the Al–Zn–Cu alloys The microstructure of the samples was observed using optical microscopy (Axio Observer.A1m, Carl Zeiss, Germany) and a field-emission scanning electron microscope (FE-SEM; Sirion 400, FEI) equipped with an energy-dispersive spectrometer (EDS). A SEM (FE-SEM, HITACH II-SU2080, Japan) was also used to carry out the fracture surface observations of all the samples. The energy-dispersive X-ray spectrum was used in determining the chemical elements and constituents of each of the phases. In order to investigate the microstructures of the Al–Zn–Cu alloys, metallographic samples were taken from the grip areas or fracture areas of the tensile samples using water cooled with a diamond disk wheel. The hot mounted samples were ground, polished, and etched in a solution composed of a 3% hydrofluoric acid (HF) solution. The α-phase grain size and the volume fraction of the non-equilibrium solidification phase were determined by electron backscattering diffraction (EBSD; S-4300, Hitachi) with a scanning step size of 0.7 μm. The EBSD data were processed using orientation imaging microscopy analysis software (TSL-OIM, EDAX); the data points with a confidence index below 0.1 were removed from the EBSD data. The melting temperatures of the Al–20~45Zn–3Cu alloys were measured by inserting a K-type thermocouple (Φ 1.0 mm) into the permanent mold. The thermal properties, such as the phase-transformation temperature, liquidus temperature, solidus temperature, and the latent heat of the Al–20~45Zn–3Cu alloys were measured repeatedly by differential scanning calorimetry (DSC 823e, Mettler Toledo, Switzerland) and differential thermal analysis (SDTA 851e, Mettler Toledo, Switzerland), during which the samples were examined at a heating rate of 20 °C/min. 2. 3. Mechanical properties of the Al–Zn–Cu alloys

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For each composition, 15 tensile test bars were prepared by gravity casting (ASTM 370-05), and their tensile properties were measured using an Instron-5989 tensile test machine with a crosshead speed of 1 mm/min, which is equivalent to a strain rate of 2.78

10−4 s−1. Non-contacting advanced

video extensometers (AVEs) of the tester were utilized to measure specimen elongations. The hardness of the Al–Zn–Cu alloys was determined using a Rockwell Hardness Tester (Wilson, RB 2000, USA) with a B-scale indenter (HRB) loading time of 1.0 sec. Impact test bars were produced after having poured the degassed molten metal into a preheated permanent metallic mold (SKD 61) with the dimensions of w 75 × t 60 × h 240 mm. These test bars were cut from the casting (ingot) and then machined to the required ASTM: B 108 specifications, as shown in Fig. 11. Impact tests were performed on un-notched samples. A computer-aided instrumented IT 542 E-Horizon version 10.0.9.5 Universal Impact Testing Machine (Tinius Olsen, USA) was used during testing. The average values of the energies obtained from the samples were tested for each alloy at room temperature (approximately 20 ℃). The damping capacity is obtained by the measurement of internal friction at room temperature in an air atmosphere using a free-decay method of flexural resonant vibration with both free ends (strain amplitude: about 2

10−6) at room temperature (approximately 19 ℃) [12, 13]. The

dimensions of the specimens for internal friction measurements were 1 mm small piece of iron foil (8 mm

4 mm

10 mm

120 mm. A

0.2 mm) was attached to both ends of the specimen in order

to control the mechanical vibration efficiently due to electromagnetic force and to sensitively monitor the vibration from the induced current.

3. Results and discussion 3.1. Thermal properties of the Al–Zn–Cu alloys In order to measure the thermal properties (e.g. the melting points and phase transformations of the Al–Zn–Cu alloys), a K-type thermocouple was inserted into the molten metal alloy in the permanent mold. Fig. 1. shows the cooling curves obtained from the Al–Zn–Cu alloys solidified at a constant cooling rate of 2.5 ℃/sec. The melting point of the alloys decreased with increasing Zn content. The undercooling curve was not observed in Fig. 1(a). However, a phase transformation was observed at 200~300 ℃. It is evident that the eutectoid reaction for Al–Zn (α + η) occurs in the Al– 40Zn–3Cu samples, as shown Fig. 1(b). This SEM image shows the distribution of typical precipitates via the eutectoid reaction β  α + η at 260 °C. Therefore, the grain boundary region of the Al–Zn–Cu alloys is predicted to exhibit a very complex microstructure, such as an α + η phase, an η-rich phase, and Cu-related intermetallic compounds. DSC and DTA analyses were carried out to investigate the more detailed thermal properties of the alloys and the results are shown in Table 3.

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Table 3 shows the solidus temperature, liquidus temperature, and latent heat of the alloys, which dramatically decrease with increasing Zn content. As a result, the solidification interval range is very wide (from 130 to 165 ℃). This characteristic suggests that these alloys are relatively easy to apply a semi-solid process to. Meanwhile, the fluidity is generally inversely proportional to the solidification range [14, 15]. However, a recent study by Han [16] found that under die-casting conditions, the fluidity length of an aluminum alloy increases with an increasing solidification interval and a decreasing solidus temperature at a given superheat, which is opposite to the fluidity of aluminum alloys in gravity casting. Considering these aspects, the HPDC Al–Zn–Cu alloys are expected to exhibit a very high flow length. The variable parameters for the fluidity of Al alloys are the latent heat, surface tension of the flow head, solidification behavior, and the pouring temperature [17]. Assuming that there are no the other variables related to fluidity in the Al–Zn–Cu–Fe–Si alloy (a multicomponent alloy), the fluidity of the alloy is dependent on the solidification interval range, and the solidification behavior of the flow head [18]. Meanwhile, the developed Al–Zn–Cu alloys show excellent fluidity in the gravity condition compared with ADC12 [9]. From this experimental result, the fluidity of the alloy system is presumably related to the solidification behavior of the flow head. The developed Al–Zn–Cu alloys, despite having a high Zn content, do not form specific intermetallic compounds [19]. Thus, when the developed alloys started to solidify in the flow head, the primary α phase initially crystallized from the melt via a monotectic reaction (L1 = α + L2). Here, L2 remained in the liquid phase, which improves the fluidity of the molten metal. In addition, the Cu-related particles in the Al–Zn–Cu alloys are present in a relatively spherical form, unlike the flake-type β-Al5FeSi phase, and thus we determined that this did not inhibit the fluidity of the flow head in the alloys. As shown in table 3, the density of the alloys exhibits a high value due to the addition of a large amount of Zn. Regarding this finding, the developed Al–Zn–Cu alloys are not light in weight, thus leading to certain restrictions in their use as structural materials. Although the produced Al–Zn– Cu alloys do not belong in the lightweight alloy category due to the large amounts of Zn compared with typical commercial alloys produced by die-casting (e.g. the ADC12 alloy produced by the gravity-casting technique), these alloys do have a high specific strength. 3.2. Microstructure of the Al–Zn–Cu alloys The microstructures of Al–Zn–Cu alloys with various high Zn contents are shown in Fig. 2. From a previous study [9], the microstructure of Al–Zn–Cu alloys consists of an α-Al phase, α + η phase, and an η-rich phase, with Al4Cu3Zn and Cu-related intermetallic compounds. The Ɵ phase can act as a reinforcing phase in the alloys. The Ɵ phase is not visible in Fig. 2. However, the microstructure was refined by further increasing the Zn content. By increasing the Zn content, the microstructure changed from that of a large to small dendrite morphology. The Zn-rich grain

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boundary region increased with increasing Zn content. The microstructure of the alloys produced by gravity casting exhibited a much coarser morphology compared to those alloys produced by HPDC [9], as the high-pressure die casting method involves a rapid cooling rate and the application of static pressure to the sample [20]. Quantitative image analysis was performed using the EBSD technique to study the microstructural evolution in the alloys with different Zn additions. The results for the measured average grain size and the fraction of the grain boundary regions are presented in Fig. 3. It can be seen that with increasing Zn content, the average grain size of the α-phase decreases from 321.16 to 101.41 μm. Fig. 4 also demonstrates that in the presence of Zn, the fraction of the grain boundary regions increases from 10 to 24%. In addition, the uniform distribution of the α-Al phase with increasing Zn content can be seen. The grain refinement of the developed Al–Zn–Cu alloys as a function of Zn addition is associated with the melt superheat of the alloys. In general, melt undercooling increases with increasing melt superheat, which results in an increase in the nucleation rate and a decrease in grain size [21]. The reduced solidus temperature of the Al–Zn–Cu alloys with high Zn content induced the increase in the melt superheat, which results in the large undercooling and grain refinement. In order to investigate the effect of high Zn addition on the intergranular fraction of the Al– Zn–Cu alloys with an α-Al phase distribution, the microstructure was characterized by SEM, as shown in Fig. 4. Fig. 4(a) shows the SEM micrograph of the α-Al phase with a Zn-rich grain boundary in the Al–20Zn–3Cu alloy, which reveals a coarse α-Al phase with intermittent distribution of the Fe–Si and Cu-related phases. It can be clearly observed that the grain size significantly decreases and the alloy’s intergranular fraction increases with the addition of Zn. However, the morphology of the Fe–Si and Cu-related phases is still flake-like and relatively spherical, respectively, as shown in Fig. 4(b) – (d). Besides, the Fe–Si-related phase is present in a random distribution near the grain and the Cu-related phase is located inside the Zn-rich grain boundary. The intergranular Znrich non-equilibrium solidification region is expected to be composed mostly of a fine α + η lamellar structure. In addition, this region is found to be included in the AlCu phase, the other Cu-related phase, and the Zn phase [9]. Fig. 5 shows the magnified SEM micrographs of the intergranular region for the Al–45Zn– 3Cu gravity-cast and HPDC condition alloys. This magnified image highlights the typically ultra-fine α + η phase lamellar structure of the alloys. Obviously, the lamellar structure of the α + η phase of the HPDC Al–45Zn–3Cu alloys was finer than that of the gravity-cast alloys (Fig. 5(a) and 5(b)). For this reason, the gravity-cast lamellar spacing is wider due to the relatively slower cooling rate than that of the die-casting method. In particular, static pressure is applied with the die-casting method after

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injecting the Al molten metal into the cavity. Thus, the grain size can be refined and the porosity removed, which results in a reduction in the lamellar spacing of the alloy. Moreover, the Cu-related phase of the alloys produced by the gravity-casting process was non-uniformly distributed compared with the alloys produced with the HPDC method, and the Cucontaining intermetallic compounds were coarser than those in the HPDC samples, as shown in Fig. 5(a) and 5(c). This is attributed to the difference in the cooling rate, which is relatively slow in gravity casting. Consequently, there is enough time for the Cu element’s solute redistribution process. The Cu and Zn present in the Al–Zn–Cu alloys had an equilibrium distribution coefficient of k < 1 (k = Cs/Cl, where Cs and Cl are the solute concentrations of the solid and liquid, respectively, at the interface), and the elemental Cu and Zn remained in the Zn-rich grain boundary or in the non-equilibrium solidification-phase region. Thus, it is presumed that the Cu-related phase was formed in the Zn-rich grain boundary regions under the gravity-casting condition. Another result of the slow cooling rate is the Al–Zn (α + η) precipitate forming via the eutectoid reaction, as shown in Fig. 1(b). This precipitate was formed via the eutectoid reaction β  α + η at 260 °C. Therefore, the grain boundary of Al–Zn–Cu alloys was preferentially created by the formation of the β + η phase. The final microstructure was completed by the eutectoid reaction β  α + η at 260 °C. On the other hand, the high cooling rate of HPDC obstructed the eutectoid reaction, and the angular β and η phases formed by the eutectic reaction remained at room temperature. A more detailed phase and composition analysis was carried out using the EDS technique for the Al–45Zn–3Cu alloys and the results are shown in Fig. 6. EDS 1 shows the results for the α-Al phase, which employed a large amount of Zn and a small amount of Cu. These Zn and Cu elements are introduced into the Al matrix in the form of a solid solution, as the melt superheat of the developed alloys increased with high Zn addition. As a result, with the increase in supercooling, the solidification rate increased, which produced a supersaturated solid solution [22]. Therefore, the developed alloys are presumed to improve their strength due to solid-solution strengthening. Fe and Si are added to Al–Zn–Cu alloys to improve fluidity and prevent mold burning. They are combined with each other to form coarse phase morphologies, as shown in EDS 2. These Fe–Sirelated phases have a flake-like or polygonal shape, as shown in Fig. 4(b). The sharp edges of the Fe– Si-related polygon phases act as stress concentration sites, which assist crack initiation, and thus decrease the ductility of the castings. EDS 4, EDS 6, and EDS 7 reveal the Cu-related phases, such as Zn5Cu, AlCu, and Al4Cu3Zn (the T′ phase) [9]. On the other hand, the EDS 5 result shows the lamellar structure of the α + η phases, and in this image, the morphology of the microstructure can be seen, as shown in Figure 5(d). Moreover, this microstructure contained a small amount of Cu, and so the developed Al–Zn–Cu alloys contained a large amount of Zn in the Al matrix. From these results, the developed alloys can be strengthened by solid-solution strengthening, as mentioned previously. Additionally, the θ phase and the Cu intermetallic compounds were evenly distributed in the alloys in

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a transgranular and intergranular manner. Consequently, the developed Al–Zn–Cu alloys can exhibit excellent mechanical properties. 3.3. Mechanical proprieties of the Al–Zn–Cu alloys As the amount of Zn increased, the fraction of Zn-rich grain boundary regions in the Al–Zn–Cu alloys also significantly increased, and as a result, the mechanical properties of the alloys changed. The results of the hardness tests are shown in Fig. 7. It can be seen that the Rockwell hardness of the alloys linearly increased with increasing Zn content. According to the microstructural evolution of the α + η phase with Zn-rich grain boundary regions, the hardness should enhance with a rise in Zn content. Furthermore, the hardness of the α + η phase, θ phase, and the Cu-related phase is higher than for the α-Al phase in the Al–Zn–Cu alloys [9]. Therefore, it is evident that the presence of the α + η phase and θ phase improves the hardness of the alloys. From these results, we would expect the strength of the developed Al–Zn–Cu alloys to improve proportionally with increasing Zn content. We next measured the tensile properties of the alloys to evaluate the influence of the Zn content. Fig. 8 presents the results obtained from the tensile test, and shows the changes in the yield, tensile strength, and fracture strain measured as a function of the Zn content. The dog-bone shape of the tensile specimens used in the test (ASTM a370-05) is shown in the insert in Fig. 8. The average yield, tensile strength, and fracture strain values of the alloys with various Zn additions are shown in Table 4. It is evident that high Zn additions improve both the yield strength and ultimate tensile strength (UTS) values of the Al–Zn–Cu alloys. The yield strength and UTS increased from 203 to 402 and from 315 to 455, respectively. However, the elongation values decreased from 4.5 to 3.1%. The developed Al–Zn–Cu alloys should therefore have very good strength caused by containing the reinforcing phase (e.g. the θ phase or Cu-related phases). Moreover, this can be attributed to the effect of the high Zn addition refining the α-Al phase and this being surrounded by a fine lamellar α + η structure and thus extending the grain boundary regions, but the increase in the brittle hard phase (e.g. the θ phase or Cu-related phases) resulted in a decrease in elongation. As mentioned previously, the presence of the fine α-Al phase surrounded by a fine lamellar α + η structure with reinforcing phases can restrain the creation and movement of dislocations during shear stress exertion [23], thereby enhancing the tensile strength of the alloy. Furthermore, the solid-solution hardening effect of Cu may have been one of the reasons for the improvement in the tensile strength of the Al–Zn–Cu alloys. Therefore, Al–Zn–Cu alloys exhibit work-hardening in uniaxial tensile stress–strain curves, as shown in Fig. 8. Such excellent mechanical properties have not been reported for other aluminum-cast alloy systems in previous studies.

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Fig. 9 shows the fracture surfaces after the tensile testing of the various Al–Zn–Cu alloys. It is evident from Fig. 9 that the fracture features exhibited by the samples with low Zn (20 wt.%) have more ductile behavior compared to the relatively brittle mode of the high Zn-containing alloys. A number of cleavage facets suggest an increase in the degree of the brittle mode of fracture with increasing Zn content. In the more than 30% Zn alloy, large particles of the Fe–Si-related phase and non-equilibrium solidification phase caused high levels of stress concentration, and consequently, a reduction in tensile ductility. It is obvious that the number of fine dimples decreases with the addition of Zn, which is consistent with the tensile test results, as shown in Table 4. In addition, the increasing fraction of the Zn-rich phase can promote the brittle fracture mode of the Al-Zn based alloy, as can be seen in Fig. 9. Furthermore, the developed Al-Zn-based alloy has a long solidification range, which can lead to casting defects such as microsegregation, porosity and shrinkage. As a result, the initial crack is initiated in the pores or shrinkage defects, as shown in Fig. 9 (a) and (b). Then, the cracks are expected to propagate along the Zn-rich grain boundaries or Cu-, Fe-Si related phase. In order to observe more details of the secondary crack propagation path, the side fracture of the tensile testing specimen was observed. Fig. 10(a) and 10(b) is the fractograph recorded from the regions near the fracture of the tensile testing sample, showing that the Al–45Zn–3Cu alloys contain α-Al dendrites, an α + η phase, a Zn-rich phase, and Cu-related intermetallic compounds. The specimen mainly fractured across the α + η phase, and Cu-containing intermetallic compounds of the grain boundary regions. In addition, the cracks were formed along the grain boundary regions. The cracks for the Cu-related intermetallic compounds are observed at the fracture surface, and these cracks lead to the formation of secondary cracks, as shown Fig. 10(a) and 10(b). Moreover, Fig. 10 shows the micrograph recorded from the region ∼1.0 mm away from the fracture, the Cu-related intermetallic compounds, and the α + η phase with an intercrystalline fracture feature with its crack direction normal to the tensile direction. Moreover, the intercrystalline cracks of the Fe–Si containing intermetallic compounds can be observed at the fracture surface, as shown in Fig. 10(c). From the Fig. 10 (c) EDS results, oxygen component around secondary phase of Cu- or Fe-Si related phase was not detected. Therefore, the main cause of fracture is determined as porosity and shrinkages defect rather than double oxide film. It is evident from the various studies that if there is no defect such as pore and shrinkage defects, the reinforcing phase acts as originating of a crack in the Al-based alloys. [32-47]. Because of the reinforcing or secondary phase has a high Young's modulus as compared with the Al matrix. Thus, when the tensile load is applied, reinforcing phase act as plausible sites for stress concentration which can lead to premature failure. In conclusion, the primary crack was formed by pore and shrinkage defects in the developed Al-Zn alloy and the secondary crack was formed in the Cu particles and Fe– Si containing intermetallic compounds. The cracks were propagated along the Zn-rich grain boundary

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regions, at the same time, secondary phases are the weak locations in the matrix and are inclined to crack easily when the tensile load is applied.

3.4. Impact properties of the Al–Zn–Cu alloys The developed alloys exhibit superior mechanical properties, as shown in Fig. 8. Therefore, the impact toughness was also expected to increase with increasing Zn content. Fig. 11 presents the Charpy absorbed energy values as a function of Zn content for Al–Zn–Cu alloys in the as-cast condition. The unnotched specimen used for the impact test is shown in the insert in Fig. 11. Unlike the predictions, Charpy absorbed energy decreased with high Zn additions, as the non-equilibrium solidification region or grain boundary area increased due to Zn additions. The grain boundary regions contained a reinforcing phase (e.g. Cu-related intermetallic compounds and an η-Zn phase), which increased due to Zn addition. However, the impact absorption energy of the alloys represents a value of approximately 10~14 J across the whole sample. Impact strength is a more dominant factor with small variations in the microstructure than is elongation of the alloys [24]. The morphology of the grain boundary regions is extended by the addition of Zn for Al–Zn–Cu alloys. On the other hand, the addition of lower levels of Zn to the alloys appears to have a favorable effect on the impact strength. Meanwhile, the impact strength of the Al–Zn–Cu alloys was reduced by Fe–Si and Cu-related intermetallic compounds with increasing grain boundary regions due to high Zn additions, as shown in Fig. 10(c). Fe and Si are added to Al–Zn–Cu alloys in order to improve castability and to prevent metallic mold burning [25, 26]. These Fe and Si elements are purposefully added to decrease the toughness of the alloy, as mentioned previously. Therefore, in order to improve the elongation and impact toughness of the alloys, it is considered best not to add Fe and Si.

3.5. Damping characteristics of the Al–Zn–Cu alloys The developed Al–Zn-based alloys can be used as damping material because of their fine α + η phase lamellar structure with the η-phase having an HCP lattice structure [27, 28]. The flake graphite cast iron and Zn–Al alloys exhibited a dual phase, which was composed of a hard and soft phase, and resulted in these alloys exhibiting the damping effect due to the viscoelastic flow at the interface of each phase one. In addition, many previous studies on Zn–Al and Zn–Al-based alloys have reported excellent damping properties due to the fine lamellar structure of the α-Al phase and ηZn phase [29, 30]. Therefore, the damping capacity of the developed Al–Zn–Cu alloys is expected to be good due to the fine lamellar structure. Fig. 12 shows the relationship between the logarithmic decrement and Zn content in the Al–Zn–Cu alloys at room temperature. Contrary to the expected result, the damping capacity of the alloy decreased with increasing Zn additions. The fraction of the Zn-rich phases increased with increasing Zn content. This basal plane (0001) of the η-Zn phase

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indicated a lower plane index and a slip system due to the HCP lattice structure. Consequently, a dislocation can be easily moved by vibration, resulting in the vibration energy being consumed, and thus this alloy exhibits a high damping capacity [31]. However, the damping capacity of the Al–Zn– Cu alloys is highest in the sample with 20% Zn content, and more Zn addition reduces the damping capacity of the alloys. Owing to the non-equilibrium solidification phase and Cu-related phase being enclosed entirely by the α-Al phase in the higher Zn content alloys, we propose that the developed high Zn content alloy would not easily be able to move the dislocation caused by the vibration. Moreover, the hardness of the secondary phase and the reinforcing phase is high, which makes it difficult for interfacial movement to occur [28]. Among each phase of the developed Al–Zn–Cu alloys, the θ phases and η-Zn phases exhibit higher hardness values than do the α phases [9]. Therefore, the lower damping capacity induced by increasing Zn additions causes enhancement of the θ phase and ηZn phase in the non-equilibrium solidification region or grain boundary area, which is entirely consistent with tensile plasticity results in this research. In conclusion, the dominant factor is thought to be the existence of the strengthening phase reducing the damping capacity, rather than the effect of the fine lamellar structure of the α + η phase on the vibration absorption of the alloy.

4. Conclusion 11

In this present study, the effects of a high Zn content on the microstructural evolution, mechanical properties, impacting property, and damping characteristics of Al–Zn–Cu gravity-cast alloys were examined. The following results were obtained: 1) When the developed alloys started to solidify in the flow head, the primary α phase initially crystallized from the melt via a monotectic reaction (L1 = α + L2). Here, L2 remains as a liquid phase, which improves the fluidity of the molten metal. In addition, the melting temperature of the Al–Zn– Cu alloys decreased with increasing Zn content. Moreover, the Cu-related particles in the Al–Zn–Cu alloys are present in a relatively spherical form, unlike the flake-type β-Al5FeSi phase, and thus we propose that this does not inhibit the fluidity in the flow head of the alloy. 2) The size of the α-Al phase significantly decreased and the fraction of the hard phases (such as α + η and θ) increased with increasing Zn content. Furthermore, with increasing Zn content, the developed alloys’ average grain size in the α-phase decreased from 321.16 to 101.41 μm. 3) The increasing fraction of the hard phases with increasing Zn content improved the hardness and strength of the developed alloys. The Al–45Zn–3Cu alloy exhibited UTS of 455 MPa and an elongation of 3.1% without melt modification or post-casting heat treatment.

4) The fractographic formation in the Al–Zn–3Cu alloys transitions from the ductile dimple type to the brittle cleavage-failure type with an increase in Zn addition. 5) The impact strength of the developed Al–Zn–Cu alloys decreased with increasing Zn content. The low Zn content sample showed high Charpy absorbed energy; however, the impact absorption energy of the alloys had a value of approximately 10~14 J in the Al–20~45Zn–3Cu alloys across the whole sample. The damping capacity of the Al–Zn–Cu alloys decreased with increasing Zn additions. The impact strength and damping capacity of the Al–Zn–Cu alloys were reduced by Fe–Si and Cu-related intermetallic compounds as increasing grain boundary regions formed due to high Zn additions.

Acknowledgements This research was funded and supported by the Industrial Core Technology Development Project (10033222). S. S. Shin is grateful to Professor I. M. Park at Pusan National University and Professor J. C. Lee at Korea University for their guidance. 12

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(a)

Temperature ( oC )

Temperature ( oC )

600

400

200

400

Temperature ( oC )

600

Al-Zn20 613 ℃ 600 Al-Zn30 595 ℃ Al-Zn40 575 ℃ Al-Zn45 560 ℃

Al-Zn20 Al-Zn20 Al-Zn30 Al-Zn30 Al-Zn40 Al-Zn40 Al-Zn45 Al-Zn45

(b)

400 260 ℃ (eutectoid reaction)

Al-Zn20 Al-Zn30 Al-Zn40 Al-Zn45

Al-Zn20 Al-Zn30 Al-Zn40 Al-Zn45

200

14

200 0 0

2000

4000

6000

Cooling Time (Sec)

0 0

2000

0 4000 0

Cooling Time (Sec)

60002000

4000

6000

Cooling Time (Sec)

Fig. 1. (a) Cooling curves of the Al–Zn–Cu alloys with various Zn contents, (b) SEM images of the Al–40Zn–3Cu alloys fabricated by gravity casting. This image shows the distribution of precipitates via the eutectoid reaction β  α + η at 260 °

Fig. 2. Optical microstructure of the Al–Zn–3Cu alloys: (a) 20Zn, (b) 30Zn, (c) 40Zn, and (d) 45Zn wt%, 3% HF solution-etched.

15

Fig. 3. (a) – (d) EBSD orientation map (inverse pole figure map) showing the grain distribution of Al– Zn–3Cu alloys with (a) 20Zn, (b) 30Zn, (c) 40Zn, and (d) 45Zn wt. % samples.

Fig. 4. (a) – (d) SEM microscope images recorded from the Al–xZn–3Cu gravity-cast alloys with (a) 20Zn, (b) 30Zn, (c) 40Zn, and (d) 45Zn wt. % samples. 16

Cu related phase

Fig. 5. (a) and (b): SEM micrographs of the HPDC Al–45Zn–3Cu alloys. (c) and (d): SEM micrographs of the gravity-cast Al–45Zn–3Cu alloys. These figures show the typically ultra-fine α + η phase lamellar structure.

EDS 2 EDS 1

EDS 7

EDS 3

EDS 6

EDS 5 EDS 4

EDS No

Al

Zn

Cu

Fe

Si (wt.%)

1

74.41

24.79

0.79

-

2

53.93

21.58

4.01

16.32

4.16

3

8.85

17.89

4.43

-

68.82

4

36.97

28.20

34.83

-

-

5

62.92

35.42

1.66

-

-

6

21.32

71.75

6.93

-

-

7

31.58

30.58

37.84

-

-

-

Fig. 6. SEM micrographs of the microstructure for Al–45Zn–3Cu alloys. The EDS analysis results are shown. 17

75

Rockwell hardness (HRB)

70

65

60

55

50 20

25

30

35

40

45

Zn content ( wt% )

Fig. 7. Variation in the Rockwell hardness of the Al–Zn–3Cu alloys as a function of Zn content.

Engineering stress ( MPa )

500

400

300

Al-20Zn Al-30Zn Al-40Zn Al-45Zn

200

100

strain rate: 1mm/min 0 0

1

2

3

4

5

Engeneering strain ( % )

Fig. 8. Uniaxial tensile stress–strain curves obtained from Al–Zn–Cu alloys. The tensile test specimens are inserted at the bottom of the graph.

18

Fig. 9. Fracture surfaces of Al–Zn–Cu alloys with (a) 20Zn, (b) 30Zn, (c) 40Zn, and (d) 45 wt. % Zn.

(c)

Fig. 10. (a) Optical microscope and (b) Secondary electron image recorded from the side surface of the fractured tensile test on the Al–45Zn–3Cu alloy specimen, showing that crack propagation of the alloy occurred predominantly in the grain boundary region or in the broken reinforcing phase in the specimen. (c) SEM with the EDS results for the fractured tensile Al–45Zn–3Cu alloy specimens.

Fig. 11. Charpy absorbed energy (J) as a function of Zn content at room temperature.

19

Logarithmic decrement (x 10-3 )

0.60

0.56

0.52

0.48

0.44 20

25

30

35

40

45

Zn content ( wt% )

Fig. 12. Relationship between logarithmic decrement and Zn content in Al–Zn alloys at room temperature.

Table 1 The chemical compositions of the Al–Zn–Cu alloys (wt. %) Alloy

Zn Cu

Fe

Si

Al

(a)

20

3.0

0.4

0.5 Balance

(b)

30

3.0

0.4

0.5 Balance

(c)

40

3.0

0.4

0.5 Balance

(d)

45

3.0

0.4

0.5 Balance

20

Table 2 Chemical compositions of the Al–Zn–Cu gravity-casting alloys, as determined by ICP-AES (wt. %) Alloy

Zn

Cu

Fe

Si

Al

(a)

18.70 2.82 0.32 0.68 Balance

(b)

27.59 2.71 0.48 0.54 Balance

(c)

36.24 2.87 0.27 0.64 Balance

(d)

42.86 3.09 0.26 0.63 Balance

Table 3 Thermal properties recorded from the DSC and DTA measurements of the Al–Zn–Cu alloys. Al–Zn20Al–Zn30Al–Zn40Al–Zn45 Liquidus (℃)

608.5 589.05 565.37 551.47

Solidus (℃)

475.1

Latent heat (J/g) Density

432.4

420.9

416.58

291

238

217

182

3.12

3.41

3.52

3.74

Table 4 Tensile properties of the Al–xZn–3Cu gravity-cast alloys tested with different Zn contents. Alloy

Yield Strength (MPa) Tensile Strength (MPa) Elongation (%)

Al–20Zn–3Cu

203

315

4.5

Al–30Zn–3Cu

257

369

3.8

Al–40Zn–3Cu

354

434

3.4

Al–45Zn–3Cu

402

455

3.1

21