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Effect of transition elements on the microstructure and tensile properties of Al–12Si alloy cast under ultrasonic melt treatment
Journal of Alloys and Compounds 712 (2017) 277e287
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Effect of transition elements on the microstructure and tensile properties of Ale12Si alloy cast under ultrasonic melt treatment Jae-Gil Jung a, Sang-Hwa Lee a, b, Young-Hee Cho a, Woon-Ha Yoon a, Tae-Young Ahn c, Yong-Sik Ahn b, Jung-Moo Lee a, * a b c
Implementation Research Division, Korea Institute of Materials Science, Changwon 51508, Republic of Korea Department of Materials Science and Engineering, Pukyong National University, Busan 48513, Republic of Korea Nuclear Materials Safety Research Division, Korea Atomic Energy Research Institute, Daejeon 34057, Republic of Korea
a r t i c l e i n f o
a b s t r a c t
Article history: Received 25 January 2017 Received in revised form 5 April 2017 Accepted 10 April 2017 Available online 12 April 2017
This paper shows that the addition of transition elements significantly influences the microstructure and tensile properties of Ale12Si alloy cast under ultrasonic melt treatment (UST). The alloys contained primary Si, eutectic Si and a variety of intermetallic compounds (IMCs); the addition of Mn and Ni changed the solidification sequence by forming primary a-Al15(Fe,Mn)4Si2 and ε-Al3Ni prior to primary Si formation. The primary Si and primary IMCs competitively consumed the nucleation sites introduced by UST, thereby influencing their refining efficiencies; the efficiency of primary Si refinement by UST increased with increasing the fraction of primary Si formed prior to IMC formation, whereas it deteriorated with increasing primary IMC fraction. The refining efficiency of eutectic Si and eutectic IMCs was not affected by the type and fraction of the primary phase. Using a Ti sonotrode for UST caused Ti contamination, resulting in grain refinement by forming Ti(Al,Si)3 particles and increasing the amount of Ti solutes. Regardless of transition elements content, the application of UST improved the tensile strength at 25 C and 350 C. The refinement of IMCs caused by UST allowed the alloy to contain more IMCs without deterioration of room-temperature tensile strength. This enhanced the increment of hightemperature tensile strength by IMCs. Ductility of heavily alloyed systems was also improved by UST. © 2017 Elsevier B.V. All rights reserved.
Keywords: Aluminum alloy Ultrasonics Nucleation Microstructure Tensile property
1. Introduction Multicomponent AleSi alloys have been widely used in piston production due to their high strength, low coefficient of thermal expansion and excellent castability [1e3]. In addition to Si, these piston alloys generally contain transition elements (e.g., Cu, Ni, Fe, Mn) that form thermally stable intermetallic compounds (IMCs), which make it possible to attain the required mechanical properties [1e3]. With increasing demands for greater fuel efficiency, there is a need to develop high-strength piston alloys that are able to withstand severe operating conditions. To this end, many researchers are currently engaged in optimizing their chemical compositions [4e10] and process variables such as the casting mold material [11] and heat treatment [12]. Ultrasonic melt treatment (UST) has been applied to AleSi piston alloys because it effectively refines the microstructure, thereby
* Corresponding author. E-mail address: [email protected] (J.-M. Lee). http://dx.doi.org/10.1016/j.jallcom.2017.04.084 0925-8388/© 2017 Elsevier B.V. All rights reserved.
improving the mechanical properties [13,14]. The application of UST expands the upper limit of the transition element content that can be used in the alloy by refining the coarse Si and IMCs that form in highly alloyed systems, further improving the mechanical properties of AleSi piston alloys [13]. Additional studies have investigated the combined effects of UST and alloying elements for developing high strength AleSi piston alloys [15e17]. Recently, the present authors [15] examined the combined effects of UST and Si addition on the microstructure and tensile properties of Ale(12,15,18)Sie4Cue3Nie1Mge0.5Fe (wt%) alloys and found a pronounced improvement in the tensile properties at room and elevated temperatures. Sha et al. [16] reported a similar result: the UST-induced improvement in high-temperature tensile strength increased from 23 to 32 MPa with increasing Co content from 0 to 1.05 wt% in the hypereutectic Ale20Sie2Cue1Nie0.6Mge0.7Fe (wt%) system. Lin et al. [17] also showed that the improvement in high-temperature tensile strength with UST increased with increasing concentration of Mn or Fe using a hypereutectic Ale17Sie2Cue1Nie0.4Mg (wt%) alloy.
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The additions of transition elements change the type and formation sequence of secondary phases, thereby influencing the refining behavior of each phase under UST. However, the effect of transition elements on the solidification microstructure and tensile properties of AleSi alloy cast under UST has not been fully understood yet. Therefore, the present study systematically investigates the effects of UST and transition elements (Cu, Ni, Fe and Mn) on the microstructure and mechanical properties of Ale12Si (wt%) piston alloy. The primary phases, eutectic phases, and grains were quantitatively evaluated for distribution, size, and area fraction in alloy microstructures processed with UST. The microstructures were analyzed in relation to the tensile properties at room and elevated temperatures. Based on the findings in this study, we suggest that the refinement of primary phases (Si, IMCs) is dependent on their competitive nucleation under ultrasonic irradiations. 2. Experimental procedure 2.1. Materials Ingots of five Ale12Si piston alloys were provided by Dong Yang Piston Co. (Ansan, Republic of Korea), and their chemical compositions were measured by ARL3460 optical emission spectroscopy and iCAP6500 inductively coupled plasma-optical emission spectroscopy (see Table 1). Note that in addition to different concentrations of transition elements (Cu, Ni, Fe, and Mn), the ingots all contain phosphorus (20e34 ppm), which is necessary to refine primary Si because AlP inclusions act as heterogeneous nucleation sites [18e21]. To examine the role of AlP in refinement of the primary Si particles, Ale25Si binary alloy with and without 0.05 wt% P was also prepared. 2.2. Casting and heat treatment Ale12Si piston alloys and Ale25Si binary alloy were remelted at 750 C and 850 C in an electric resistance furnace, and then degassed by Ar gas bubbling filtration (GBF) during 20 min. The degassed melts, each weighing approximately 1 kg, were then poured into a copper book mold (245 mm 70 mm 200 mm) that had been preheated to 200 C. Considering the formation temperatures of primary Si, the melt pouring temperatures of Ale12Si and Ale25Si alloys were determined as 700 C and 800 C, respectively. The cooling rate of solidifying melts from 650 C to the eutectic arrest temperature (~560 C) was ~7.5 C/s [15]. A titanium sonotrode preheated to 200 C was injected into each 1 kg melt to introduce ultrasonic waves with an amplitude and frequency of 20 mm and 19 kHz, respectively. The UST was performed for 1 min within a temperature range of 750 to 700 C for Ale12Si piston alloys and a temperature range of 850 to 800 C for Ale25Si binary alloy, after which the melts were poured into the same mold. The details about the ultrasonic process can be found in a previous article [15]. The ingots were solution-treated at 490 C for 2 h, and then aged at 230 C for 5 h (i.e., T7 heat treatment).
2.3. Microstructural observations Samples for microstructural observation were taken from each ingot at a position one half of its length and width, and one quarter of its height. The microstructure was observed using an optical microscope (OM, Nikon, MA200) and a scanning electron microscope (SEM, JEOL, JSM-6610LV) equipped with energy diverse X-ray spectroscopy (EDXS, JEOL, INCA Energy). An image analyzer (IMT, iSolution) was used to quantitatively measure the size (maximum length) and area fraction of the secondary phases. Three OM images (200 magnification) were used for analyzing primary Si particles, while ten OM images (1000 magnification) were used for analyzing eutectic Si and IMC particles. The grain structure was examined using an electron backscatter diffraction (EBSD) instrument installed in the SEM (TESCAN, CZ/MIRA I LMH) with a step size of 10 mm. The fcc-Al grains with sizes greater than 50 mm were averaged to separate them from primary Si particles having a diamond cubic structure. The precipitates were observed through a transmission electron microscope (TEM, JEOL, JEM 2100F). The cross-sectional TEM samples were fabricated by using a focused ion beam and they were placed on Mo grids. To estimate the amount of precipitates formed during T7 heat treatment, the electrical resistivity of as-cast and T7 treated alloys was measured at 25 ± 1 C using a sourcemeter (Keithley, Model 6220) and a voltmeter (Keithley, Model 2182), as outlined in ASTM F76e08 [22]. 2.4. Tensile tests Room-temperature tensile tests of T7 treated alloys were performed using an Instron 4206 universal testing machine with a crosshead speed of 1.5 mm/min, as outlined in ASTM E8/E8Me13a [23]. Tensile tests were also performed at 350 C in accordance with ASTM E21e09 [24] after isothermally maintaining each sample at 350 C for 100 h to simulate the piston operating condition. The crosshead speeds during high-temperature tensile tests were 0.125 and 1.5 mm/min before and after the yield point. Dogbone-shaped (gage section: Ø6 25 mm) samples were used for tensile tests. 3. Results 3.1. Thermodynamic calculation Fig. 1 shows the plot of temperature versus solid mole fraction (fs) during Scheil-Gulliver solidification of Alloy 1, which was calculated using Thermo-Calc software [25] with the TCAL3 database. A minor alloying element, Ti, was excluded in the calculation. Primary Si and ε-Al3Ni phases start to form at 577.3 C and 568.9 C, respectively. Then, complex reactions take place at temperatures ranging from 560.0 to 509.8 C, resulting in the formation of eutectic Si, various IMCs (e.g., ε-Al3Ni, d-Al3CuNi, g-Al7Cu4Ni, QAl5Cu2Mg8Si6, q-Al2Cu, M-Mg2Si, etc.) and an fcc-Al matrix. From the thermodynamic calculations, it is expected that various IMCs
Table 1 Chemical composition of Ale12Si piston alloys and their mole fraction of intermetallic compounds. Alloy
1 2 3 4 5
Chemical composition (wt.%)
Mole fraction of intermetallic compound (%)
Al
Si
Cu
Ni
Mg
Fe
Mn
Ti
P
Transition element
Non-equilibrium solidification
Equilibrium at 350 C
Bal. Bal. Bal. Bal. Bal.
12.08 11.90 12.15 12.81 11.97
3.24 3.22 3.32 4.17 3.79
2.36 2.34 2.51 2.26 3.39
0.81 0.97 0.82 0.86 0.82
0.24 0.41 0.24 0.43 0.50
0.20 0.39 0.59 0.35 0.18
0.08 0.09 0.08 0.09 0.10
0.0021 0.0020 0.0022 0.0034 0.0023
6.12 6.45 6.74 7.30 7.96
6.23 8.44 9.08 8.07 11.80
10.45 11.91 12.00 12.41 13.92
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Fig. 1. Temperature-solid mole fraction curve during Scheil-Gulliver solidification of Alloy 1. Secondary phases are indicated as: a-Al15(Fe,Mn)4Si2, b-Al9Fe2Si2, g-Al7Cu4Ni, dAl3CuNi, ε-Al3Ni, p-Al8FeMg3Si6, q-Al2Cu, Q-Al5Cu2Mg8Si6 and M-Mg2Si.
will form similarly to other AleSi piston alloys. The initiation temperature (Ti), accumulated solid mole fraction (fs), and constituent phase of Scheil solidification reactions for the five Ale12Si piston alloys are listed in Table 2. The first solidification reaction is different for each alloy, depending on the concentrations of Si and the transition elements. It is primary Si that forms first in Alloy 1 and 4, as these contain relatively low concentrations of transition elements (or a high concentration of Si). In contrast, IMCs (a-Al15(Fe,Mn)4Si2 or ε-Al3Ni) form first in Alloys 2, 3 and 5 due to their higher concentration of transition elements. Note also that the Ti and fs values for the first solidification reaction increase with increasing concentration of Si or transition elements. The mole fraction of IMC formed during solidification and its equilibrium value at 350 C are shown in Table 1, which reveals an overall increase in fraction with increasing concentration of transition elements. The Fe, Mn and Ni effectively increase the mole fraction of IMCs by forming b-Al9Fe2Si2, a-Al15(Fe,Mn)4Si2 and ε-Al3Ni, but Cu is less effective because of the relatively high concentration required to form IMCs such as g-Al7Cu4Ni and q-Al2Cu as well as its higher solubility in the Al matrix. 3.2. Microstructural observations 3.2.1. As-cast alloy Fig. 2 shows the effect of UST on the microstructure of as-cast Alloys 1, 3 and 5, in which transition elements amount to 6.12, 6.74 and 7.96 wt%, respectively. The as-cast alloys without UST contain dendritic microstructures, whereas equiaxed
microstructures are prevalent in the as-cast alloys with UST. No pores were observed in any of the samples because the melts were sufficiently degassed prior to casting by means of Ar GBF or UST. Significantly, primary Si particles are severely segregated in those alloys without UST, whereas a homogenous distribution of smaller primary Si particles is seen in the alloys with UST. In either case, most of the eutectic Si and IMCs are observed along the cell boundaries. Fig. 3 shows SEM images revealing different sized IMCs with different morphologies in as-cast alloy. The type of IMCs was identified based on their measured chemical compositions, as listed in Table 3. All samples contained eutectic IMCs such as ε-Al3Ni, dAl3CuNi, g-Al7Cu4Ni, Q-Al5Cu2Mg8Si6, M-Mg2Si and q-Al2Cu. In addition to this, coarse primary a-Al15(Fe,Mn)4Si2 was observed in Alloys 2 and 3 (Fig. 3(b) and (e)), while primary ε-Al3Ni was observed in Alloy 5 (Fig. 3(c) and (f)). These primary IMCs present in Alloys 2, 3 and 5 are considered to form first during solidification, as predicted by thermodynamic calculation results (Table 2). It is clearly seen in Fig. 3 that all of the IMCs decreased in size with the application of UST, regardless of the transition element content. Moreover, the agglomeration of needle-like ε-Al3Ni and d-Al3CuNi, which was frequently observed in the alloys without UST, was rarely observed in alloys with UST. Further, the flake-like Ti(Al,Si)3 particles were observed in the alloys with UST. This is attributed to the increased Ti concentration (see Table 4) caused by dissolution of Ti from the ultrasonic horn into the melt during UST [26]. The effect of UST on the grain structure was examined by EBSD analysis. Fig. 4 shows the EBSD microstructure of the fcc a-Al matrix
Table 2 Initiation temperature (Ti), accumulated solid mole fraction (fs), and constituent phase of Scheil solidification reaction for Ale12Si piston alloys. Alloy
1 2 3 4 5
1st reaction
2nd reaction
Ti ( C)
fs (%)
Phase constituent
577.3 578.3 588.8 589.1 589.3
0.58 0.20 0.36 1.00 0.93
L L L L L
þ þ þ þ þ
Si
a a Si ε
3rd reaction
Ti ( C)
fs (%)
Phase constituent
568.9 572.2 578.3 574.7 579.7
1.41 0.47 0.98 2.42 2.88
L L L L L
þ þ þ þ þ
Si þ ε a þ Si a þ Si Si þ a ε þ Si
4th reaction
Ti ( C)
fs (%)
Phase constituent
563.1 569.9 572.8 562.1 567.6
2.60 1.90 2.83 2.69 4.05
L L L L L
þ þ þ þ þ
Si þ ε þ a-Al a þ Si þ ε a þ Si þ ε Si þ a þ ε ε þ Si þ b
Ti ( C)
fs (%)
Phase constituent
560.0 562.0 562.6 560.5 561.1
49.23 36.52 63.30 22.12 15.16
L L L L L
þ þ þ þ þ
Si þ ε þ a þ a-Al a þ Si þ ε þ a-Al a þ Si þ ε þ a-Al Si þ a þ ε þ a-Al ε þ Si þ b þ a-Al
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Fig. 2. OM images of as-cast alloys (aec) without and (def) with UST: (a,d) Alloy 1, (b,e) Alloy 3, (c,f) Alloy 5.
in the as-cast alloys. The effect of UST on grain structure shows a contradiction between the alloys; the application of UST refined the grain structure of Alloy 1, whereas it slightly coarsened the grain structure of Alloy 4. The effect of UST on the size and area fraction of primary Si, eutectic Si, and primary and eutectic IMCs and grain of as-cast alloy was quantitatively evaluated, and the results are summarized in Table 5. In the absence of UST, the size of primary Si in Alloy 1 (15.8 mm) is smaller than those in Alloy 2 (18.7 mm) and Alloy 3 (19.3 mm), whereas it is larger than that in Alloy 5 (14.5 mm). It is thought that the formation of Si-containing a-Al15(Fe,Mn)4Si2 in Alloys 2 and 3 inhibits the nucleation of primary Si by forming a Sidepleted zone near a-Al15(Fe,Mn)4Si2 particles. In contrast, the formation of ε-Al3Ni in Alloy 5 might accelerate the nucleation of primary Si by forming a Si-enriched zone near ε-Al3Ni particles. Moreover, the larger primary Si size for Alloy 4 (21.5 mm) is due to its higher Si concentration (12.8 wt%) in comparison to those (11.9e12.1 wt%) of other alloys. Table 5 also shows that UST reduces the size and increases the area fraction of primary Si particles. In addition, UST reduces the size of eutectic Si in as-cast alloys except in Alloy 3, which has fine Chinese-script eutectic Si particles in the absence of UST (Fig. 2(b)). The area fraction of eutectic Si is reduced in UST alloys through an increase in the area fraction of primary Si. The UST also decreases
the size of IMCs in UST alloys. However, it would seem that UST does not significantly change the total area fraction of IMCs, although it causes the formation of some Ti(Al,Si)3 particles. Meanwhile, there was a significant decrease in the average grain size with UST except for the Alloy 4. 3.2.2. Heat-treated alloy Fig. 5(a) and (b) present OM images of as-cast and T7 treated Alloy 5 without UST. The constituent phases (i.e., primary Si, eutectic Si and IMCs) are the same in both specimens. Notably, however, the interfaces between the primary Si and Al matrix, as well as between eutectic Si and the Al matrix, are smoother in the T7 alloy. The T7 alloys also show evidence of a breakup of eutectic Si particles that is caused by their spheroidization during solution treatment at 490 C. In addition, some IMCs (e.g. Mg2Si, QAl5Cu2Mg8Si6, g-Al7Cu4Ni and Al2Cu), which are thermally unstable at 490 C were partially dissolved during solution treatment. The artificial aging at 230 C results in the formation of fine precipitates within the Al matrix, as shown in Fig. 5(c). The needlelike particles are q'-Al2Cu and relatively thicker ones are q-Al2Cu, which are confirmed by analyses of their selected area diffraction patterns taken along the [110]Al. Needle-like q'-Al2Cu particles generate streaks along the [200]Al [27] and q-Al2Cu particles are formed with a crystallographic orientation of [110]Al//[531]q.
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Fig. 3. SEM images of as-cast alloys (aec) without and (def) with UST: (a,d) Alloy 1, (b,e) Alloy 3, (c,f) Alloy 5.
Table 3 Chemical composition of intermetallic compounds observed in Ale12Si piston alloys. No.
1 2 3 4 5 6 7 8
Chemical composition (at. %)
Suggested phase
Mg
Al
Si
Fe
Mn
Ni
Cu
Ti
e 59.4 0.2 e 33.3 0.6 e e
67.4 7.5 73.3 68.1 28.6 77.4 65.1 42.1
0.9 33.1 3.7 12.1 30.4 3.4 e 35.0
0.1 e 1.5 3.8 e 0.9 e e
e e 0.6 12.7 e 3.3 e e
0.5 e 10.2 3.3 0.2 12.7 9.7 e
23.1 e 10.8 e 7.5 1.8 25.1 e
e e e e e e e 17.0
Table 4 Ti concentrations of Ale12Si piston alloys (wt%). Alloy
w/o UST
w/UST
1 2 3 4 5
0.08 0.09 0.08 0.09 0.10
0.21 0.24 0.24 0.15 0.26
q-Al2Cu M-Mg2Si d-Al3CuNi a-Al15(Mn,Fe)4Si2 Q-Al5Cu2Mg8Si6 ε-Al3Ni g-Al7Cu4Ni Ti(Al,Si)3
The decreased electrical resistivity values of the T7 samples in comparison to those of the as-cast samples (Fig. 6) also indicate that the precipitation occurs during aging at 230 C [28]. The difference in electrical resistivity values (Dr) between as-cast and T7 alloys increases with increasing transition element content. This means that the amount of precipitates (e.g., q'-Al2Cu) formed during aging at 230 C is larger in highly alloyed system. However, these precipitates formed during aging are known to have an insignificant effect on strength at 350 C [29]. This is because most of q'-Al2Cu particles were already transformed to coarse q-Al2Cu particles
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Fig. 4. EBSD microstructures of as-cast (a,b) Alloy 1 and (c,d) Alloy 4 (a,c) without and (b,d) with UST.
Table 5 Effect of UST on microstructural characteristics of as-cast Ale12Si piston alloys. Alloy
UST
Primary Si Area fraction (%)
1 2 3 4 5
w/o w/ w/o w/ w/o w/ w/o w/ w/o w/
2.7 2.5 1.1 2.1 1.5 5.4 2.1 4.6 1.4 3.8
± ± ± ± ± ± ± ± ± ±
0.1 0.2 1.1 0.7 0.5 0.6 0.9 1.1 0.1 0.7
Eutectic Si Size (mm) 15.8 10.8 18.7 12.5 19.3 15.7 21.5 13.4 14.5 13.5
± ± ± ± ± ± ± ± ± ±
2.1 1.3 2.2 0.3 1.2 0.9 1.0 2.1 2.5 1.0
Primary and eutectic IMCs
Refining efficiency (%)
Area fraction (%)
Size (mm)
Refining efficiency (%)
Area fraction (%)
31.6
7.9 ± 0.6 6.0 ± 0.6 8.5 ± 1.9 6.4 ± 1.0 10.8 ± 2.5 9.3 ± 1.1 7.2 ± 1.6 6.3 ± 0.9 8.1 ± 2.0 6.0 ± 0.4
10.3 ± 1.0 9.2 ± 0.8 10.9 ± 1.0 9.6 ± 1.4 3.3 ± 0.3 8.1 ± 1.7 12.8 ± 2.0 10.4 ± 0.7 9.2 ± 0.9 8.3 ± 0.9
10.7
6.3 5.5 5.9 6.9 6.5 6.5 7.8 7.3 8.3 8.7
33.2 18.7 37.7 6.9
11.9 e 18.8 9.8
± ± ± ± ± ± ± ± ± ±
0.7 0.9 2.1 1.2 2.2 1.7 1.6 1.2 2.4 1.1
Size (mm)
Refining efficiency (%)
8.7 ± 0.9 6.9 ± 0.1 9.0 ± 0.9 7.7 ± 0.3 10.9 ± 3.0 9.8 ± 0.4 9.1 ± 0.8 7.2 ± 0.3 9.5 ± 1.8 7.6 ± 0.8
20.7 14.4 10.1 20.9 20.0
Grain size (mm)
577 241 318 235 598 272 264 362 313 166
± ± ± ± ± ± ± ± ± ±
252 87 147 86 320 95 97 143 135 62
during isothermal holding at 350 C for 100 h before the tensile test, as shown in Fig. 5(d).
Alloy 5 (7.96 wt%), however, the decrease in dimple size from 22.8 to 14.9 mm with UST results in improved ductility (Fig. 7(c) and (d)).
3.3. Tensile properties
4. Discussion
Table 6 lists the tensile strength and elongation at room and elevated temperatures for the T7 alloys. It is clear that both the tensile strength and elongation seem to decrease with increasing transition element content due to the increased area fraction and size of brittle IMCs, which make the alloy vulnerable to tensile cracking at room temperature [13,15]. At the elevated temperature of 350 C, it can be seen that the yield strength and tensile strength increase with the transition element content due to the increase in the IMC fraction. In contrast, the elongation decreases with increasing transition element content. For all alloys, UST improves the tensile strength at room and elevated temperatures. However, the effect of UST on elongation is different depending on the amount of transition elements. The UST increases the elongation of Alloys 3, 4 and 5, whereas it deteriorates the elongation of Alloys 1 and 2 despite refinement of their microstructures. The contradictory effects of UST on the elongation as a function of transition element content are also visible in the tensile-fractured surfaces at 350 C. In Alloy 1 with fewer transition elements (6.12 wt%), the dimple size increased with UST from 14.6 to 16.1 mm, which is in agreement with the observed deterioration in ductility (Fig. 7(a) and (b)). In heavily alloyed systems such as
4.1. Transition elements combined with UST effects on solidification microstructure 4.1.1. Refining mechanism of UST The UST decreased the sizes of secondary phases as shown in Table 5. The refinement of secondary phases by UST is generally explained by cavitation-induced dendrite fragmentation and cavitation-induced heterogeneous nucleation [30e33]. Since the UST was terminated above the liquidus temperature in this study, dendrite fragmentation is not responsible for the refinement of secondary phases [15]. Although nucleation possibly occurs near the cold sonotrode during UST, this effect on the microstructural refinement is not significant, according to Qian et al. [33]. Thus, the cavitation-induced heterogeneous nucleation mechanism is considered to be responsible for the refinement of phases. The cavitation-induced heterogeneous nucleation mechanism can be further divided into three different subcategories [33]. The first is the pressure pulseemelting point mechanism based on the Clapeyron equation. The collapse of bubbles increases the pressure on the melt, and thus, increases the melting point. The increased melting point is equivalent to increased undercooling, leading to
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Fig. 5. OM images of (a) as-cast and (b) T7 Alloy 5 without UST. TEM images of T7 Alloy 5 (c) before and (d) after isothermal holding at 350 C for high-temperature tensile tests.
Fig. 6. Electrical resistivity values of as-cast and T7 alloys without UST.
enhanced nucleation. The second mechanism is related to the fact that the adiabatic expansion of cavitation bubbles undercools the bubbleeliquid interfaces, enhancing the nucleation on the bubble surfaces. The third mechanism is the cavitation-enhanced wetting of inclusions. Non-metallic inclusions (e.g., MgAl2O4 [34], AlP [20]) exist in the Al melt containing Mg and P, and their wettability is improved by the formation of bubbles through the degassing of hydrogen gas trapped at the surface of inclusions. The wetted inclusions can act as effective nucleation sites during the casting. Among these mechanisms, the first may not be prevalent here because the change in pressure due to bubble collapse disappears immediately after UST is terminated. This is supported by the fact that the fraction of primary Si was increased by the application of UST (Table 5); the activation of a pressure pulseemelting point mechanism is known to decrease the fraction of primary Si by shifting the AleSi eutectic point toward the high Si region [16]. It is thought that the wettable inclusions such as MgAl2O4 (cubic; a ¼ 8.08 Å [34]) and AlP (cubic; a ¼ 5.42 Å [20]) might remain wet in the melt even after ultrasonic treatment, thereby acting as nucleation sites for primary Si during casting. Wang et al. [35] and
Table 6 Effect of UST and transition elements on tensile properties at room and elevated temperatures of T7 piston alloys. Alloy
UST
25 C
350 C
Tensile strength (MPa) 1 2 3 4 5
w/o w/ w/o w/ w/o w/ w/o w/ w/o w/
281 296 298 305 243 309 242 285 248 285
± ± ± ± ± ± ± ± ± ±
27 22 24 17 31 23 23 8 10 7
Elongation (%) 0.71 0.65 0.89 0.78 0.49 0.95 0.39 0.42 0.41 0.53
± ± ± ± ± ± ± ± ± ±
0.26 0.15 0.23 0.17 0.15 0.24 0.07 0.03 0.05 0.04
Elongation (%)
Yield strength (MPa) 38.5 39.7 39.8 39.8 41.5 43.7 41.0 42.0 45.8 45.3
± ± ± ± ± ± ± ± ± ±
0.8 1.2 1.0 3.1 2.6 0.9 1.6 0.8 3.0 4.1
Tensile strength (MPa) 63.0 65.5 65.5 68.5 67.0 68.1 67.7 68.3 71.9 76.1
± ± ± ± ± ± ± ± ± ±
0.1 1.4 2.0 0.4 0.9 0.6 0.7 1.3 2.2 2.3
21.2 ± 4.0 17.9 ± 1.2 16.6 ± 1.2 13.8 ± 1.9 9.9 ± 2.0 11.1 ± 4.7 3.7 ± 1.2 12.4 ± 4.5 6.9 ± 3.3 9.1 ± 2.3
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Fig. 7. Tensile-fractured surfaces of (a,b) Alloy 1 and (c,d) Alloy 5 (a,c) without and (b,d) with UST. Tensile tests were performed at 350 C.
Komarov et al. [36] also suggested that the cavitation-induced nucleation via improved wetting of inclusions plays the dominant role in the refinement of primary phases when UST is performed above the liquidus temperature. The roles of oxide and AlP in refining primary Si particles under the UST were investigated further using the Ale25Si binary alloy. It is seen from Fig. 8 that the UST decreased the size of primary Si in the Ale25Si binary alloy with and without AlP inclusions; the size of primary Si was decreased from 38.8 to 27.5 mm and from 143 to 90 mm with and without P addition, respectively. The results suggest that the wetted oxides can act as nucleation sites for the primary Si particles despite different lattice constants between oxide (e.g., MgAl2O4, cubic; a ¼ 8.08 Å [34]) and Si (cubic; 5.43 Å [20]). In a similar fashion, the wetted oxides and AlP may enhance the nucleation of primary IMCs such as a-Al15(Fe,Mn)4Si2 (cubic; a ¼ ~12.68 Å [37]) and ε-Al3Ni (orthorhombic; a ¼ 6.60 Å, b ¼ 7.35 Å, c ¼ 4.80 Å [38]). Meanwhile, it is unclear that the second mechanism (i.e., enhanced nucleation near bubble surfaces) is active in this system. In addition to cavitation-induced heterogeneous nucleation, acoustic streaming that homogenizes the alloying elements and distributes the wetted inclusions over the whole melt also plays a significant role in refining the primary phases. 4.1.2. Competitive refinement of primary phase It is worth noting that the efficiency of primary Si refinement by UST (i.e., the decrease in primary Si size) was different in each alloy as shown in Table 5: Alloy 1 (31.6%), Alloy 2 (33.2%), Alloy 3 (18.7%), Alloy 4 (37.7%) and Alloy 5 (6.9%). Fig. 9(a) shows the relationship between refining efficiency of primary Si and the calculated mole fraction of the primary Si (fs of primary Si for the first reaction in Table 2). For a better understanding, the experimental results obtained in Ale(12,15,18)Si piston alloys [15] and Ale25Si binary alloys are also plotted. It is evident that the refining efficiency of primary Si is greatly increased with increasing primary Si fraction up to 1%, but becomes saturated at approximately 35%. Conversely, the refining efficiency of primary Si decreases with
increasing mole fraction of primary IMC (Fig. 9(b)), which means that any primary IMC formed prior to primary Si consumes nucleation sites. This is supported by the drastic refinement of primary IMC in Alloy 3 and Alloy 5 with poor refining efficiency of primary Si; the sizes of primary a-Al15(Fe,Mn)4Si2 in Alloy 3 and primary ε-Al3Ni in Alloy 5 decreased significantly from 36 to 23 mm and from 50 to 17 mm, respectively. The consumption of nucleation sites by primary IMC limits the opportunity for primary Si nucleation and, consequently, reduces the efficiency of primary Si refinement by UST. This implies that AleSi piston alloys require a sufficient amount (>1%) of pre-existing primary Si particles prior to IMC formation to achieve the maximum refining efficiency using ultrasonic irradiation. Consequently, it is necessary to optimize the amounts of added Si and transition elements in relation to the formation sequence and mole fraction of the primary phase (i.e., Si or IMC). 4.1.3. Eutectic phase and grain The refining efficiency of eutectic Si by UST ranges from 9.8 to 18.8% as shown in Table 5. It is thought that the refining efficiency of eutectic Si was not affected by the type and mole fraction of the primary phase because the eutectic Si phases are formed after the ultrasonically introduced nucleation sites are nearly consumed by the formation of primary phases. Also, the efficiency of IMC refinement (10.1e20.9%) was not influenced by the primary phase, because most of the IMCs are eutectic phases. According to Das and Kotadia [39,40], the refining efficiency of eutectic phases can be improved by the application of UST during solidification. They showed that the UST performed during solidification significantly decreased the size of eutectic Si particles in hypoeutectic Ale10Si [39] and Ale7Si [40] alloys, thereby achieving greater refining efficiency (>70%) of the eutectic Si. They suggested that such a significant refinement of eutectic Si was attributed to the intense convection from cavitation near the sonotrode that homogenizes the solutes at the eutectic growth front [40]. It is well known that both TiAl3 particles and the solute Ti atoms cause grain refinement by providing potential nucleation sites for
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Fig. 8. OM images of as-cast (a,b) Ale25Si and (c,d) Ale25Sie0.05P alloys (a,c) without and (b,d) with UST.
Fig. 9. Effects of mole fraction of (a) primary Si and (b) primary IMC formed during 1st solidification reaction on efficiency of primary Si refinement in UST alloys. Relationships observed in Ale(12,15,18)Si piston alloys [15] and Ale25Si binary alloys are plotted together.
a-Al and restrict the growth of a-Al [41,42]. In this study, the UST using the Ti sonotrode increased the Ti concentration from 0.08e0.10 to 0.15e0.26 wt% (Table 4). Taking into account a peritectic composition in an AleTi binary alloy (~0.15 wt%) [41], UST alloys probably contain numerous TiAl3 and/or Ti(Al,Si)3 particles and a large amount of solute Ti atoms. Thus, the effect of Ti concentration on the grain size of as-cast alloys should be considered along with the UST effect on grain size. Fig. 10 shows the effect of Ti concentration on the grain size of as-cast Ale12Si piston alloys. The grain size gradually decreases with increasing Ti concentration. Overall, the UST alloys having higher Ti concentrations exhibit smaller grain sizes than the alloys without UST. The grain size of Alloy 4 was slightly increased by application of UST. However, it is not reasonable to conclude that the UST causes the grain coarsening in Alloy 4 considering the large standard deviation values. This analysis strongly suggests that the grain size of as-cast alloys depends on the Ti concentration rather than on the application of UST [15,43].
Fig. 10. Effect of Ti content on the grain size of as-cast alloys without and with UST.
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Fig. 11. Effect of IMC fraction on tensile strength at (a) 25 C and (b) 350 C.
4.2. Transition elements combined with UST effects on tensile properties Fig. 11(a) shows the relationship between the measured IMC fraction and tensile strength at 25 C. In the absence of UST, the room-temperature tensile strength decreases with increasing IMC fraction. In the presence of UST, however, the room-temperature tensile strength remains at ~300 MPa until the IMC fraction reaches 7%, and then decreases with further IMC fraction. This implies that the refinement of IMCs by UST allows the piston alloy to contain more IMCs without deterioration of strength. The increased IMC fraction leads to an improvement in high-temperature tensile strength as shown in Fig. 11(b). Furthermore, the UST increases the slope of the plot of tensile strength versus the IMC fraction from 2.6 to 3.2 MPa/%. This is likely due to the refinement of primary and eutectic phases. It is worth noting that while UST causes a decrease in the elongation of Alloy 1 and 2, it actually increases the elongation of Alloys 3, 4 and 5. The alloy specimens subjected to UST were found to contain coarse flake-like Ti(Al,Si)3 particles that are deleterious to ductility. With those alloys containing fewer transition elements, (i.e., Alloys 1 and 2), the elongation is relatively high due to a smaller IMC fraction, and so these are more significantly affected by deleterious Ti(Al,Si)3 particles. In contrast, alloys with a higher transition element content (i.e., Alloys 3, 4 and 5) experience a beneficial refinement of microstructure with UST that can overcome the deleterious effect of Ti(Al,Si)3 particles.
by UST. The decreased fraction of primary Si indicates that the refinement of primary phases can be attributed to the enhanced nucleation on the cavitation bubble surfaces and the wetted inclusions when UST is used. (3) The application of UST also reduces the sizes of eutectic Si and IMC, but the refining efficiency in this case is not influenced by the type or fraction of the primary phase. The formation of coarse flake-like Ti(Al,Si)3 particles caused by the dissolution of Ti elements from the titanium ultrasonic horn can also occur, resulting in the grain refinement. (4) At room temperature, both the tensile strength and ductility simultaneously decrease with increasing transition element content. At 350 C, however, the yield strength and tensile strength increase with increasing transition element content at the expense of ductility. Regardless of transition element content, the UST improves the tensile strength at room and elevated temperatures, which is mainly attributed to a refinement of the microstructure. (5) There is also an improvement in ductility at elevated temperatures with UST when the transition element content is relatively high, implying that the beneficial effect of UST is greater in heavily alloyed systems. The deterioration in ductility that occurs in alloys with lower transition element content is most likely because the deleterious effect of coarse flake Ti(Al,Si)3 is greater than the beneficial effect of UST (i.e., microstructure refinement) in these alloys. Acknowledgements
5. Conclusions (1) The microstructure of Ale12Si piston alloys is composed of primary Si, eutectic Si and a variety of IMCs (e.g., ε-Al3Ni, dAl3CuNi, g-Al7Cu4Ni, Q-Al5Cu2Mg8Si6, Mg2Si and q-Al2Cu). The addition of Mn and Ni can change the solidification sequence by forming primary a-Al15(Fe,Mn)4Si2 or ε-Al3Ni prior to primary Si formation. Both the formation temperature and fraction of primary phase (Si or IMC) increase with increasing Si or transition element content. (2) The UST significantly decreases the size of primary Si and primary IMCs. The refinement efficiency of primary Si increases with an increase in primary Si fraction or decrease in primary IMC fraction formed during the first solidification reaction. The consumption of nucleation sites by primary IMC limits the opportunity for primary Si nucleation and, consequently, reduces the efficiency of primary Si refinement
The authors would like to acknowledge the financial support received from the R&D Convergence Program of the MSIP (Ministry of Science, ICT and Future Planning) and NST (National Research Council of Science and Technology) of the Republic of Korea (Grant CCPe13e17eKIMS), and from the Main Research Program of the Korea Institute of Materials Science (PNK4711). References [1] J.R. Davies, Aluminum and Aluminum Alloys, ASM International, OH, 1993. [2] N. Belov, D. Eskin, N. Avxentieva, Constituent phase diagrams of the AleCueFeeMgeNieSi system and their application to the analysis of aluminum piston alloys, Acta Mater 53 (2005) 4709e4722. [3] R. Gholizadeh, S.G. Shabestari, Investigation of the effects of Ni, Fe, and Mn on the formation of complex intermetallic compounds in AleSieCueMgeNi alloys, Metall, Mater. Trans. A 42A (2011) 3447e3458. [4] A.J. Moffat, S. Barnes, B.G. Mellor, P.A.S. Reed, The effect of silicon content on long crack fatigue behavior of aluminum-silicon piston alloys at elevated
J.-G. Jung et al. / Journal of Alloys and Compounds 712 (2017) 277e287 temperature, Inter. J. Fatigue 27 (2005) 1564e1570. [5] Z. Qian, X. Liu, D. Zhao, G. Zhang, Effects of trace Mn addition on the elevated temperature tensile strength and microstructure of a low-iron AleSi piston alloy, Mater. Lett. 62 (2008) 2146e2149. [6] Y. Li, Y. Wu, Z. Qian, X. Liu, Effect of co-addition of RE, Fe and Mn on the microstructure and performance of A390 alloy, Mater. Sci. Eng. A 527 (2009) 146e149. [7] Y. Li, Y. Yang, Y. Wu, L. Wang, X. Liu, Quantitative comparison of three Niecontaining phases to the elevated-temperature properties of AleSi piston alloys, Mater. Sci. Eng. A 527 (2010) 7132e7137. [8] Y. Yang, K. Yu, Y. Li, D. Zhao, X. Liu, Evolution of nickel-rich phases in AleSieCueNieMg piston alloys with different Cu additions, Mater. Des. 33 (2012) 220e225. [9] Y. Li, Y. Yang, Y. Wu, Z. Wei, X. Liu, Supportive strengthening role of Crerich phase on AleSi multicomponent piston alloy at elevated temperature, Mater. Sci. Eng. A 528 (2011) 4427e4430. [10] C.-Y. Jeong, Effect of alloying elements on high temperature mechanical properties for piston alloy, Mater. Trans. 53 (2012) 234e239. [11] M.M. Haque, M.A. Maleque, Effect of process variables on structure and properties of aluminiumesilicon piston alloy, J. Mater. Proc. Tech. 77 (1998) 122e128. [12] S. Manasijevi c, S. Markovi c, Z. A cimovi c-Pavlovi c, K. Rai c, R. Radisa, Effect of heat treatment on the microstructure and mechanical properties of piston alloys, Mater Tech. 47 (2013) 585e591. [13] J.-G. Jung, S.-H. Lee, J.-M. Lee, Y.-H. Cho, S.-H. Kim, W.-H. Yoon, Improved mechanical properties of near-eutectic AleSi piston alloy through ultrasonic melt treatment, Mater. Sci. Eng. A 669 (2016) 187e195. [14] W. Khalifa, S. El-Hadad, Y. Tsunekawa, Microstructure and wear behavior of solidification sonoprocessed B390 hypereutectic AleSi alloy, Metall, Mater. Trans. A 44A (2013) 5817e5824. [15] J.-G. Jung, J.-M. Lee, Y.-H. Cho, W.-H. Yoon, Combined effects of ultrasonic melt treatment, Si addition and solution treatment on the microstructure and tensile properties of multicomponent AleSi alloys, J. Alloys Compd. 693 (2017) 201e210. [16] M. Sha, S. Wu, L. Wan, Combined effects of cobalt addition and ultrasonic vibration on microstructure and mechanical properties of hypereutectic AleSi alloys with 0.7%Fe, Mater. Sci. Eng. A 554 (2012) 142e148. } , P. An, L. Wan, Effects of ultrasonic vibration and manga[17] C. Lin, S. Wu, S. Lu nese on microstructure and mechanical properties of hypereutectic AleSi alloys with 2%Fe, Intermetallics 32 (2013) 176e183. [18] P.B. Crosley, L.F. Mondolfo, The modification of aluminum-silicon alloys, AFS Trans. 74 (1966) 53e64. [19] C.R. Ho, B. Cantor, Heterogeneous nucleation of solidification of Si in AleSi and AleSieP alloys, Acta Metall. Mater. 43 (1995) 3231e3246. [20] K. Nogita, S.D. McDonald, K. Tsujimoto, K. Yasuda, A.K. Dahle, Aluminium phosphide eutectic grain nucleus in hypoeutectic AleSi alloys, J. Electron Microsc. 53 (2004) 361e369. [21] J. Li, F.S. Hage, X. Liu, Q. Ramasse, P. Schumacher, Revealing heterogeneous nucleation of primary Si and eutectic Si by AlP in hypereutectic AleSi alloys, Sci. Rep. 6 (2016), 25244-1e8. [22] ASTM F76-08, Standard Test Methods for Measuring Resistivity and Hall Coefficient and Determining Hall Mobility in Singleecrystal Semiconductors, ASTM International, PA, 2008.
287
[23] ASTM E8/E8Me13a, Standard Test Methods for Tension Testing of Metallic Materials, ASTM International, PA, 2013. [24] ASTM E21e09, Standard Test Methods for Elevated Temperature Tension Tests of Metallic Materials, ASTM International, PA, 2009. [25] N. Saunders, A.P. Miodownik, CALPHAD, Elsevier, Oxford, 1998. [26] F. Dong, X. Li, L. Zhang, L. Ma, R. Li, Cavitation erosion mechanism of titanium alloy radiation rods in aluminum melt, Ultrason. Sonochem. 31 (2016) 150e156. [27] A.K. Shukla, W.A. Baeslack III, Study of microstructural evolution in frictionstir welded thin-sheet AleCueLi alloy using transmission-electron microscopy, Scr. Mater 56 (2007) 513e516. [28] J.-G. Jung, J.-S. Park, Y.-K. Lee, Quantitative analyses of dissolution and precipitation kinetics of q-Al2Cu phase in an Ale6.2Sie2.9Cu alloy using electrical resistivity, Met. Mater. Int. 19 (2013) 147e152. [29] Z. Asghar, G. Requena, E. Boller, Three-dimensional rigid multiphase networks providing high-temperature strength to cast AlSi10Cu5Ni1-2 piston alloys, Acta Mater 59 (2011) 6420e6432. [30] M.C. Flemings, Solidification Processing, McGraw-Hill, New York, 1974. [31] G.I. Eskin, Ultrasonic Treatment of Light Alloy Melts, Gordon & Breach, Amsterdam, 1998. [32] J.D. Hunt, K.A. Jackson, Nucleation of solid in an undercooled liquid by cavitation, J. Appl. Phys. 37 (1966) 254e257. [33] M. Qian, A. Ramirez, A. Das, Ultrasonic refinement of magnesium by cavitation: clarifying the role of wall crystals, J. Cryst. Growth 311 (2009) 3708e3715. [34] Y. Wang, H.-T. Li, Z. Fan, Oxidation of aluminium alloy melts and inoculation by oxide particles, Trans. Indian Inst. Met. 65 (2012) 653e661. [35] F. Wang, D. Eskin, T. Connolley, J. Mi, Effect of ultrasonic melt treatment on the refinement of primary Al3Ti intermetallic in an Ale0.4Ti alloy, J. Cryst. Growth 435 (2016) 24e30. [36] S. Komarov, Y. Ishiwata, I. Mikhailov, Industrial application of ultrasonic vibrations to improve the structure of AleSi hypereutectic alloys: potential and limitations, Metall. Mater. Trans. A 46A (2015) 2876e2883. [37] A.K. Srivastaba, S. Ranganathan, Quasicrystals, crystals and multiple twins in rapidly solidified AleCreSi, AleMneSi and AleMneCreSi alloys, Acta Mater. 44 (1996) 2935e2946. [38] N.D. Broom, G.J. Davies, Plastic deformation, fracture and dislocation structures in the orthorhombic intermetallic phase Al3Ni, Acta Metall. 23 (1975) 537e546. [39] A. Das, H.R. Kotadia, Effect of high-intensity ultrasonic irradiation on the modification of solidification microstructure in a Si-rich hypoeutectic AleSi alloy, Mater. Chem. Phys. 125 (2011) 853e859. [40] H.R. Kotadia, A. Das, Modification of solidification microstructure in hypo- and hyper-eutectic Al-Si alloys under high-intensity ultrasonic irradiation, J. Alloys Compd. 620 (2015) 1e4. [41] M. Easton, D. StJohn, Grain refinement of aluminum alloys: Part. I. the nucleant and solute paradigmsea review of the literature, Metall. Mater. Trans. A 30A (1999) 1613e1623. [42] Z. Fan, Y. Wang, Y. Zhang, T. Qin, X.R. Zhou, G.E. Thompson, T. Pennycook, T. Hashimoto, Grain refining mechanism in the Al/AleTieB system, Acta. Mater. 84 (2015) 292e304. [43] D. Eskin, Ultrasonic melt processing: opportunities and misconceptions, Mater. Sci. Forum 794e796 (2014) 101e106.
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Effect of transition elements on the microstructure and tensile properties of Ale12Si alloy cast under ultrasonic melt treatment Jae-Gil Jung a, Sang-Hwa Lee a, b, Young-Hee Cho a, Woon-Ha Yoon a, Tae-Young Ahn c, Yong-Sik Ahn b, Jung-Moo Lee a, * a b c
Implementation Research Division, Korea Institute of Materials Science, Changwon 51508, Republic of Korea Department of Materials Science and Engineering, Pukyong National University, Busan 48513, Republic of Korea Nuclear Materials Safety Research Division, Korea Atomic Energy Research Institute, Daejeon 34057, Republic of Korea
a r t i c l e i n f o
a b s t r a c t
Article history: Received 25 January 2017 Received in revised form 5 April 2017 Accepted 10 April 2017 Available online 12 April 2017
This paper shows that the addition of transition elements significantly influences the microstructure and tensile properties of Ale12Si alloy cast under ultrasonic melt treatment (UST). The alloys contained primary Si, eutectic Si and a variety of intermetallic compounds (IMCs); the addition of Mn and Ni changed the solidification sequence by forming primary a-Al15(Fe,Mn)4Si2 and ε-Al3Ni prior to primary Si formation. The primary Si and primary IMCs competitively consumed the nucleation sites introduced by UST, thereby influencing their refining efficiencies; the efficiency of primary Si refinement by UST increased with increasing the fraction of primary Si formed prior to IMC formation, whereas it deteriorated with increasing primary IMC fraction. The refining efficiency of eutectic Si and eutectic IMCs was not affected by the type and fraction of the primary phase. Using a Ti sonotrode for UST caused Ti contamination, resulting in grain refinement by forming Ti(Al,Si)3 particles and increasing the amount of Ti solutes. Regardless of transition elements content, the application of UST improved the tensile strength at 25 C and 350 C. The refinement of IMCs caused by UST allowed the alloy to contain more IMCs without deterioration of room-temperature tensile strength. This enhanced the increment of hightemperature tensile strength by IMCs. Ductility of heavily alloyed systems was also improved by UST. © 2017 Elsevier B.V. All rights reserved.
Keywords: Aluminum alloy Ultrasonics Nucleation Microstructure Tensile property
1. Introduction Multicomponent AleSi alloys have been widely used in piston production due to their high strength, low coefficient of thermal expansion and excellent castability [1e3]. In addition to Si, these piston alloys generally contain transition elements (e.g., Cu, Ni, Fe, Mn) that form thermally stable intermetallic compounds (IMCs), which make it possible to attain the required mechanical properties [1e3]. With increasing demands for greater fuel efficiency, there is a need to develop high-strength piston alloys that are able to withstand severe operating conditions. To this end, many researchers are currently engaged in optimizing their chemical compositions [4e10] and process variables such as the casting mold material [11] and heat treatment [12]. Ultrasonic melt treatment (UST) has been applied to AleSi piston alloys because it effectively refines the microstructure, thereby
* Corresponding author. E-mail address: [email protected] (J.-M. Lee). http://dx.doi.org/10.1016/j.jallcom.2017.04.084 0925-8388/© 2017 Elsevier B.V. All rights reserved.
improving the mechanical properties [13,14]. The application of UST expands the upper limit of the transition element content that can be used in the alloy by refining the coarse Si and IMCs that form in highly alloyed systems, further improving the mechanical properties of AleSi piston alloys [13]. Additional studies have investigated the combined effects of UST and alloying elements for developing high strength AleSi piston alloys [15e17]. Recently, the present authors [15] examined the combined effects of UST and Si addition on the microstructure and tensile properties of Ale(12,15,18)Sie4Cue3Nie1Mge0.5Fe (wt%) alloys and found a pronounced improvement in the tensile properties at room and elevated temperatures. Sha et al. [16] reported a similar result: the UST-induced improvement in high-temperature tensile strength increased from 23 to 32 MPa with increasing Co content from 0 to 1.05 wt% in the hypereutectic Ale20Sie2Cue1Nie0.6Mge0.7Fe (wt%) system. Lin et al. [17] also showed that the improvement in high-temperature tensile strength with UST increased with increasing concentration of Mn or Fe using a hypereutectic Ale17Sie2Cue1Nie0.4Mg (wt%) alloy.
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The additions of transition elements change the type and formation sequence of secondary phases, thereby influencing the refining behavior of each phase under UST. However, the effect of transition elements on the solidification microstructure and tensile properties of AleSi alloy cast under UST has not been fully understood yet. Therefore, the present study systematically investigates the effects of UST and transition elements (Cu, Ni, Fe and Mn) on the microstructure and mechanical properties of Ale12Si (wt%) piston alloy. The primary phases, eutectic phases, and grains were quantitatively evaluated for distribution, size, and area fraction in alloy microstructures processed with UST. The microstructures were analyzed in relation to the tensile properties at room and elevated temperatures. Based on the findings in this study, we suggest that the refinement of primary phases (Si, IMCs) is dependent on their competitive nucleation under ultrasonic irradiations. 2. Experimental procedure 2.1. Materials Ingots of five Ale12Si piston alloys were provided by Dong Yang Piston Co. (Ansan, Republic of Korea), and their chemical compositions were measured by ARL3460 optical emission spectroscopy and iCAP6500 inductively coupled plasma-optical emission spectroscopy (see Table 1). Note that in addition to different concentrations of transition elements (Cu, Ni, Fe, and Mn), the ingots all contain phosphorus (20e34 ppm), which is necessary to refine primary Si because AlP inclusions act as heterogeneous nucleation sites [18e21]. To examine the role of AlP in refinement of the primary Si particles, Ale25Si binary alloy with and without 0.05 wt% P was also prepared. 2.2. Casting and heat treatment Ale12Si piston alloys and Ale25Si binary alloy were remelted at 750 C and 850 C in an electric resistance furnace, and then degassed by Ar gas bubbling filtration (GBF) during 20 min. The degassed melts, each weighing approximately 1 kg, were then poured into a copper book mold (245 mm 70 mm 200 mm) that had been preheated to 200 C. Considering the formation temperatures of primary Si, the melt pouring temperatures of Ale12Si and Ale25Si alloys were determined as 700 C and 800 C, respectively. The cooling rate of solidifying melts from 650 C to the eutectic arrest temperature (~560 C) was ~7.5 C/s [15]. A titanium sonotrode preheated to 200 C was injected into each 1 kg melt to introduce ultrasonic waves with an amplitude and frequency of 20 mm and 19 kHz, respectively. The UST was performed for 1 min within a temperature range of 750 to 700 C for Ale12Si piston alloys and a temperature range of 850 to 800 C for Ale25Si binary alloy, after which the melts were poured into the same mold. The details about the ultrasonic process can be found in a previous article [15]. The ingots were solution-treated at 490 C for 2 h, and then aged at 230 C for 5 h (i.e., T7 heat treatment).
2.3. Microstructural observations Samples for microstructural observation were taken from each ingot at a position one half of its length and width, and one quarter of its height. The microstructure was observed using an optical microscope (OM, Nikon, MA200) and a scanning electron microscope (SEM, JEOL, JSM-6610LV) equipped with energy diverse X-ray spectroscopy (EDXS, JEOL, INCA Energy). An image analyzer (IMT, iSolution) was used to quantitatively measure the size (maximum length) and area fraction of the secondary phases. Three OM images (200 magnification) were used for analyzing primary Si particles, while ten OM images (1000 magnification) were used for analyzing eutectic Si and IMC particles. The grain structure was examined using an electron backscatter diffraction (EBSD) instrument installed in the SEM (TESCAN, CZ/MIRA I LMH) with a step size of 10 mm. The fcc-Al grains with sizes greater than 50 mm were averaged to separate them from primary Si particles having a diamond cubic structure. The precipitates were observed through a transmission electron microscope (TEM, JEOL, JEM 2100F). The cross-sectional TEM samples were fabricated by using a focused ion beam and they were placed on Mo grids. To estimate the amount of precipitates formed during T7 heat treatment, the electrical resistivity of as-cast and T7 treated alloys was measured at 25 ± 1 C using a sourcemeter (Keithley, Model 6220) and a voltmeter (Keithley, Model 2182), as outlined in ASTM F76e08 [22]. 2.4. Tensile tests Room-temperature tensile tests of T7 treated alloys were performed using an Instron 4206 universal testing machine with a crosshead speed of 1.5 mm/min, as outlined in ASTM E8/E8Me13a [23]. Tensile tests were also performed at 350 C in accordance with ASTM E21e09 [24] after isothermally maintaining each sample at 350 C for 100 h to simulate the piston operating condition. The crosshead speeds during high-temperature tensile tests were 0.125 and 1.5 mm/min before and after the yield point. Dogbone-shaped (gage section: Ø6 25 mm) samples were used for tensile tests. 3. Results 3.1. Thermodynamic calculation Fig. 1 shows the plot of temperature versus solid mole fraction (fs) during Scheil-Gulliver solidification of Alloy 1, which was calculated using Thermo-Calc software [25] with the TCAL3 database. A minor alloying element, Ti, was excluded in the calculation. Primary Si and ε-Al3Ni phases start to form at 577.3 C and 568.9 C, respectively. Then, complex reactions take place at temperatures ranging from 560.0 to 509.8 C, resulting in the formation of eutectic Si, various IMCs (e.g., ε-Al3Ni, d-Al3CuNi, g-Al7Cu4Ni, QAl5Cu2Mg8Si6, q-Al2Cu, M-Mg2Si, etc.) and an fcc-Al matrix. From the thermodynamic calculations, it is expected that various IMCs
Table 1 Chemical composition of Ale12Si piston alloys and their mole fraction of intermetallic compounds. Alloy
1 2 3 4 5
Chemical composition (wt.%)
Mole fraction of intermetallic compound (%)
Al
Si
Cu
Ni
Mg
Fe
Mn
Ti
P
Transition element
Non-equilibrium solidification
Equilibrium at 350 C
Bal. Bal. Bal. Bal. Bal.
12.08 11.90 12.15 12.81 11.97
3.24 3.22 3.32 4.17 3.79
2.36 2.34 2.51 2.26 3.39
0.81 0.97 0.82 0.86 0.82
0.24 0.41 0.24 0.43 0.50
0.20 0.39 0.59 0.35 0.18
0.08 0.09 0.08 0.09 0.10
0.0021 0.0020 0.0022 0.0034 0.0023
6.12 6.45 6.74 7.30 7.96
6.23 8.44 9.08 8.07 11.80
10.45 11.91 12.00 12.41 13.92
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Fig. 1. Temperature-solid mole fraction curve during Scheil-Gulliver solidification of Alloy 1. Secondary phases are indicated as: a-Al15(Fe,Mn)4Si2, b-Al9Fe2Si2, g-Al7Cu4Ni, dAl3CuNi, ε-Al3Ni, p-Al8FeMg3Si6, q-Al2Cu, Q-Al5Cu2Mg8Si6 and M-Mg2Si.
will form similarly to other AleSi piston alloys. The initiation temperature (Ti), accumulated solid mole fraction (fs), and constituent phase of Scheil solidification reactions for the five Ale12Si piston alloys are listed in Table 2. The first solidification reaction is different for each alloy, depending on the concentrations of Si and the transition elements. It is primary Si that forms first in Alloy 1 and 4, as these contain relatively low concentrations of transition elements (or a high concentration of Si). In contrast, IMCs (a-Al15(Fe,Mn)4Si2 or ε-Al3Ni) form first in Alloys 2, 3 and 5 due to their higher concentration of transition elements. Note also that the Ti and fs values for the first solidification reaction increase with increasing concentration of Si or transition elements. The mole fraction of IMC formed during solidification and its equilibrium value at 350 C are shown in Table 1, which reveals an overall increase in fraction with increasing concentration of transition elements. The Fe, Mn and Ni effectively increase the mole fraction of IMCs by forming b-Al9Fe2Si2, a-Al15(Fe,Mn)4Si2 and ε-Al3Ni, but Cu is less effective because of the relatively high concentration required to form IMCs such as g-Al7Cu4Ni and q-Al2Cu as well as its higher solubility in the Al matrix. 3.2. Microstructural observations 3.2.1. As-cast alloy Fig. 2 shows the effect of UST on the microstructure of as-cast Alloys 1, 3 and 5, in which transition elements amount to 6.12, 6.74 and 7.96 wt%, respectively. The as-cast alloys without UST contain dendritic microstructures, whereas equiaxed
microstructures are prevalent in the as-cast alloys with UST. No pores were observed in any of the samples because the melts were sufficiently degassed prior to casting by means of Ar GBF or UST. Significantly, primary Si particles are severely segregated in those alloys without UST, whereas a homogenous distribution of smaller primary Si particles is seen in the alloys with UST. In either case, most of the eutectic Si and IMCs are observed along the cell boundaries. Fig. 3 shows SEM images revealing different sized IMCs with different morphologies in as-cast alloy. The type of IMCs was identified based on their measured chemical compositions, as listed in Table 3. All samples contained eutectic IMCs such as ε-Al3Ni, dAl3CuNi, g-Al7Cu4Ni, Q-Al5Cu2Mg8Si6, M-Mg2Si and q-Al2Cu. In addition to this, coarse primary a-Al15(Fe,Mn)4Si2 was observed in Alloys 2 and 3 (Fig. 3(b) and (e)), while primary ε-Al3Ni was observed in Alloy 5 (Fig. 3(c) and (f)). These primary IMCs present in Alloys 2, 3 and 5 are considered to form first during solidification, as predicted by thermodynamic calculation results (Table 2). It is clearly seen in Fig. 3 that all of the IMCs decreased in size with the application of UST, regardless of the transition element content. Moreover, the agglomeration of needle-like ε-Al3Ni and d-Al3CuNi, which was frequently observed in the alloys without UST, was rarely observed in alloys with UST. Further, the flake-like Ti(Al,Si)3 particles were observed in the alloys with UST. This is attributed to the increased Ti concentration (see Table 4) caused by dissolution of Ti from the ultrasonic horn into the melt during UST [26]. The effect of UST on the grain structure was examined by EBSD analysis. Fig. 4 shows the EBSD microstructure of the fcc a-Al matrix
Table 2 Initiation temperature (Ti), accumulated solid mole fraction (fs), and constituent phase of Scheil solidification reaction for Ale12Si piston alloys. Alloy
1 2 3 4 5
1st reaction
2nd reaction
Ti ( C)
fs (%)
Phase constituent
577.3 578.3 588.8 589.1 589.3
0.58 0.20 0.36 1.00 0.93
L L L L L
þ þ þ þ þ
Si
a a Si ε
3rd reaction
Ti ( C)
fs (%)
Phase constituent
568.9 572.2 578.3 574.7 579.7
1.41 0.47 0.98 2.42 2.88
L L L L L
þ þ þ þ þ
Si þ ε a þ Si a þ Si Si þ a ε þ Si
4th reaction
Ti ( C)
fs (%)
Phase constituent
563.1 569.9 572.8 562.1 567.6
2.60 1.90 2.83 2.69 4.05
L L L L L
þ þ þ þ þ
Si þ ε þ a-Al a þ Si þ ε a þ Si þ ε Si þ a þ ε ε þ Si þ b
Ti ( C)
fs (%)
Phase constituent
560.0 562.0 562.6 560.5 561.1
49.23 36.52 63.30 22.12 15.16
L L L L L
þ þ þ þ þ
Si þ ε þ a þ a-Al a þ Si þ ε þ a-Al a þ Si þ ε þ a-Al Si þ a þ ε þ a-Al ε þ Si þ b þ a-Al
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Fig. 2. OM images of as-cast alloys (aec) without and (def) with UST: (a,d) Alloy 1, (b,e) Alloy 3, (c,f) Alloy 5.
in the as-cast alloys. The effect of UST on grain structure shows a contradiction between the alloys; the application of UST refined the grain structure of Alloy 1, whereas it slightly coarsened the grain structure of Alloy 4. The effect of UST on the size and area fraction of primary Si, eutectic Si, and primary and eutectic IMCs and grain of as-cast alloy was quantitatively evaluated, and the results are summarized in Table 5. In the absence of UST, the size of primary Si in Alloy 1 (15.8 mm) is smaller than those in Alloy 2 (18.7 mm) and Alloy 3 (19.3 mm), whereas it is larger than that in Alloy 5 (14.5 mm). It is thought that the formation of Si-containing a-Al15(Fe,Mn)4Si2 in Alloys 2 and 3 inhibits the nucleation of primary Si by forming a Sidepleted zone near a-Al15(Fe,Mn)4Si2 particles. In contrast, the formation of ε-Al3Ni in Alloy 5 might accelerate the nucleation of primary Si by forming a Si-enriched zone near ε-Al3Ni particles. Moreover, the larger primary Si size for Alloy 4 (21.5 mm) is due to its higher Si concentration (12.8 wt%) in comparison to those (11.9e12.1 wt%) of other alloys. Table 5 also shows that UST reduces the size and increases the area fraction of primary Si particles. In addition, UST reduces the size of eutectic Si in as-cast alloys except in Alloy 3, which has fine Chinese-script eutectic Si particles in the absence of UST (Fig. 2(b)). The area fraction of eutectic Si is reduced in UST alloys through an increase in the area fraction of primary Si. The UST also decreases
the size of IMCs in UST alloys. However, it would seem that UST does not significantly change the total area fraction of IMCs, although it causes the formation of some Ti(Al,Si)3 particles. Meanwhile, there was a significant decrease in the average grain size with UST except for the Alloy 4. 3.2.2. Heat-treated alloy Fig. 5(a) and (b) present OM images of as-cast and T7 treated Alloy 5 without UST. The constituent phases (i.e., primary Si, eutectic Si and IMCs) are the same in both specimens. Notably, however, the interfaces between the primary Si and Al matrix, as well as between eutectic Si and the Al matrix, are smoother in the T7 alloy. The T7 alloys also show evidence of a breakup of eutectic Si particles that is caused by their spheroidization during solution treatment at 490 C. In addition, some IMCs (e.g. Mg2Si, QAl5Cu2Mg8Si6, g-Al7Cu4Ni and Al2Cu), which are thermally unstable at 490 C were partially dissolved during solution treatment. The artificial aging at 230 C results in the formation of fine precipitates within the Al matrix, as shown in Fig. 5(c). The needlelike particles are q'-Al2Cu and relatively thicker ones are q-Al2Cu, which are confirmed by analyses of their selected area diffraction patterns taken along the [110]Al. Needle-like q'-Al2Cu particles generate streaks along the [200]Al [27] and q-Al2Cu particles are formed with a crystallographic orientation of [110]Al//[531]q.
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Fig. 3. SEM images of as-cast alloys (aec) without and (def) with UST: (a,d) Alloy 1, (b,e) Alloy 3, (c,f) Alloy 5.
Table 3 Chemical composition of intermetallic compounds observed in Ale12Si piston alloys. No.
1 2 3 4 5 6 7 8
Chemical composition (at. %)
Suggested phase
Mg
Al
Si
Fe
Mn
Ni
Cu
Ti
e 59.4 0.2 e 33.3 0.6 e e
67.4 7.5 73.3 68.1 28.6 77.4 65.1 42.1
0.9 33.1 3.7 12.1 30.4 3.4 e 35.0
0.1 e 1.5 3.8 e 0.9 e e
e e 0.6 12.7 e 3.3 e e
0.5 e 10.2 3.3 0.2 12.7 9.7 e
23.1 e 10.8 e 7.5 1.8 25.1 e
e e e e e e e 17.0
Table 4 Ti concentrations of Ale12Si piston alloys (wt%). Alloy
w/o UST
w/UST
1 2 3 4 5
0.08 0.09 0.08 0.09 0.10
0.21 0.24 0.24 0.15 0.26
q-Al2Cu M-Mg2Si d-Al3CuNi a-Al15(Mn,Fe)4Si2 Q-Al5Cu2Mg8Si6 ε-Al3Ni g-Al7Cu4Ni Ti(Al,Si)3
The decreased electrical resistivity values of the T7 samples in comparison to those of the as-cast samples (Fig. 6) also indicate that the precipitation occurs during aging at 230 C [28]. The difference in electrical resistivity values (Dr) between as-cast and T7 alloys increases with increasing transition element content. This means that the amount of precipitates (e.g., q'-Al2Cu) formed during aging at 230 C is larger in highly alloyed system. However, these precipitates formed during aging are known to have an insignificant effect on strength at 350 C [29]. This is because most of q'-Al2Cu particles were already transformed to coarse q-Al2Cu particles
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Fig. 4. EBSD microstructures of as-cast (a,b) Alloy 1 and (c,d) Alloy 4 (a,c) without and (b,d) with UST.
Table 5 Effect of UST on microstructural characteristics of as-cast Ale12Si piston alloys. Alloy
UST
Primary Si Area fraction (%)
1 2 3 4 5
w/o w/ w/o w/ w/o w/ w/o w/ w/o w/
2.7 2.5 1.1 2.1 1.5 5.4 2.1 4.6 1.4 3.8
± ± ± ± ± ± ± ± ± ±
0.1 0.2 1.1 0.7 0.5 0.6 0.9 1.1 0.1 0.7
Eutectic Si Size (mm) 15.8 10.8 18.7 12.5 19.3 15.7 21.5 13.4 14.5 13.5
± ± ± ± ± ± ± ± ± ±
2.1 1.3 2.2 0.3 1.2 0.9 1.0 2.1 2.5 1.0
Primary and eutectic IMCs
Refining efficiency (%)
Area fraction (%)
Size (mm)
Refining efficiency (%)
Area fraction (%)
31.6
7.9 ± 0.6 6.0 ± 0.6 8.5 ± 1.9 6.4 ± 1.0 10.8 ± 2.5 9.3 ± 1.1 7.2 ± 1.6 6.3 ± 0.9 8.1 ± 2.0 6.0 ± 0.4
10.3 ± 1.0 9.2 ± 0.8 10.9 ± 1.0 9.6 ± 1.4 3.3 ± 0.3 8.1 ± 1.7 12.8 ± 2.0 10.4 ± 0.7 9.2 ± 0.9 8.3 ± 0.9
10.7
6.3 5.5 5.9 6.9 6.5 6.5 7.8 7.3 8.3 8.7
33.2 18.7 37.7 6.9
11.9 e 18.8 9.8
± ± ± ± ± ± ± ± ± ±
0.7 0.9 2.1 1.2 2.2 1.7 1.6 1.2 2.4 1.1
Size (mm)
Refining efficiency (%)
8.7 ± 0.9 6.9 ± 0.1 9.0 ± 0.9 7.7 ± 0.3 10.9 ± 3.0 9.8 ± 0.4 9.1 ± 0.8 7.2 ± 0.3 9.5 ± 1.8 7.6 ± 0.8
20.7 14.4 10.1 20.9 20.0
Grain size (mm)
577 241 318 235 598 272 264 362 313 166
± ± ± ± ± ± ± ± ± ±
252 87 147 86 320 95 97 143 135 62
during isothermal holding at 350 C for 100 h before the tensile test, as shown in Fig. 5(d).
Alloy 5 (7.96 wt%), however, the decrease in dimple size from 22.8 to 14.9 mm with UST results in improved ductility (Fig. 7(c) and (d)).
3.3. Tensile properties
4. Discussion
Table 6 lists the tensile strength and elongation at room and elevated temperatures for the T7 alloys. It is clear that both the tensile strength and elongation seem to decrease with increasing transition element content due to the increased area fraction and size of brittle IMCs, which make the alloy vulnerable to tensile cracking at room temperature [13,15]. At the elevated temperature of 350 C, it can be seen that the yield strength and tensile strength increase with the transition element content due to the increase in the IMC fraction. In contrast, the elongation decreases with increasing transition element content. For all alloys, UST improves the tensile strength at room and elevated temperatures. However, the effect of UST on elongation is different depending on the amount of transition elements. The UST increases the elongation of Alloys 3, 4 and 5, whereas it deteriorates the elongation of Alloys 1 and 2 despite refinement of their microstructures. The contradictory effects of UST on the elongation as a function of transition element content are also visible in the tensile-fractured surfaces at 350 C. In Alloy 1 with fewer transition elements (6.12 wt%), the dimple size increased with UST from 14.6 to 16.1 mm, which is in agreement with the observed deterioration in ductility (Fig. 7(a) and (b)). In heavily alloyed systems such as
4.1. Transition elements combined with UST effects on solidification microstructure 4.1.1. Refining mechanism of UST The UST decreased the sizes of secondary phases as shown in Table 5. The refinement of secondary phases by UST is generally explained by cavitation-induced dendrite fragmentation and cavitation-induced heterogeneous nucleation [30e33]. Since the UST was terminated above the liquidus temperature in this study, dendrite fragmentation is not responsible for the refinement of secondary phases [15]. Although nucleation possibly occurs near the cold sonotrode during UST, this effect on the microstructural refinement is not significant, according to Qian et al. [33]. Thus, the cavitation-induced heterogeneous nucleation mechanism is considered to be responsible for the refinement of phases. The cavitation-induced heterogeneous nucleation mechanism can be further divided into three different subcategories [33]. The first is the pressure pulseemelting point mechanism based on the Clapeyron equation. The collapse of bubbles increases the pressure on the melt, and thus, increases the melting point. The increased melting point is equivalent to increased undercooling, leading to
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Fig. 5. OM images of (a) as-cast and (b) T7 Alloy 5 without UST. TEM images of T7 Alloy 5 (c) before and (d) after isothermal holding at 350 C for high-temperature tensile tests.
Fig. 6. Electrical resistivity values of as-cast and T7 alloys without UST.
enhanced nucleation. The second mechanism is related to the fact that the adiabatic expansion of cavitation bubbles undercools the bubbleeliquid interfaces, enhancing the nucleation on the bubble surfaces. The third mechanism is the cavitation-enhanced wetting of inclusions. Non-metallic inclusions (e.g., MgAl2O4 [34], AlP [20]) exist in the Al melt containing Mg and P, and their wettability is improved by the formation of bubbles through the degassing of hydrogen gas trapped at the surface of inclusions. The wetted inclusions can act as effective nucleation sites during the casting. Among these mechanisms, the first may not be prevalent here because the change in pressure due to bubble collapse disappears immediately after UST is terminated. This is supported by the fact that the fraction of primary Si was increased by the application of UST (Table 5); the activation of a pressure pulseemelting point mechanism is known to decrease the fraction of primary Si by shifting the AleSi eutectic point toward the high Si region [16]. It is thought that the wettable inclusions such as MgAl2O4 (cubic; a ¼ 8.08 Å [34]) and AlP (cubic; a ¼ 5.42 Å [20]) might remain wet in the melt even after ultrasonic treatment, thereby acting as nucleation sites for primary Si during casting. Wang et al. [35] and
Table 6 Effect of UST and transition elements on tensile properties at room and elevated temperatures of T7 piston alloys. Alloy
UST
25 C
350 C
Tensile strength (MPa) 1 2 3 4 5
w/o w/ w/o w/ w/o w/ w/o w/ w/o w/
281 296 298 305 243 309 242 285 248 285
± ± ± ± ± ± ± ± ± ±
27 22 24 17 31 23 23 8 10 7
Elongation (%) 0.71 0.65 0.89 0.78 0.49 0.95 0.39 0.42 0.41 0.53
± ± ± ± ± ± ± ± ± ±
0.26 0.15 0.23 0.17 0.15 0.24 0.07 0.03 0.05 0.04
Elongation (%)
Yield strength (MPa) 38.5 39.7 39.8 39.8 41.5 43.7 41.0 42.0 45.8 45.3
± ± ± ± ± ± ± ± ± ±
0.8 1.2 1.0 3.1 2.6 0.9 1.6 0.8 3.0 4.1
Tensile strength (MPa) 63.0 65.5 65.5 68.5 67.0 68.1 67.7 68.3 71.9 76.1
± ± ± ± ± ± ± ± ± ±
0.1 1.4 2.0 0.4 0.9 0.6 0.7 1.3 2.2 2.3
21.2 ± 4.0 17.9 ± 1.2 16.6 ± 1.2 13.8 ± 1.9 9.9 ± 2.0 11.1 ± 4.7 3.7 ± 1.2 12.4 ± 4.5 6.9 ± 3.3 9.1 ± 2.3
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Fig. 7. Tensile-fractured surfaces of (a,b) Alloy 1 and (c,d) Alloy 5 (a,c) without and (b,d) with UST. Tensile tests were performed at 350 C.
Komarov et al. [36] also suggested that the cavitation-induced nucleation via improved wetting of inclusions plays the dominant role in the refinement of primary phases when UST is performed above the liquidus temperature. The roles of oxide and AlP in refining primary Si particles under the UST were investigated further using the Ale25Si binary alloy. It is seen from Fig. 8 that the UST decreased the size of primary Si in the Ale25Si binary alloy with and without AlP inclusions; the size of primary Si was decreased from 38.8 to 27.5 mm and from 143 to 90 mm with and without P addition, respectively. The results suggest that the wetted oxides can act as nucleation sites for the primary Si particles despite different lattice constants between oxide (e.g., MgAl2O4, cubic; a ¼ 8.08 Å [34]) and Si (cubic; 5.43 Å [20]). In a similar fashion, the wetted oxides and AlP may enhance the nucleation of primary IMCs such as a-Al15(Fe,Mn)4Si2 (cubic; a ¼ ~12.68 Å [37]) and ε-Al3Ni (orthorhombic; a ¼ 6.60 Å, b ¼ 7.35 Å, c ¼ 4.80 Å [38]). Meanwhile, it is unclear that the second mechanism (i.e., enhanced nucleation near bubble surfaces) is active in this system. In addition to cavitation-induced heterogeneous nucleation, acoustic streaming that homogenizes the alloying elements and distributes the wetted inclusions over the whole melt also plays a significant role in refining the primary phases. 4.1.2. Competitive refinement of primary phase It is worth noting that the efficiency of primary Si refinement by UST (i.e., the decrease in primary Si size) was different in each alloy as shown in Table 5: Alloy 1 (31.6%), Alloy 2 (33.2%), Alloy 3 (18.7%), Alloy 4 (37.7%) and Alloy 5 (6.9%). Fig. 9(a) shows the relationship between refining efficiency of primary Si and the calculated mole fraction of the primary Si (fs of primary Si for the first reaction in Table 2). For a better understanding, the experimental results obtained in Ale(12,15,18)Si piston alloys [15] and Ale25Si binary alloys are also plotted. It is evident that the refining efficiency of primary Si is greatly increased with increasing primary Si fraction up to 1%, but becomes saturated at approximately 35%. Conversely, the refining efficiency of primary Si decreases with
increasing mole fraction of primary IMC (Fig. 9(b)), which means that any primary IMC formed prior to primary Si consumes nucleation sites. This is supported by the drastic refinement of primary IMC in Alloy 3 and Alloy 5 with poor refining efficiency of primary Si; the sizes of primary a-Al15(Fe,Mn)4Si2 in Alloy 3 and primary ε-Al3Ni in Alloy 5 decreased significantly from 36 to 23 mm and from 50 to 17 mm, respectively. The consumption of nucleation sites by primary IMC limits the opportunity for primary Si nucleation and, consequently, reduces the efficiency of primary Si refinement by UST. This implies that AleSi piston alloys require a sufficient amount (>1%) of pre-existing primary Si particles prior to IMC formation to achieve the maximum refining efficiency using ultrasonic irradiation. Consequently, it is necessary to optimize the amounts of added Si and transition elements in relation to the formation sequence and mole fraction of the primary phase (i.e., Si or IMC). 4.1.3. Eutectic phase and grain The refining efficiency of eutectic Si by UST ranges from 9.8 to 18.8% as shown in Table 5. It is thought that the refining efficiency of eutectic Si was not affected by the type and mole fraction of the primary phase because the eutectic Si phases are formed after the ultrasonically introduced nucleation sites are nearly consumed by the formation of primary phases. Also, the efficiency of IMC refinement (10.1e20.9%) was not influenced by the primary phase, because most of the IMCs are eutectic phases. According to Das and Kotadia [39,40], the refining efficiency of eutectic phases can be improved by the application of UST during solidification. They showed that the UST performed during solidification significantly decreased the size of eutectic Si particles in hypoeutectic Ale10Si [39] and Ale7Si [40] alloys, thereby achieving greater refining efficiency (>70%) of the eutectic Si. They suggested that such a significant refinement of eutectic Si was attributed to the intense convection from cavitation near the sonotrode that homogenizes the solutes at the eutectic growth front [40]. It is well known that both TiAl3 particles and the solute Ti atoms cause grain refinement by providing potential nucleation sites for
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Fig. 8. OM images of as-cast (a,b) Ale25Si and (c,d) Ale25Sie0.05P alloys (a,c) without and (b,d) with UST.
Fig. 9. Effects of mole fraction of (a) primary Si and (b) primary IMC formed during 1st solidification reaction on efficiency of primary Si refinement in UST alloys. Relationships observed in Ale(12,15,18)Si piston alloys [15] and Ale25Si binary alloys are plotted together.
a-Al and restrict the growth of a-Al [41,42]. In this study, the UST using the Ti sonotrode increased the Ti concentration from 0.08e0.10 to 0.15e0.26 wt% (Table 4). Taking into account a peritectic composition in an AleTi binary alloy (~0.15 wt%) [41], UST alloys probably contain numerous TiAl3 and/or Ti(Al,Si)3 particles and a large amount of solute Ti atoms. Thus, the effect of Ti concentration on the grain size of as-cast alloys should be considered along with the UST effect on grain size. Fig. 10 shows the effect of Ti concentration on the grain size of as-cast Ale12Si piston alloys. The grain size gradually decreases with increasing Ti concentration. Overall, the UST alloys having higher Ti concentrations exhibit smaller grain sizes than the alloys without UST. The grain size of Alloy 4 was slightly increased by application of UST. However, it is not reasonable to conclude that the UST causes the grain coarsening in Alloy 4 considering the large standard deviation values. This analysis strongly suggests that the grain size of as-cast alloys depends on the Ti concentration rather than on the application of UST [15,43].
Fig. 10. Effect of Ti content on the grain size of as-cast alloys without and with UST.
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Fig. 11. Effect of IMC fraction on tensile strength at (a) 25 C and (b) 350 C.
4.2. Transition elements combined with UST effects on tensile properties Fig. 11(a) shows the relationship between the measured IMC fraction and tensile strength at 25 C. In the absence of UST, the room-temperature tensile strength decreases with increasing IMC fraction. In the presence of UST, however, the room-temperature tensile strength remains at ~300 MPa until the IMC fraction reaches 7%, and then decreases with further IMC fraction. This implies that the refinement of IMCs by UST allows the piston alloy to contain more IMCs without deterioration of strength. The increased IMC fraction leads to an improvement in high-temperature tensile strength as shown in Fig. 11(b). Furthermore, the UST increases the slope of the plot of tensile strength versus the IMC fraction from 2.6 to 3.2 MPa/%. This is likely due to the refinement of primary and eutectic phases. It is worth noting that while UST causes a decrease in the elongation of Alloy 1 and 2, it actually increases the elongation of Alloys 3, 4 and 5. The alloy specimens subjected to UST were found to contain coarse flake-like Ti(Al,Si)3 particles that are deleterious to ductility. With those alloys containing fewer transition elements, (i.e., Alloys 1 and 2), the elongation is relatively high due to a smaller IMC fraction, and so these are more significantly affected by deleterious Ti(Al,Si)3 particles. In contrast, alloys with a higher transition element content (i.e., Alloys 3, 4 and 5) experience a beneficial refinement of microstructure with UST that can overcome the deleterious effect of Ti(Al,Si)3 particles.
by UST. The decreased fraction of primary Si indicates that the refinement of primary phases can be attributed to the enhanced nucleation on the cavitation bubble surfaces and the wetted inclusions when UST is used. (3) The application of UST also reduces the sizes of eutectic Si and IMC, but the refining efficiency in this case is not influenced by the type or fraction of the primary phase. The formation of coarse flake-like Ti(Al,Si)3 particles caused by the dissolution of Ti elements from the titanium ultrasonic horn can also occur, resulting in the grain refinement. (4) At room temperature, both the tensile strength and ductility simultaneously decrease with increasing transition element content. At 350 C, however, the yield strength and tensile strength increase with increasing transition element content at the expense of ductility. Regardless of transition element content, the UST improves the tensile strength at room and elevated temperatures, which is mainly attributed to a refinement of the microstructure. (5) There is also an improvement in ductility at elevated temperatures with UST when the transition element content is relatively high, implying that the beneficial effect of UST is greater in heavily alloyed systems. The deterioration in ductility that occurs in alloys with lower transition element content is most likely because the deleterious effect of coarse flake Ti(Al,Si)3 is greater than the beneficial effect of UST (i.e., microstructure refinement) in these alloys. Acknowledgements
5. Conclusions (1) The microstructure of Ale12Si piston alloys is composed of primary Si, eutectic Si and a variety of IMCs (e.g., ε-Al3Ni, dAl3CuNi, g-Al7Cu4Ni, Q-Al5Cu2Mg8Si6, Mg2Si and q-Al2Cu). The addition of Mn and Ni can change the solidification sequence by forming primary a-Al15(Fe,Mn)4Si2 or ε-Al3Ni prior to primary Si formation. Both the formation temperature and fraction of primary phase (Si or IMC) increase with increasing Si or transition element content. (2) The UST significantly decreases the size of primary Si and primary IMCs. The refinement efficiency of primary Si increases with an increase in primary Si fraction or decrease in primary IMC fraction formed during the first solidification reaction. The consumption of nucleation sites by primary IMC limits the opportunity for primary Si nucleation and, consequently, reduces the efficiency of primary Si refinement
The authors would like to acknowledge the financial support received from the R&D Convergence Program of the MSIP (Ministry of Science, ICT and Future Planning) and NST (National Research Council of Science and Technology) of the Republic of Korea (Grant CCPe13e17eKIMS), and from the Main Research Program of the Korea Institute of Materials Science (PNK4711). References [1] J.R. Davies, Aluminum and Aluminum Alloys, ASM International, OH, 1993. [2] N. Belov, D. Eskin, N. Avxentieva, Constituent phase diagrams of the AleCueFeeMgeNieSi system and their application to the analysis of aluminum piston alloys, Acta Mater 53 (2005) 4709e4722. [3] R. Gholizadeh, S.G. Shabestari, Investigation of the effects of Ni, Fe, and Mn on the formation of complex intermetallic compounds in AleSieCueMgeNi alloys, Metall, Mater. Trans. A 42A (2011) 3447e3458. [4] A.J. Moffat, S. Barnes, B.G. Mellor, P.A.S. Reed, The effect of silicon content on long crack fatigue behavior of aluminum-silicon piston alloys at elevated
J.-G. Jung et al. / Journal of Alloys and Compounds 712 (2017) 277e287 temperature, Inter. J. Fatigue 27 (2005) 1564e1570. [5] Z. Qian, X. Liu, D. Zhao, G. Zhang, Effects of trace Mn addition on the elevated temperature tensile strength and microstructure of a low-iron AleSi piston alloy, Mater. Lett. 62 (2008) 2146e2149. [6] Y. Li, Y. Wu, Z. Qian, X. Liu, Effect of co-addition of RE, Fe and Mn on the microstructure and performance of A390 alloy, Mater. Sci. Eng. A 527 (2009) 146e149. [7] Y. Li, Y. Yang, Y. Wu, L. Wang, X. Liu, Quantitative comparison of three Niecontaining phases to the elevated-temperature properties of AleSi piston alloys, Mater. Sci. Eng. A 527 (2010) 7132e7137. [8] Y. Yang, K. Yu, Y. Li, D. Zhao, X. Liu, Evolution of nickel-rich phases in AleSieCueNieMg piston alloys with different Cu additions, Mater. Des. 33 (2012) 220e225. [9] Y. Li, Y. Yang, Y. Wu, Z. Wei, X. Liu, Supportive strengthening role of Crerich phase on AleSi multicomponent piston alloy at elevated temperature, Mater. Sci. Eng. A 528 (2011) 4427e4430. [10] C.-Y. Jeong, Effect of alloying elements on high temperature mechanical properties for piston alloy, Mater. Trans. 53 (2012) 234e239. [11] M.M. Haque, M.A. Maleque, Effect of process variables on structure and properties of aluminiumesilicon piston alloy, J. Mater. Proc. Tech. 77 (1998) 122e128. [12] S. Manasijevi c, S. Markovi c, Z. A cimovi c-Pavlovi c, K. Rai c, R. Radisa, Effect of heat treatment on the microstructure and mechanical properties of piston alloys, Mater Tech. 47 (2013) 585e591. [13] J.-G. Jung, S.-H. Lee, J.-M. Lee, Y.-H. Cho, S.-H. Kim, W.-H. Yoon, Improved mechanical properties of near-eutectic AleSi piston alloy through ultrasonic melt treatment, Mater. Sci. Eng. A 669 (2016) 187e195. [14] W. Khalifa, S. El-Hadad, Y. Tsunekawa, Microstructure and wear behavior of solidification sonoprocessed B390 hypereutectic AleSi alloy, Metall, Mater. Trans. A 44A (2013) 5817e5824. [15] J.-G. Jung, J.-M. Lee, Y.-H. Cho, W.-H. Yoon, Combined effects of ultrasonic melt treatment, Si addition and solution treatment on the microstructure and tensile properties of multicomponent AleSi alloys, J. Alloys Compd. 693 (2017) 201e210. [16] M. Sha, S. Wu, L. Wan, Combined effects of cobalt addition and ultrasonic vibration on microstructure and mechanical properties of hypereutectic AleSi alloys with 0.7%Fe, Mater. Sci. Eng. A 554 (2012) 142e148. } , P. An, L. Wan, Effects of ultrasonic vibration and manga[17] C. Lin, S. Wu, S. Lu nese on microstructure and mechanical properties of hypereutectic AleSi alloys with 2%Fe, Intermetallics 32 (2013) 176e183. [18] P.B. Crosley, L.F. Mondolfo, The modification of aluminum-silicon alloys, AFS Trans. 74 (1966) 53e64. [19] C.R. Ho, B. Cantor, Heterogeneous nucleation of solidification of Si in AleSi and AleSieP alloys, Acta Metall. Mater. 43 (1995) 3231e3246. [20] K. Nogita, S.D. McDonald, K. Tsujimoto, K. Yasuda, A.K. Dahle, Aluminium phosphide eutectic grain nucleus in hypoeutectic AleSi alloys, J. Electron Microsc. 53 (2004) 361e369. [21] J. Li, F.S. Hage, X. Liu, Q. Ramasse, P. Schumacher, Revealing heterogeneous nucleation of primary Si and eutectic Si by AlP in hypereutectic AleSi alloys, Sci. Rep. 6 (2016), 25244-1e8. [22] ASTM F76-08, Standard Test Methods for Measuring Resistivity and Hall Coefficient and Determining Hall Mobility in Singleecrystal Semiconductors, ASTM International, PA, 2008.
287
[23] ASTM E8/E8Me13a, Standard Test Methods for Tension Testing of Metallic Materials, ASTM International, PA, 2013. [24] ASTM E21e09, Standard Test Methods for Elevated Temperature Tension Tests of Metallic Materials, ASTM International, PA, 2009. [25] N. Saunders, A.P. Miodownik, CALPHAD, Elsevier, Oxford, 1998. [26] F. Dong, X. Li, L. Zhang, L. Ma, R. Li, Cavitation erosion mechanism of titanium alloy radiation rods in aluminum melt, Ultrason. Sonochem. 31 (2016) 150e156. [27] A.K. Shukla, W.A. Baeslack III, Study of microstructural evolution in frictionstir welded thin-sheet AleCueLi alloy using transmission-electron microscopy, Scr. Mater 56 (2007) 513e516. [28] J.-G. Jung, J.-S. Park, Y.-K. Lee, Quantitative analyses of dissolution and precipitation kinetics of q-Al2Cu phase in an Ale6.2Sie2.9Cu alloy using electrical resistivity, Met. Mater. Int. 19 (2013) 147e152. [29] Z. Asghar, G. Requena, E. Boller, Three-dimensional rigid multiphase networks providing high-temperature strength to cast AlSi10Cu5Ni1-2 piston alloys, Acta Mater 59 (2011) 6420e6432. [30] M.C. Flemings, Solidification Processing, McGraw-Hill, New York, 1974. [31] G.I. Eskin, Ultrasonic Treatment of Light Alloy Melts, Gordon & Breach, Amsterdam, 1998. [32] J.D. Hunt, K.A. Jackson, Nucleation of solid in an undercooled liquid by cavitation, J. Appl. Phys. 37 (1966) 254e257. [33] M. Qian, A. Ramirez, A. Das, Ultrasonic refinement of magnesium by cavitation: clarifying the role of wall crystals, J. Cryst. Growth 311 (2009) 3708e3715. [34] Y. Wang, H.-T. Li, Z. Fan, Oxidation of aluminium alloy melts and inoculation by oxide particles, Trans. Indian Inst. Met. 65 (2012) 653e661. [35] F. Wang, D. Eskin, T. Connolley, J. Mi, Effect of ultrasonic melt treatment on the refinement of primary Al3Ti intermetallic in an Ale0.4Ti alloy, J. Cryst. Growth 435 (2016) 24e30. [36] S. Komarov, Y. Ishiwata, I. Mikhailov, Industrial application of ultrasonic vibrations to improve the structure of AleSi hypereutectic alloys: potential and limitations, Metall. Mater. Trans. A 46A (2015) 2876e2883. [37] A.K. Srivastaba, S. Ranganathan, Quasicrystals, crystals and multiple twins in rapidly solidified AleCreSi, AleMneSi and AleMneCreSi alloys, Acta Mater. 44 (1996) 2935e2946. [38] N.D. Broom, G.J. Davies, Plastic deformation, fracture and dislocation structures in the orthorhombic intermetallic phase Al3Ni, Acta Metall. 23 (1975) 537e546. [39] A. Das, H.R. Kotadia, Effect of high-intensity ultrasonic irradiation on the modification of solidification microstructure in a Si-rich hypoeutectic AleSi alloy, Mater. Chem. Phys. 125 (2011) 853e859. [40] H.R. Kotadia, A. Das, Modification of solidification microstructure in hypo- and hyper-eutectic Al-Si alloys under high-intensity ultrasonic irradiation, J. Alloys Compd. 620 (2015) 1e4. [41] M. Easton, D. StJohn, Grain refinement of aluminum alloys: Part. I. the nucleant and solute paradigmsea review of the literature, Metall. Mater. Trans. A 30A (1999) 1613e1623. [42] Z. Fan, Y. Wang, Y. Zhang, T. Qin, X.R. Zhou, G.E. Thompson, T. Pennycook, T. Hashimoto, Grain refining mechanism in the Al/AleTieB system, Acta. Mater. 84 (2015) 292e304. [43] D. Eskin, Ultrasonic melt processing: opportunities and misconceptions, Mater. Sci. Forum 794e796 (2014) 101e106.