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Microstructural evolution and mechanical properties of Al-7Si-0.3Mg alloys with erbium additions
Accepted Manuscript Microstructural evolution and mechanical properties of Al-7Si-0.3Mg alloys with erbium additions P. Pandee, U. Patakham, C. Limmaneevichitr PII:
S0925-8388(17)33083-9
DOI:
10.1016/j.jallcom.2017.09.054
Reference:
JALCOM 43110
To appear in:
Journal of Alloys and Compounds
Received Date: 4 February 2017 Revised Date:
4 September 2017
Accepted Date: 5 September 2017
Please cite this article as: P. Pandee, U. Patakham, C. Limmaneevichitr, Microstructural evolution and mechanical properties of Al-7Si-0.3Mg alloys with erbium additions, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2017.09.054. 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 proof before it is published in its final 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.
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Microstructural Evolution and Mechanical Properties of Al-7Si-0.3Mg Alloys with Erbium Additions
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P. Pandee1* U. Patakham2 and C. Limmaneevichitr1
Department of Production Engineering, Faculty of Engineering, King Mongkut’s University of Technology Thonburi, 126 Pracha-Utid Rd., Bangmod, Tungkhru, Bangkok 10140, Thailand
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National Metal and Material Technology Center, National Science and Technology Development
*Corresponding author
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Agency, Pathumthani 12120, Thailand
Phromphong Pandee
Telephone: (+66) 82 654 8183, Fax: (+66) 2470 9198
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Email: [email protected]
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ACCEPTED MANUSCRIPT Abstract The mechanical properties of hypoeutectic Al-Si foundry alloys are generally controlled by their microstructures. The present study was performed to determine the influence of erbium (Er) in Al7Si-0.3Mg alloys on the microstructures and mechanical properties. The evolution of microstructure
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was studied by thermal analysis and interrupted solidification technique. Optical microscopy, scanning electron microscope, and energy-dispersive X-ray were employed to characterize the alloy microstructures. The results demonstrated that the addition of Er significantly refined both the
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secondary dendrite arm spacing (SDAS) and the primary α-Al grain. Simultaneously, Er modified eutectic Si phase from coarse acicular plate-like structures into fine fibrous features. We observed
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that Er additions strongly suppress the nucleation of eutectic Si due to the preferential formation of ErP instead of AlP. Er additions change the eutectic growth mode from nucleation and growth on AlP particles to the propagation of a defined eutectic front from the mold walls. Moreover, Er addition that was cast for different holding times indicated that there was no fading effect on the
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SDAS, primary α-Al grain size, and eutectic Si morphology. Importantly, the tensile properties of the Er-modified alloys at as-cast and T6 conditions can be enhanced due to simultaneous refinement
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both the SDAS and the α-Al grains and modification of the eutectic Si. Keyword: Grain refinement, Modification, Secondary dendrite arm spacing, Aluminum-silicon
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alloys, Eutectic solidification, Rare earth
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ACCEPTED MANUSCRIPT 1. Introduction The major microstructural features of commercial hypoeutectic Al-Si foundry alloys are primary αAl dendrites, secondary dendrite arm spacing (SDAS), and Al-Si eutectic structures. Control of these
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microstructures is important in achieving high-quality casting products. It is well-known that reducing SDAS, refining primary α-Al dendrites, and modifying eutectic silicon morphology improves the mechanical properties of hypoeutectic Al-Si foundry alloys. The addition of alloying elements is the main way for obtaining proper microstructure. The addition of alloying elements,
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such as boron, titanium, strontium, or sodium, is one of many methods to achieve grain refinement
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and eutectic silicon modification. However, these elements may only refine the primary α-Al grains or modify the eutectic silicon and cannot do both functions.
In the past several years, many rare earth elements have been studied as micro alloying elements. Rare earth elements can refine primary α-Al grain [1, 2] and/or change the eutectic silicon from a
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coarse plate or flake to a fibrous or lamellar morphology [3-9]. Recently, there has been growing interest in adding rare earth element erbium (Er) into aluminum casting alloys. For example, Zhao et al. [10] revealed that the addition of Er in Al-Zn-Mg-Cu alloys mainly forms an Al3Er, which could
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act as a heterogeneous nucleus to refine grains. Hu et al. [5] reported that Er modified the eutectic silicon from coarse plate-like to fine-fibrous in Al-Si eutectic alloy. Moreover, our previous study
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suggested that the addition of Er to the hypoeutectic Al-Si alloy resulted in modify the eutectic silicon [11]. However, the mechanisms of grain refinement and eutectic silicon modification due to erbium additions are not well understood. The mechanism of grain refinement is usually explained by both the nucleant potency and the solute segregating power [12-14]. It has been suggested that an increase of the nucleant with high potency significantly increases the nucleation rate, resulting in effective grain refinement. In terms of the solute segregating power, it has also been suggested that the contribution of solute is related to the
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ACCEPTED MANUSCRIPT role of solute segregation near the liquid-solid interface in restricting grain growth and constitutional undercooling as a driving-force to further activate nucleation. Several theories, such as restricted nucleation theory and restricted growth theory, were proposed to
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explain the mechanism of chemical modification in hypoeutectic Al-Si alloy. The restricted growth theory is explained by the impurity induced twinning (IIT) mechanism [15, 16] and the twin-plane re-entrant edge (TPRE) mechanism [9, 16-19]. The IIT mechanism proposes that the atoms of the modifier are absorbed on the growth steps of the silicon solid-liquid interface, which causes an
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increase in the twin density of silicon. The TPRE mechanism postulates that a modifier prevents or
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retards the subsequent attachment of silicon atoms during eutectic silicon growth. For restricted nucleation theory, it has been proposed that the AlP particles act as the nucleation site of the eutectic silicon phases [20-23]. The modifying elements can react with AlP nuclei in the melt and alter the number density of eutectic silicon nucleation events, resulting in modification of eutectic silicon. Moreover, chemical modification mechanism is explained with other important factors, such as
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intermetallic clusters and altered surface tension [24-28]. Although considerable research has been devoted to the benefits of erbium additions, microstructure
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evolution during solidification has received considerably less attention. The purpose of this study is to investigate microstructural evolution during solidification to clarify the influence of erbium on the
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hypoeutectic Al-Si alloy and explain the grain refinement and eutectic silicon modification mechanisms as well as the tensile properties. The aim of this research is to compare the chemical grain refinement and chemical modification mechanisms caused by erbium with other previously investigated elements.
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ACCEPTED MANUSCRIPT 2. Experimental Procedure Commercial ingots of Al-7Si-0.3Mg alloy were used in the experiment. The Al-7Si-0.3Mg alloy ingots were manufactured from primary aluminum ingot to minimize the effects of minor element
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contamination, and samples were prepared from the same batch to avoid any undesired variation. Approximately 1 kg of the base alloy (Al-7Si-0.3Mg) was weighed and melted in a silicon carbide crucible at 750 ºC using an induction furnace. Er was added in varying amounts (0.2, 0.4, and 0.6
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wt.%) and the alloys were made using Al-10 wt.% Er master alloy.
Cooling curve analysis was conducted for each alloy in order to determine the characteristic
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temperatures of primary α-Al dendrite and Al-Si eutectic solidification, i.e., nucleation temperature (TN), minimum temperature (TMin), and growth temperature (TG). The cooling rate of the liquid at 650 ºC, just prior to nucleation of primary α-Al phase, was approximately 0.1 ºC/s. Each alloy was then poured at 720 ºC into a graphite cup (outer diameter 50 mm, inner diameter 30 mm, height 50 mm
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and a uniform wall thickness of 10 mm) and a stainless steel cup (outer diameter of bottom 30 mm, outer diameter of upper 34 mm, height 40 mm, and a uniform wall thickness of 0.5 mm). The cooling rate in the graphite cup was approximately 3 ºC/s. The cooling rate in the stainless steel cup was
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approximately 1 ºC/s.
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The fading phenomenon is a key drawback for both grain refinement and eutectic modification in an aluminum foundry alloy. The fading characteristics of Er addition were subsequently evaluated on the SDAS, primary α-Al grain size, and eutectic silicon morphology. The addition of 0.4 wt.% Er was held at 750 ºC for different time intervals. Parts of the melt were poured at different intervals (30, 60, 90, and 120 min) into graphite molds to later observe the resultant microstructure for possible fading of Er. Interrupted solidification experiments were performed to investigate the nucleation behavior of eutectic silicon. A schematic diagram of the experimental set-up is shown in our previous work [29]. 5
ACCEPTED MANUSCRIPT All samples were slowly cooled at approximately 0.1 ºC/s and immediately dropped into an ice-brine quench bath located underneath the holding furnace at a desired time. The fully solidified samples of all experiments were cross-sectioned at 15 mm from the bottom of the
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samples, and the obtained samples were subsequently prepared with standard metallographic procedures. The samples were then etched in Tucker’s reagent (45 ml HCl, 15 ml HNO3, 15 ml HF, 25 ml H2O) to highlight the grain boundaries. The samples were scanned on a high-resolution flatbed scanner to obtain macrographs for the investigation of the grain size. Grain sizes were measured by
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the lineal intercept method with the help of an optical microscope image analyzer system. The SDAS
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values were determined by measuring the SDAS from the corresponding metallography samples using an optical microscope image analyzer system. At least 50 measurements were made for each sample, and the average value of these measurements was used to represent the SDAS value at the corresponding compositions. Moreover, an image analyzer was used to measure three parameters of eutectic silicon particles, i.e., silicon particle length, mean area, and aspect ratio. All three parameters
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were the average values from 20 different images. The microstructures from all samples were recorded by an optical microscope (OM) and JEOL JSM-78000F scanning electron microscope
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(SEM). Intermetallic phases were identified using a Hitachi S-3400N scanning electron microscope equipped with an energy-dispersive X-ray (EDX) detector, EDAX_TSL EBSD hardware and
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OIM_TSL Collection 5.3 software for automated recognition and indexing of EBSD patterns. Castings were also made to prepare tensile test bars according to the ASTM B-108. The mold temperature was preheated at 250 ºC and the pouring temperature is 720 ºC. Tensile testing of as-cast and T6 (solution-treated for 8 h at 540 ºC and aging for 8 h at 155 ºC) samples was carried out at room temperature and a strain rate of 5 mm/min. The present values of tensile properties are the average calculated from five tests for each set of alloy.
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ACCEPTED MANUSCRIPT 3. Results Thermal Analysis An assembly of cooling curves taken from Al-7Si-0.3Mg alloys with and without various Er
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additions is presented in Figure 1. It should be noted that these plots are offset on the time scale to better identify the trend. Three characteristic temperatures are measured from the cooling curves of each alloy, i.e., the nucleation temperature, TN, at which nucleation starts, the minimum temperature,
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TMin, at which nucleation ends, and the growth temperature, TG. The characteristic temperatures of the primary α-Al and Al-Si eutectic solidification obtained from the cooling curves are summarized
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in Table 1.
The nucleation and growth temperatures of primary α-Al dendrite of the Al-7Si-0.3Mg alloy without Er addition are approximately 606.3 ºC and 607.9 ºC. In alloys with Er, the nucleation and growth temperatures of primary α-Al dendrite solidification are slightly increased. At Al-Si eutectic
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solidification, the nucleation temperature was reduced by approximately 5.9 ºC (574.9 ºC - 569.0 ºC) with the addition of 0.2 wt.% Er and later stabilized with an increase in the Er level. Low recalescence undercooling (0.6 ºC) was observed in the Al-7Si-0.3Mg alloy without Er addition.
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Recalescence undercooling was significantly increased with 0.2 wt.% Er addition and then appeared
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to stabilize at approximately 2.0 ºC for higher Er levels. Table 1 shows the growth temperature with similar trends. Upon adding 0.2 wt.% Er, the growth temperature is reduced by approximately 6.1 ºC (573.8 ºC - 567.7 ºC), and it remains constant with further increases in Er content. Primary α-Aluminum Grain
The macrographs of Al-7Si-0.3Mg alloys with various Er additions solidified with low cooling rate (1 ºC/s) are shown in Figures 2(a) to (d) and with high cooling rate (3 ºC/s) are shown in Figures 2(e) to (h). The primary α-Al grain structures of Er-free alloy with both cooling rates are shown in
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ACCEPTED MANUSCRIPT Figures 2(a) and (e). For low cooling rates, large equiaxed grains are observed in samples. For high cooling rates, the sample consists of large equiaxed grains in the center and large columnar grains at the edge. In samples with 0.2 wt.% Er addition, there were slight changes in grain structure with both cooling rates, as shown in Figures 2(b) and (f). However, with an addition of high Er (0.6 wt.%),
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grain refinement was evident, as shown in Figures 2(d) and (h). It is clear that fine equiaxed grains dominated the samples with 0.6 wt.% Er addition. The average grain size and standard deviation measurements with different Er additions are shown in Figure 3. At low cooling rate, adding Er
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significantly reduces the average grain size from 5011 µm to 1319 µm. However, at high cooling
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rate, adding Er slightly reduces the average grain size from 2004 µm to 982 µm. Typical dendritic microstructures of Al-7Si-0.3Mg alloys with various Er additions at a low cooling rate (1 ºC/s) are shown in Figures 4(a) to (d) and a high cooling rate (3 ºC/s) are shown in Figures 4(e) to (h). It can be observed from the micrographs that the addition of Er significantly reduced the SDAS. Figure 5 shows the average SDAS and standard deviation measurements as a function of the
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percentage of Er addition. It can be observed that increasing the levels of Er solute results in a reduction in SDAS to a certain extent. As 0.6 wt.% Er was added to the Al-7Si-0.3Mg alloy, SDAS
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of the Er-free alloy decreased from 79 µm to 63 µm at a low cooling rate and from 32 µm to 18 µm at a high cooling rate. Er exhibits the strongest refining effect on the SDAS for a high cooling rate (3
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ºC/s).
SEM micrograph and EDX point analysis of Al-7Si-0.3Mg alloy with 0.6 wt.% Er addition are presented in Figures 6(a) and (b). The SEM image in Figure 6(a) shows some typical primary phases at the grain boundaries. The intermetallics (as pointed out in Figure 6(a)) were identified by SEMEDX point analysis (Figure 6(b)) to be Al42Si30Er28 in at.%, which is close to Al3Si2Er2 [30]. Al-Si Eutectic Morphology
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ACCEPTED MANUSCRIPT Figure 7 shows Al-Si eutectic structures from an optical microscope taken from samples with different cooling rates and Er additions. The eutectic silicon in the Al-7Si-0.3Mg alloy without Er addition is very coarse and plate-like, as shown in Figure 7(a) (at a low cooling rate) and Figure 7(e) (at a high cooling rate). With the addition of 0.2 wt.% Er (Figures 7(b) and (f)), a decrease in the
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eutectic size was observed, however, the eutectic silicon is still plate-like. In the sample with 0.6 wt.% Er addition, fine fibrous eutectic silicon can be obtained (Figures 7(d) and (h)). The high magnification 3D morphologies of eutectic silicon at high cooling rate (3 ºC/s) with various Er
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additions are depicted in Figure 8. Figure 8(a) show the clustered plate-like eutectic Si occurred in the Er-free alloy. The addition of 0.2 wt.% Er resulted in the refinement of eutectic Si but the
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morphology was still the mixture of plate-like and branched structure, as shown in Figure 8(b). While from Figures 8(c) and (d), it can be seen that the 0.4 and 0.6 wt.% Er modified alloy exhibited a fine fibrous eutectic silicon. The length, mean area, and aspect ratio of eutectic silicon particles of Al-7Si-0.3Mg alloys at different Er additions were determined by an image analyzer and are
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summarized in Table 2. It is clear that all parameters of eutectic silicon decrease with an increase of
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Figure 9 shows the micrographs of Al-7Si-0.3Mg alloys with and without Er additions taken from the quenched samples (at 100 s after the start of Al-Si eutectic solidification). The quenched sample
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without Er addition is shown in Figure 9(a). We found that the eutectic is located at or close to the tips of the primary α-Al dendrites throughout the cross-section of the quenched sample. Polyhedral silicon crystal with a dark particle at the center is shown in Figure 9(c). Our previous work on similar particles [29] has shown that they contain Al, P, and O. A quenched eutectic microstructure of the sample with 0.6 wt.% Er addition is shown in Figure 9(b). When adding Er, Al-Si eutectic is only present adjacent to the mold wall, and there is a eutectic growth front propagating in an opposite direction to the heat flux direction. The typical eutectic growth front of Al-7Si-0.3Mg with 0.6 wt.%
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ACCEPTED MANUSCRIPT Er addition as shown in Figure 9(d). We found a eutectic growth front propagating from bottom to top of the micrograph which corresponds to the direction opposite to the thermal gradient. SEM analysis of Er-containing samples found much fewer P-containing particles. Figure 10 shows a
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typical case in a sample containing 0.6 wt.% Er. The secondary electron (SE) image in Figure 10(a) shows a cubic particle in the sample. EDX point analysis in Figure 10(b) shows that the particle contains Er and P (and negligible Al). In order to confirm that this phase is ErP and Kα peak of Al is overlapping, the crystal structure was determined using EBSD, as shown in Figure 11. Figure 11(a)
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shows the Kikuchi pattern for the ErP phase. The results of the EBSD pattern analysis indicate that
Er-Containing Intermetallic Phases
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the ErP phase has a rock salt structure, as shown in Figure 11(b).
Figure 12 shows SEM analysis of intermetallic in eutectic region in a 0.6 wt.% Er addition sample after natural solidification at 3 ºC/s. The SE image (Figure 12(a)) shows that there are intermetallic
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phases in the eutectic region. EDX mapping such as Figures 12(b) to (f) suggests that intermetallics containing Er. To identify the Er-containing intermetallics, EDX point analysis was performed. Figure 12(g) is EDX point analysis at point A in Figure 12(a) and Figure 12(h) is EDX point analysis
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at point B in Figure 12(a). Figure 12(i) shows EDX results of both points indicating the presence of
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Al, Si, Mg, Er, and Fe for point A and Al, Si, Mg, and Er for point B. Fading Behavior of Erbium
The macrographs and micrographs of samples with 0.4 wt.% Er addition cast after various holding times (30, 60, 90 and 120 min) are presented in Figure 13. It can be clearly observed from Figures 13(a) to (d) that the primary α-Al grain size does not show any significant change with increasing holding time. Moreover, it can be observed from the micrographs of SDAS (Figures 13(e) to (h)),
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ACCEPTED MANUSCRIPT and silicon morphology (Figures 13(i) to (l)) that the holding time has no effect on the SDAS and silicon morphology. It can be concluded that the Er additions do not show any fading effect. Tensile Testing
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The tensile properties of Al-7Si-0.3Mg alloy with various Er additions at as-cast and T6 conditions are listed in Table 3. At the as-cast condition, both the tensile strength and the elongation increased with 0.2 wt.% Er addition and then slightly decreased with an increase of Er content. At the T6
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condition, the tensile strength was increased from 242.6 MPa to 280.3 MPa and elongation was increased from 2.7 % to 3.3 % with 0.2 wt.% Er addition. However, they were decreased when the Er
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content was increased to 0.4 wt.% and 0.6 wt.%.
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ACCEPTED MANUSCRIPT 4. Discussion Refinement of Primary Aluminum Dendrite Grain The mechanisms of grain refinement are usually explained by the addition of numerous potent
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heterogeneous nuclei and/or the addition of growth restricting solute. According to the Al-Er binary phase diagram [31], the eutectic reaction temperature of the A13Er intermetallic phase and α-Al is 655 ºC at 6.0 wt.% Er. The Er content needs to be higher than 6.0 wt.% for the formation of the
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primary Al3Er as the nuclei.
The heterogeneous nuclei for the α-Al grains nucleation may come from the pre-existed Al3Er
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particles in the melts. Al3Er has the same crystal structure (FCC) with α-Al. In addition, the lattice parameter of Al3Er (a = 0.4215 nm) is very close to that of α-Al (a = 0.4049 nm). Therefore, A13Er can act as the heterogeneous nuclei for the α-Al grains. In the present study, 0.6 wt.% Er addition leads to grain refinement. However, we could not find Al3Er particles in the center of the primary α-
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Al dendrite grain. Further work is, therefore, required to prove the identity of the active nucleant introduced by Er additions to Al-7Si-0.3Mg alloy.
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Grain refinement directly corresponds to the liquidus region on the cooling curve. Therefore, the change of the characteristic temperature in the liquidus region was evaluated by grain refinement.
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Thermal analysis has been used to assess the performance of grain refiners and it is an efficient tool to determine the refining efficiency of a grain refiner [32]. The results clearly show that the addition of erbium to an Al-7Si-0.3Mg alloy has no effect on the recalescence after primary α-Al nucleation, (Figure 1 and Table 1). However, the nucleation temperature of primary α-Al was increased with Er addition. It can be observed that the grain size decreased with an increase in the erbium concentration in both cooling rates (1 ºC/s and 3 ºC/s). This finding may indicate that the addition of Er increased the nucleation temperature of the primary α-Al, thereby, refining the α-Al grains.
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ACCEPTED MANUSCRIPT The effect of the solute atoms in promoting grain refinement is explained in terms of the growth restriction factor of a solute (Q), mC0(k-1), where m is the gradient of the liquidus line, C0 is the solute concentration, and k is the partition coefficient between the melt and the primary α-Al grains. The larger the growth restriction factor, the better grain refining efficiency. For Er in Al, k is ≈ 0 at ≈
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0 C0,max [33]. The very low partition coefficient of Er in Al indicates that grain refinement by Er addition is not the result of the increase of the growth restriction factor.
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Refinement or Modification of Eutectic Silicon
It is well-established that the solidification of Al-Si eutectic proceeds according to three different
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eutectic nucleation and growth modes [26, 34]: (i) nucleation and growth of eutectic silicon from numerous nucleant particles in the vicinity of the primary α-Al dendrite tips, (ii) independent heterogeneous nucleation of eutectic grains in the interdendritic spaces, and (iii) nucleation near the mold wall and growth opposite to the thermal gradient. In the case of unmodified Al-7Si-0.3Mg
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alloy (as shown in Figure 9(a)), plenty of nucleation events for eutectic silicon were found; subsequently, the eutectic grows from many centers throughout the sample in quenched samples. In commercial-purity aluminum foundry alloys, it is well-known that there is a potent nucleant for Si,
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which is typically AlP present due to P as an impurity [20, 22, 35, 36]. The crystal structure of AlP and Si are zincblende and diamond cubic structure; respectively, which is closely related and the
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lattice mismatch is very small (AlP = 5.42 Å [37], Si = 5.43 Å [38]) makes AlP a potent nucleating substrate for Si [22]. The formation of AlP might be present prior to primary α-Al nucleation and then be pushed ahead of aluminum-liquid interfaces, or AlP could nucleate near primary α-Al dendrite tips as the liquid becomes supersaturated in P, or both. Under this situation, the undercooling for the nucleation and growth of Al-Si eutectic only requires approximately 0.6 ºC (Table 1), due to the presence of numerous AlP particles of a few µm in size that are sufficiently large that high undercooling for free growth is not necessary [39].
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ACCEPTED MANUSCRIPT The addition of Er to an Al-7Si-0.3Mg alloy appears to have a significant effect on the nucleation of eutectic Si. In the current study, ErP was directly observed. It is clear from the SEM results (Figure 10) that Er reacts with impurity P to form ErP in preference to AlP. The formation of ErP may also reduce the amount of nucleate particles (AlP). The measured undercooling of eutectic Si from
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thermal analysis (2.5 ºC with 0.6 wt.% Er; Table 1) clearly indicates that there is an interaction between Er and P. It has been suggested that Na3P is formed in the case of Na addition [21], and Sr3P2 in the case of Sr addition [40] instead of AlP. The same hypothesis has been found in other rare
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earth additions, i.e., Sc [29], and Y [41]. The lattice parameters of ErP are a = 5.60 Å [42], and the crystal structure is a rocksalt structure (Figure 11). The crystal structure of ErP is fundamentally
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different from the diamond cubic and zincblende structures of Si and AlP; therefore, ErP is not a good nucleant for eutectic Si compared to AlP. In the present study, there was no microstructural evidence for ErP nucleating Si. According to the IIT theory, Lu and Hellawell [15] stated that an ideal ratio of the atomic radii of the impurity element must be the correct size relative to the radii of
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Si (Si(ri:rSi) = 1.65) and that Er has an atomic ratio (rEr:rSi) of 1.49 [15], which is further from the ideal ratio than potent modifiers [43]. Therefore, Er might not be expected to be a good modifier. However, to analyze the distribution of modifier atoms in the eutectic silicon crystal, many
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techniques have been used, i.e., micro X-ray fluorescence (µ-XRF) [29, 43, 44], nano-secondary ion
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mass spectrometry (NanoSIMS) [45, 46] or atom probe tomography (APT) [27, 47, 48]. Further research could be performed using the techniques mentioned above to clarify the mechanism of Er modification.
The significant change in nucleation frequency as shown in Figure 9 is likely to have significant effect in altering eutectic growth. The decrease in the nucleation frequency of eutectic grains (or a transition from numerous equiaxed eutectic grains to a front growing from the mold walls) can be related to a decrease of the solid-liquid interface area at a given solid fraction. Flood and Hunt [49] note that for constant heat removal, the growth rate is inversely proportional to the solid-liquid 14
ACCEPTED MANUSCRIPT interface area. Therefore, the transition from numerous eutectic grains to a eutectic front growing from the mold walls will increase the mean interface velocity at a given solid fraction. Because the interphase spacing depends on interface velocity, this transition leads to smaller plate spacing and smaller plate thickness. However, to achieve full fibrous eutectic silicon at a satisfying degree of
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modification with Er, an adequate cooling rate is critical. Fading phenomenon is always a problem for conventional grain refiners and modifiers in hypoeutectic Al-Si alloys such as Al-5Ti-B as a grain refiner [50] and Na as a modifier [51]. It is observed that the holding time has no effect on the
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the fading phenomenon by Er addition was not observed.
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SDAS, grain size, and eutectic silicon morphology as shown in Figure 13. It can be concluded that
Mechanical Properties
It is well-know that mechanical properties of Al-Si casting alloy are completely depending on the size of the α-Al primary phase and the morphology, size, and distribution of the eutectic Si. As
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discussed above, refined α-Al primary phase and eutectic Si particles in the Er addition alloys confirm that the Er element influences the mechanical properties. From the results, it can be deduced that the addition of Er causes refinement of the α-Al grains and eutectic Si, which increases the
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tensile strength of the Al-7Si-0.3Mg alloys. Moreover, the addition of Er can reduced the harmful effect of β-phase (Al5FeSi) by changing the morphology of Fe-rich intermetallic compounds from β-
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phase (plate-like) to α-phase (Al, Si, Mg, Er, Fe), as shown in Figure 12. However, it has been found in this study that addition of Er more than 0.40 wt.% leads to a reduction of the tensile strength. The formation of Er-containing phase (Al, Si, Mg, Er), as shown in Figure 12, seems to be responsible for the reduction of tensile properties.
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ACCEPTED MANUSCRIPT 5. Conclusions The microstructure evolution and mechanical properties due to Er additions have been investigated in an Al-7Si-0.3Mg foundry alloy. The key conclusions are as follows:
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1. Er refines the SDAS and α-Al grain. Moreover, Er modifies the eutectic Si morphology from plate toward fibrous morphology.
2. Er additions significantly decrease the nucleation frequency of eutectic Si, due to the
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formation of ErP instead of AlP.
3. The macroscopic eutectic growth mode is the propagation of a defined eutectic front from
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the mold walls opposite to the heat flux direction in this Al-7Si-0.3Mg alloy. 4. The fading phenomenon of 0.4 wt.% Er addition in grain size, SDAS, and eutectic silicon morphology was not observed.
5. The addition of 0.2 wt.% Er significantly improved the tensile strength of the casting in both
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as-cast and T6 conditions, due to the microstructure refinement. However, adding higher amounts of Er, the tensile strength slightly decreased due to the formation of Er-containing intermetallics.
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6. Acknowledgments
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Financial supports from the Thailand Research Fund through the TRF Research Grant for New Researcher (Grant No. TRG5980021), the National Research University Project of Thailand’s Office of the Higher Education Commission and King Mongkut’s University of Technology Thonburi through the “KMUTT 55th Anniversary Commemorative Fund” are acknowledged. The authors would like to express gratitude to Professor Supapan Seraphin for her valuable comments.
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ACCEPTED MANUSCRIPT 7. References [1] U. Patakham, J. Kajornchaiyakul, C. Limmaneevichitr, Grain refinement mechanism in an AlSi-Mg alloy with scandium, Journal of Alloys and Compounds, 542 (2012) 177-186.
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ACCEPTED MANUSCRIPT [42] L. Calvert, P. Villars, Pearson’s handbook of crystallographic data for intermetallic phases, ASM, Materials Park, OH, 1991. [43] K. Nogita, H. Yasuda, M. Yoshiya, S.D. McDonald, K. Uesugi, A. Takeuchi, Y. Suzuki, The
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Determination of strontium segregation in modified hypoeutectic Al-Si alloy by micro X-ray
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fluorescence analysis, Scripta Materialia, 55 (2006) 787-790.
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[49] S.C. Flood, J.D. Hunt, Modification of Al-Si eutectic alloys with Na, Metal Science, 15 (1981) 287-294. [50] C. Limmaneevichitr, W. Eidhed, Fading mechanism of grain refinement of aluminum-silicon alloy with Al-Ti-B grain refiners, Materials Science and Engineering: A, 349 (2003) 197-206.
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ACCEPTED MANUSCRIPT [51] B.M. Thall, B. Chalmers, Modification in aluminium silicon alloys, Journal of the Institute of Metals, 77 (1950) 79-97. 8. List of Table
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Table 1 Variation of primary α-Al dendrite reaction temperatures and Al-Si eutectic reaction temperature: TN, Tmin, TG and ∆TR obtained from the cooling curves in Figure 1 in ºC.
Table 2 Parameters of eutectic silicon particles of Al-7Si-0.3Mg alloys with different
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concentrations of Er addition measurement by image analyzer.
in as-cast and T6 heat-treated conditions. 9. List of Figures
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Table 3 The tensile properties of Al-7Si-0.3Mg alloys with different concentrations of Er addition
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Fig. 1 Assembly of cooling curves take from Al-7Si-0.3Mg alloys with and without Er. Fig. 2 Macrographs of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er at different cooling rates. Images are arranged with increasing Er content within a row: (a) and (e) Er
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free, (b) and (f) 0.2 wt.% Er, (c) and (g) 0.4 wt.% Er, (d) and (h) 0.6 wt.% Er, whereas the cooling rate is different in each column: (a)-(d) 1 ºC/s, (e)-(h) 3 ºC/s.
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Fig. 3 Comparative grain size analysis of Al-7Si-0.3Mg alloys with and without Er addition at different cooling rate.
Fig. 4 Dendritic microstructure of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er at different cooling rates. Images are arranged with increasing Er content within a row: (a) and (e) Er free, (b) and (f) 0.2 wt.% Er, (c) and (g) 0.4 wt.% Er, (d) and (h) 0.6 wt.% Er, whereas the cooling rate is different in each column: (a)-(d) 1 ºC/s, (e)-(h) 3 ºC/s.
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ACCEPTED MANUSCRIPT Fig. 5 Comparative secondary dendrite arm spacing analysis of Al-7Si-0.3Mg alloys with and without Er addition at different cooling rate. Fig. 6 SEM micrograph and EDX analysis at the front of dendrite of the fully sample of Er
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containing: (a) SE image and (b) EDX point analysis. Fig. 7 Optical micrographs at eutectic area of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er at different cooling rates. Images are arranged with increasing Er content within a
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row: (a) and (e) Er free, (b) and (f) 0.2 wt.% Er, (c) and (g) 0.4 wt.% Er, (d) and (h) 0.6 wt.% Er, whereas the cooling rate is different in each column: (a)-(d) 1 ºC/s, (e)-(h) 3 ºC/s.
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Fig. 8 SEM micrographs at eutectic area of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er in graphite mold: (a) Er free, (b) 0.2 wt.% Er, (c) 0.40 wt.% Er, and (d) 0.6 wt.% Er. Fig. 9 Micrographs of quenched samples of Al-7Si-0.3Mg alloys at low magnification: (a) Er-free,
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and (b) 0.6 wt.% Er, and at high magnification: (c) Er-free, and (d) 0.6 wt.% Er. Fig. 10 SEM micrograph and EDX analysis of a eutectic region in quenched sample of Er containing: (a) SE image and (b) EDX point analysis.
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Kikuchi lines.
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Fig. 11 EBSD patterns (a) indexed Kikuchi pattern of the ErP phase and (b) identification of
Fig. 12 SEM micrograph and EDS analysis of an intermetallic in sample of Er-containing: (a) BSE image, EDX mapping (b) Al, (c) Si, (d) Mg and (e) Er, (f) Fe, (g) EDX point analysis at A, (h) EDX point analysis at B and (i) EDX results. Fig. 13 Macrographs and micrographs structure of Al-7Si-0.3Mg-0.4Er with various holding times of (a) 30, (b) 60, (c) 90 and (d) 120 min cast in graphite mold.
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ACCEPTED MANUSCRIPT Table 1 Al-Si eutectic reaction
TN
TMin
TG
∆TR
TN
TMin
TG
∆TR
Er-free
606.3
604.8
607.6
2.8
574.9
573.2
573.8
0.6
0.2 Er
606.5
605.4
608.0
2.6
569.0
565.7
567.7
2.0
0.4 Er
607.6
606.3
608.9
2.6
568.9
567.4
569.6
2.2
0.6 Er
607.4
605.9
608.6
2.7
567.8
565.1
567.6
2.5
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Primary α-Al dendrite reaction
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ACCEPTED MANUSCRIPT Table 2 Er Content Average length Average mean Average aspect (wt.%) of Si phases area of Si ratio of Si (µm) phases (µm2) phases
1 ºC/s
0
12.4 [1.4]
167.5 [19.2]
3.7 [0.4]
0.2
7.5 [1.1]
48.7 [6.4]
3.4 [0.4]
0.4
6.5 [0.5]
35.3 [2.7]
3.3 [0.2]
0.6
5.3 [0.4]
27.8 [4.2]
0
4.9 [0.4]
18.6 [3.8]
0.2
4.3 (0.3]
16.9 [2.0]
0.4
2.5 [0.2]
8.0 [1.6]
2.5 [0.2]
0.6
2.2 [0.2]
7.1 [1.3]
2.5 [0.1]
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4.3 [0.5]
3.4 [0.5]
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3 ºC/s
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Cooling rate
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T6
Tensile strength (MPa)
Elongation (%)
Tensile strength (MPa)
Elongation (%)
Er-free
145.2 [3.9]
2.1 [0.3]
242.6 [11.9]
2.7 [1.7]
0.2 Er
159.2 [9.0]
3.2 [0.8]
280.3 [12.6]
3.3 [0.2]
0.4 Er
147.0 [6.8]
1.7 [0.2]
252.7 [11.9]
1.3 [0.2]
0.6 Er
148.0 [3.7]
1.1 [0.4]
251.1 [10.9]
1.2 [0.7]
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ACCEPTED MANUSCRIPT Highlights •
The addition of Er refines the secondary dendrite arm spacing, primary α-Al grain size and eutectic silicon particles.
•
Erbium additions cause ErP to form instead of AlP and strongly decreases the number
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Erbium modification is therefore mostly due to altered silicon nucleation.
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density of Si nucleation events.
S0925-8388(17)33083-9
DOI:
10.1016/j.jallcom.2017.09.054
Reference:
JALCOM 43110
To appear in:
Journal of Alloys and Compounds
Received Date: 4 February 2017 Revised Date:
4 September 2017
Accepted Date: 5 September 2017
Please cite this article as: P. Pandee, U. Patakham, C. Limmaneevichitr, Microstructural evolution and mechanical properties of Al-7Si-0.3Mg alloys with erbium additions, Journal of Alloys and Compounds (2017), doi: 10.1016/j.jallcom.2017.09.054. 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 proof before it is published in its final 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.
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Microstructural Evolution and Mechanical Properties of Al-7Si-0.3Mg Alloys with Erbium Additions
1
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P. Pandee1* U. Patakham2 and C. Limmaneevichitr1
Department of Production Engineering, Faculty of Engineering, King Mongkut’s University of Technology Thonburi, 126 Pracha-Utid Rd., Bangmod, Tungkhru, Bangkok 10140, Thailand
2
National Metal and Material Technology Center, National Science and Technology Development
*Corresponding author
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Agency, Pathumthani 12120, Thailand
Phromphong Pandee
Telephone: (+66) 82 654 8183, Fax: (+66) 2470 9198
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Email: [email protected]
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ACCEPTED MANUSCRIPT Abstract The mechanical properties of hypoeutectic Al-Si foundry alloys are generally controlled by their microstructures. The present study was performed to determine the influence of erbium (Er) in Al7Si-0.3Mg alloys on the microstructures and mechanical properties. The evolution of microstructure
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was studied by thermal analysis and interrupted solidification technique. Optical microscopy, scanning electron microscope, and energy-dispersive X-ray were employed to characterize the alloy microstructures. The results demonstrated that the addition of Er significantly refined both the
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secondary dendrite arm spacing (SDAS) and the primary α-Al grain. Simultaneously, Er modified eutectic Si phase from coarse acicular plate-like structures into fine fibrous features. We observed
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that Er additions strongly suppress the nucleation of eutectic Si due to the preferential formation of ErP instead of AlP. Er additions change the eutectic growth mode from nucleation and growth on AlP particles to the propagation of a defined eutectic front from the mold walls. Moreover, Er addition that was cast for different holding times indicated that there was no fading effect on the
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SDAS, primary α-Al grain size, and eutectic Si morphology. Importantly, the tensile properties of the Er-modified alloys at as-cast and T6 conditions can be enhanced due to simultaneous refinement
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both the SDAS and the α-Al grains and modification of the eutectic Si. Keyword: Grain refinement, Modification, Secondary dendrite arm spacing, Aluminum-silicon
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alloys, Eutectic solidification, Rare earth
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ACCEPTED MANUSCRIPT 1. Introduction The major microstructural features of commercial hypoeutectic Al-Si foundry alloys are primary αAl dendrites, secondary dendrite arm spacing (SDAS), and Al-Si eutectic structures. Control of these
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microstructures is important in achieving high-quality casting products. It is well-known that reducing SDAS, refining primary α-Al dendrites, and modifying eutectic silicon morphology improves the mechanical properties of hypoeutectic Al-Si foundry alloys. The addition of alloying elements is the main way for obtaining proper microstructure. The addition of alloying elements,
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such as boron, titanium, strontium, or sodium, is one of many methods to achieve grain refinement
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and eutectic silicon modification. However, these elements may only refine the primary α-Al grains or modify the eutectic silicon and cannot do both functions.
In the past several years, many rare earth elements have been studied as micro alloying elements. Rare earth elements can refine primary α-Al grain [1, 2] and/or change the eutectic silicon from a
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coarse plate or flake to a fibrous or lamellar morphology [3-9]. Recently, there has been growing interest in adding rare earth element erbium (Er) into aluminum casting alloys. For example, Zhao et al. [10] revealed that the addition of Er in Al-Zn-Mg-Cu alloys mainly forms an Al3Er, which could
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act as a heterogeneous nucleus to refine grains. Hu et al. [5] reported that Er modified the eutectic silicon from coarse plate-like to fine-fibrous in Al-Si eutectic alloy. Moreover, our previous study
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suggested that the addition of Er to the hypoeutectic Al-Si alloy resulted in modify the eutectic silicon [11]. However, the mechanisms of grain refinement and eutectic silicon modification due to erbium additions are not well understood. The mechanism of grain refinement is usually explained by both the nucleant potency and the solute segregating power [12-14]. It has been suggested that an increase of the nucleant with high potency significantly increases the nucleation rate, resulting in effective grain refinement. In terms of the solute segregating power, it has also been suggested that the contribution of solute is related to the
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ACCEPTED MANUSCRIPT role of solute segregation near the liquid-solid interface in restricting grain growth and constitutional undercooling as a driving-force to further activate nucleation. Several theories, such as restricted nucleation theory and restricted growth theory, were proposed to
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explain the mechanism of chemical modification in hypoeutectic Al-Si alloy. The restricted growth theory is explained by the impurity induced twinning (IIT) mechanism [15, 16] and the twin-plane re-entrant edge (TPRE) mechanism [9, 16-19]. The IIT mechanism proposes that the atoms of the modifier are absorbed on the growth steps of the silicon solid-liquid interface, which causes an
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increase in the twin density of silicon. The TPRE mechanism postulates that a modifier prevents or
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retards the subsequent attachment of silicon atoms during eutectic silicon growth. For restricted nucleation theory, it has been proposed that the AlP particles act as the nucleation site of the eutectic silicon phases [20-23]. The modifying elements can react with AlP nuclei in the melt and alter the number density of eutectic silicon nucleation events, resulting in modification of eutectic silicon. Moreover, chemical modification mechanism is explained with other important factors, such as
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intermetallic clusters and altered surface tension [24-28]. Although considerable research has been devoted to the benefits of erbium additions, microstructure
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evolution during solidification has received considerably less attention. The purpose of this study is to investigate microstructural evolution during solidification to clarify the influence of erbium on the
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hypoeutectic Al-Si alloy and explain the grain refinement and eutectic silicon modification mechanisms as well as the tensile properties. The aim of this research is to compare the chemical grain refinement and chemical modification mechanisms caused by erbium with other previously investigated elements.
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ACCEPTED MANUSCRIPT 2. Experimental Procedure Commercial ingots of Al-7Si-0.3Mg alloy were used in the experiment. The Al-7Si-0.3Mg alloy ingots were manufactured from primary aluminum ingot to minimize the effects of minor element
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contamination, and samples were prepared from the same batch to avoid any undesired variation. Approximately 1 kg of the base alloy (Al-7Si-0.3Mg) was weighed and melted in a silicon carbide crucible at 750 ºC using an induction furnace. Er was added in varying amounts (0.2, 0.4, and 0.6
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wt.%) and the alloys were made using Al-10 wt.% Er master alloy.
Cooling curve analysis was conducted for each alloy in order to determine the characteristic
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temperatures of primary α-Al dendrite and Al-Si eutectic solidification, i.e., nucleation temperature (TN), minimum temperature (TMin), and growth temperature (TG). The cooling rate of the liquid at 650 ºC, just prior to nucleation of primary α-Al phase, was approximately 0.1 ºC/s. Each alloy was then poured at 720 ºC into a graphite cup (outer diameter 50 mm, inner diameter 30 mm, height 50 mm
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and a uniform wall thickness of 10 mm) and a stainless steel cup (outer diameter of bottom 30 mm, outer diameter of upper 34 mm, height 40 mm, and a uniform wall thickness of 0.5 mm). The cooling rate in the graphite cup was approximately 3 ºC/s. The cooling rate in the stainless steel cup was
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approximately 1 ºC/s.
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The fading phenomenon is a key drawback for both grain refinement and eutectic modification in an aluminum foundry alloy. The fading characteristics of Er addition were subsequently evaluated on the SDAS, primary α-Al grain size, and eutectic silicon morphology. The addition of 0.4 wt.% Er was held at 750 ºC for different time intervals. Parts of the melt were poured at different intervals (30, 60, 90, and 120 min) into graphite molds to later observe the resultant microstructure for possible fading of Er. Interrupted solidification experiments were performed to investigate the nucleation behavior of eutectic silicon. A schematic diagram of the experimental set-up is shown in our previous work [29]. 5
ACCEPTED MANUSCRIPT All samples were slowly cooled at approximately 0.1 ºC/s and immediately dropped into an ice-brine quench bath located underneath the holding furnace at a desired time. The fully solidified samples of all experiments were cross-sectioned at 15 mm from the bottom of the
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samples, and the obtained samples were subsequently prepared with standard metallographic procedures. The samples were then etched in Tucker’s reagent (45 ml HCl, 15 ml HNO3, 15 ml HF, 25 ml H2O) to highlight the grain boundaries. The samples were scanned on a high-resolution flatbed scanner to obtain macrographs for the investigation of the grain size. Grain sizes were measured by
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the lineal intercept method with the help of an optical microscope image analyzer system. The SDAS
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values were determined by measuring the SDAS from the corresponding metallography samples using an optical microscope image analyzer system. At least 50 measurements were made for each sample, and the average value of these measurements was used to represent the SDAS value at the corresponding compositions. Moreover, an image analyzer was used to measure three parameters of eutectic silicon particles, i.e., silicon particle length, mean area, and aspect ratio. All three parameters
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were the average values from 20 different images. The microstructures from all samples were recorded by an optical microscope (OM) and JEOL JSM-78000F scanning electron microscope
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(SEM). Intermetallic phases were identified using a Hitachi S-3400N scanning electron microscope equipped with an energy-dispersive X-ray (EDX) detector, EDAX_TSL EBSD hardware and
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OIM_TSL Collection 5.3 software for automated recognition and indexing of EBSD patterns. Castings were also made to prepare tensile test bars according to the ASTM B-108. The mold temperature was preheated at 250 ºC and the pouring temperature is 720 ºC. Tensile testing of as-cast and T6 (solution-treated for 8 h at 540 ºC and aging for 8 h at 155 ºC) samples was carried out at room temperature and a strain rate of 5 mm/min. The present values of tensile properties are the average calculated from five tests for each set of alloy.
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ACCEPTED MANUSCRIPT 3. Results Thermal Analysis An assembly of cooling curves taken from Al-7Si-0.3Mg alloys with and without various Er
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additions is presented in Figure 1. It should be noted that these plots are offset on the time scale to better identify the trend. Three characteristic temperatures are measured from the cooling curves of each alloy, i.e., the nucleation temperature, TN, at which nucleation starts, the minimum temperature,
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TMin, at which nucleation ends, and the growth temperature, TG. The characteristic temperatures of the primary α-Al and Al-Si eutectic solidification obtained from the cooling curves are summarized
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in Table 1.
The nucleation and growth temperatures of primary α-Al dendrite of the Al-7Si-0.3Mg alloy without Er addition are approximately 606.3 ºC and 607.9 ºC. In alloys with Er, the nucleation and growth temperatures of primary α-Al dendrite solidification are slightly increased. At Al-Si eutectic
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solidification, the nucleation temperature was reduced by approximately 5.9 ºC (574.9 ºC - 569.0 ºC) with the addition of 0.2 wt.% Er and later stabilized with an increase in the Er level. Low recalescence undercooling (0.6 ºC) was observed in the Al-7Si-0.3Mg alloy without Er addition.
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Recalescence undercooling was significantly increased with 0.2 wt.% Er addition and then appeared
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to stabilize at approximately 2.0 ºC for higher Er levels. Table 1 shows the growth temperature with similar trends. Upon adding 0.2 wt.% Er, the growth temperature is reduced by approximately 6.1 ºC (573.8 ºC - 567.7 ºC), and it remains constant with further increases in Er content. Primary α-Aluminum Grain
The macrographs of Al-7Si-0.3Mg alloys with various Er additions solidified with low cooling rate (1 ºC/s) are shown in Figures 2(a) to (d) and with high cooling rate (3 ºC/s) are shown in Figures 2(e) to (h). The primary α-Al grain structures of Er-free alloy with both cooling rates are shown in
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ACCEPTED MANUSCRIPT Figures 2(a) and (e). For low cooling rates, large equiaxed grains are observed in samples. For high cooling rates, the sample consists of large equiaxed grains in the center and large columnar grains at the edge. In samples with 0.2 wt.% Er addition, there were slight changes in grain structure with both cooling rates, as shown in Figures 2(b) and (f). However, with an addition of high Er (0.6 wt.%),
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grain refinement was evident, as shown in Figures 2(d) and (h). It is clear that fine equiaxed grains dominated the samples with 0.6 wt.% Er addition. The average grain size and standard deviation measurements with different Er additions are shown in Figure 3. At low cooling rate, adding Er
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significantly reduces the average grain size from 5011 µm to 1319 µm. However, at high cooling
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rate, adding Er slightly reduces the average grain size from 2004 µm to 982 µm. Typical dendritic microstructures of Al-7Si-0.3Mg alloys with various Er additions at a low cooling rate (1 ºC/s) are shown in Figures 4(a) to (d) and a high cooling rate (3 ºC/s) are shown in Figures 4(e) to (h). It can be observed from the micrographs that the addition of Er significantly reduced the SDAS. Figure 5 shows the average SDAS and standard deviation measurements as a function of the
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percentage of Er addition. It can be observed that increasing the levels of Er solute results in a reduction in SDAS to a certain extent. As 0.6 wt.% Er was added to the Al-7Si-0.3Mg alloy, SDAS
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of the Er-free alloy decreased from 79 µm to 63 µm at a low cooling rate and from 32 µm to 18 µm at a high cooling rate. Er exhibits the strongest refining effect on the SDAS for a high cooling rate (3
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ºC/s).
SEM micrograph and EDX point analysis of Al-7Si-0.3Mg alloy with 0.6 wt.% Er addition are presented in Figures 6(a) and (b). The SEM image in Figure 6(a) shows some typical primary phases at the grain boundaries. The intermetallics (as pointed out in Figure 6(a)) were identified by SEMEDX point analysis (Figure 6(b)) to be Al42Si30Er28 in at.%, which is close to Al3Si2Er2 [30]. Al-Si Eutectic Morphology
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ACCEPTED MANUSCRIPT Figure 7 shows Al-Si eutectic structures from an optical microscope taken from samples with different cooling rates and Er additions. The eutectic silicon in the Al-7Si-0.3Mg alloy without Er addition is very coarse and plate-like, as shown in Figure 7(a) (at a low cooling rate) and Figure 7(e) (at a high cooling rate). With the addition of 0.2 wt.% Er (Figures 7(b) and (f)), a decrease in the
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eutectic size was observed, however, the eutectic silicon is still plate-like. In the sample with 0.6 wt.% Er addition, fine fibrous eutectic silicon can be obtained (Figures 7(d) and (h)). The high magnification 3D morphologies of eutectic silicon at high cooling rate (3 ºC/s) with various Er
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additions are depicted in Figure 8. Figure 8(a) show the clustered plate-like eutectic Si occurred in the Er-free alloy. The addition of 0.2 wt.% Er resulted in the refinement of eutectic Si but the
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morphology was still the mixture of plate-like and branched structure, as shown in Figure 8(b). While from Figures 8(c) and (d), it can be seen that the 0.4 and 0.6 wt.% Er modified alloy exhibited a fine fibrous eutectic silicon. The length, mean area, and aspect ratio of eutectic silicon particles of Al-7Si-0.3Mg alloys at different Er additions were determined by an image analyzer and are
Er additions.
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summarized in Table 2. It is clear that all parameters of eutectic silicon decrease with an increase of
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Figure 9 shows the micrographs of Al-7Si-0.3Mg alloys with and without Er additions taken from the quenched samples (at 100 s after the start of Al-Si eutectic solidification). The quenched sample
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without Er addition is shown in Figure 9(a). We found that the eutectic is located at or close to the tips of the primary α-Al dendrites throughout the cross-section of the quenched sample. Polyhedral silicon crystal with a dark particle at the center is shown in Figure 9(c). Our previous work on similar particles [29] has shown that they contain Al, P, and O. A quenched eutectic microstructure of the sample with 0.6 wt.% Er addition is shown in Figure 9(b). When adding Er, Al-Si eutectic is only present adjacent to the mold wall, and there is a eutectic growth front propagating in an opposite direction to the heat flux direction. The typical eutectic growth front of Al-7Si-0.3Mg with 0.6 wt.%
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ACCEPTED MANUSCRIPT Er addition as shown in Figure 9(d). We found a eutectic growth front propagating from bottom to top of the micrograph which corresponds to the direction opposite to the thermal gradient. SEM analysis of Er-containing samples found much fewer P-containing particles. Figure 10 shows a
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typical case in a sample containing 0.6 wt.% Er. The secondary electron (SE) image in Figure 10(a) shows a cubic particle in the sample. EDX point analysis in Figure 10(b) shows that the particle contains Er and P (and negligible Al). In order to confirm that this phase is ErP and Kα peak of Al is overlapping, the crystal structure was determined using EBSD, as shown in Figure 11. Figure 11(a)
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shows the Kikuchi pattern for the ErP phase. The results of the EBSD pattern analysis indicate that
Er-Containing Intermetallic Phases
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the ErP phase has a rock salt structure, as shown in Figure 11(b).
Figure 12 shows SEM analysis of intermetallic in eutectic region in a 0.6 wt.% Er addition sample after natural solidification at 3 ºC/s. The SE image (Figure 12(a)) shows that there are intermetallic
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phases in the eutectic region. EDX mapping such as Figures 12(b) to (f) suggests that intermetallics containing Er. To identify the Er-containing intermetallics, EDX point analysis was performed. Figure 12(g) is EDX point analysis at point A in Figure 12(a) and Figure 12(h) is EDX point analysis
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at point B in Figure 12(a). Figure 12(i) shows EDX results of both points indicating the presence of
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Al, Si, Mg, Er, and Fe for point A and Al, Si, Mg, and Er for point B. Fading Behavior of Erbium
The macrographs and micrographs of samples with 0.4 wt.% Er addition cast after various holding times (30, 60, 90 and 120 min) are presented in Figure 13. It can be clearly observed from Figures 13(a) to (d) that the primary α-Al grain size does not show any significant change with increasing holding time. Moreover, it can be observed from the micrographs of SDAS (Figures 13(e) to (h)),
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ACCEPTED MANUSCRIPT and silicon morphology (Figures 13(i) to (l)) that the holding time has no effect on the SDAS and silicon morphology. It can be concluded that the Er additions do not show any fading effect. Tensile Testing
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The tensile properties of Al-7Si-0.3Mg alloy with various Er additions at as-cast and T6 conditions are listed in Table 3. At the as-cast condition, both the tensile strength and the elongation increased with 0.2 wt.% Er addition and then slightly decreased with an increase of Er content. At the T6
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condition, the tensile strength was increased from 242.6 MPa to 280.3 MPa and elongation was increased from 2.7 % to 3.3 % with 0.2 wt.% Er addition. However, they were decreased when the Er
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content was increased to 0.4 wt.% and 0.6 wt.%.
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ACCEPTED MANUSCRIPT 4. Discussion Refinement of Primary Aluminum Dendrite Grain The mechanisms of grain refinement are usually explained by the addition of numerous potent
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heterogeneous nuclei and/or the addition of growth restricting solute. According to the Al-Er binary phase diagram [31], the eutectic reaction temperature of the A13Er intermetallic phase and α-Al is 655 ºC at 6.0 wt.% Er. The Er content needs to be higher than 6.0 wt.% for the formation of the
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primary Al3Er as the nuclei.
The heterogeneous nuclei for the α-Al grains nucleation may come from the pre-existed Al3Er
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particles in the melts. Al3Er has the same crystal structure (FCC) with α-Al. In addition, the lattice parameter of Al3Er (a = 0.4215 nm) is very close to that of α-Al (a = 0.4049 nm). Therefore, A13Er can act as the heterogeneous nuclei for the α-Al grains. In the present study, 0.6 wt.% Er addition leads to grain refinement. However, we could not find Al3Er particles in the center of the primary α-
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Al dendrite grain. Further work is, therefore, required to prove the identity of the active nucleant introduced by Er additions to Al-7Si-0.3Mg alloy.
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Grain refinement directly corresponds to the liquidus region on the cooling curve. Therefore, the change of the characteristic temperature in the liquidus region was evaluated by grain refinement.
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Thermal analysis has been used to assess the performance of grain refiners and it is an efficient tool to determine the refining efficiency of a grain refiner [32]. The results clearly show that the addition of erbium to an Al-7Si-0.3Mg alloy has no effect on the recalescence after primary α-Al nucleation, (Figure 1 and Table 1). However, the nucleation temperature of primary α-Al was increased with Er addition. It can be observed that the grain size decreased with an increase in the erbium concentration in both cooling rates (1 ºC/s and 3 ºC/s). This finding may indicate that the addition of Er increased the nucleation temperature of the primary α-Al, thereby, refining the α-Al grains.
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ACCEPTED MANUSCRIPT The effect of the solute atoms in promoting grain refinement is explained in terms of the growth restriction factor of a solute (Q), mC0(k-1), where m is the gradient of the liquidus line, C0 is the solute concentration, and k is the partition coefficient between the melt and the primary α-Al grains. The larger the growth restriction factor, the better grain refining efficiency. For Er in Al, k is ≈ 0 at ≈
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0 C0,max [33]. The very low partition coefficient of Er in Al indicates that grain refinement by Er addition is not the result of the increase of the growth restriction factor.
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Refinement or Modification of Eutectic Silicon
It is well-established that the solidification of Al-Si eutectic proceeds according to three different
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eutectic nucleation and growth modes [26, 34]: (i) nucleation and growth of eutectic silicon from numerous nucleant particles in the vicinity of the primary α-Al dendrite tips, (ii) independent heterogeneous nucleation of eutectic grains in the interdendritic spaces, and (iii) nucleation near the mold wall and growth opposite to the thermal gradient. In the case of unmodified Al-7Si-0.3Mg
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alloy (as shown in Figure 9(a)), plenty of nucleation events for eutectic silicon were found; subsequently, the eutectic grows from many centers throughout the sample in quenched samples. In commercial-purity aluminum foundry alloys, it is well-known that there is a potent nucleant for Si,
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which is typically AlP present due to P as an impurity [20, 22, 35, 36]. The crystal structure of AlP and Si are zincblende and diamond cubic structure; respectively, which is closely related and the
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lattice mismatch is very small (AlP = 5.42 Å [37], Si = 5.43 Å [38]) makes AlP a potent nucleating substrate for Si [22]. The formation of AlP might be present prior to primary α-Al nucleation and then be pushed ahead of aluminum-liquid interfaces, or AlP could nucleate near primary α-Al dendrite tips as the liquid becomes supersaturated in P, or both. Under this situation, the undercooling for the nucleation and growth of Al-Si eutectic only requires approximately 0.6 ºC (Table 1), due to the presence of numerous AlP particles of a few µm in size that are sufficiently large that high undercooling for free growth is not necessary [39].
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ACCEPTED MANUSCRIPT The addition of Er to an Al-7Si-0.3Mg alloy appears to have a significant effect on the nucleation of eutectic Si. In the current study, ErP was directly observed. It is clear from the SEM results (Figure 10) that Er reacts with impurity P to form ErP in preference to AlP. The formation of ErP may also reduce the amount of nucleate particles (AlP). The measured undercooling of eutectic Si from
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thermal analysis (2.5 ºC with 0.6 wt.% Er; Table 1) clearly indicates that there is an interaction between Er and P. It has been suggested that Na3P is formed in the case of Na addition [21], and Sr3P2 in the case of Sr addition [40] instead of AlP. The same hypothesis has been found in other rare
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earth additions, i.e., Sc [29], and Y [41]. The lattice parameters of ErP are a = 5.60 Å [42], and the crystal structure is a rocksalt structure (Figure 11). The crystal structure of ErP is fundamentally
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different from the diamond cubic and zincblende structures of Si and AlP; therefore, ErP is not a good nucleant for eutectic Si compared to AlP. In the present study, there was no microstructural evidence for ErP nucleating Si. According to the IIT theory, Lu and Hellawell [15] stated that an ideal ratio of the atomic radii of the impurity element must be the correct size relative to the radii of
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Si (Si(ri:rSi) = 1.65) and that Er has an atomic ratio (rEr:rSi) of 1.49 [15], which is further from the ideal ratio than potent modifiers [43]. Therefore, Er might not be expected to be a good modifier. However, to analyze the distribution of modifier atoms in the eutectic silicon crystal, many
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techniques have been used, i.e., micro X-ray fluorescence (µ-XRF) [29, 43, 44], nano-secondary ion
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mass spectrometry (NanoSIMS) [45, 46] or atom probe tomography (APT) [27, 47, 48]. Further research could be performed using the techniques mentioned above to clarify the mechanism of Er modification.
The significant change in nucleation frequency as shown in Figure 9 is likely to have significant effect in altering eutectic growth. The decrease in the nucleation frequency of eutectic grains (or a transition from numerous equiaxed eutectic grains to a front growing from the mold walls) can be related to a decrease of the solid-liquid interface area at a given solid fraction. Flood and Hunt [49] note that for constant heat removal, the growth rate is inversely proportional to the solid-liquid 14
ACCEPTED MANUSCRIPT interface area. Therefore, the transition from numerous eutectic grains to a eutectic front growing from the mold walls will increase the mean interface velocity at a given solid fraction. Because the interphase spacing depends on interface velocity, this transition leads to smaller plate spacing and smaller plate thickness. However, to achieve full fibrous eutectic silicon at a satisfying degree of
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modification with Er, an adequate cooling rate is critical. Fading phenomenon is always a problem for conventional grain refiners and modifiers in hypoeutectic Al-Si alloys such as Al-5Ti-B as a grain refiner [50] and Na as a modifier [51]. It is observed that the holding time has no effect on the
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the fading phenomenon by Er addition was not observed.
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SDAS, grain size, and eutectic silicon morphology as shown in Figure 13. It can be concluded that
Mechanical Properties
It is well-know that mechanical properties of Al-Si casting alloy are completely depending on the size of the α-Al primary phase and the morphology, size, and distribution of the eutectic Si. As
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discussed above, refined α-Al primary phase and eutectic Si particles in the Er addition alloys confirm that the Er element influences the mechanical properties. From the results, it can be deduced that the addition of Er causes refinement of the α-Al grains and eutectic Si, which increases the
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tensile strength of the Al-7Si-0.3Mg alloys. Moreover, the addition of Er can reduced the harmful effect of β-phase (Al5FeSi) by changing the morphology of Fe-rich intermetallic compounds from β-
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phase (plate-like) to α-phase (Al, Si, Mg, Er, Fe), as shown in Figure 12. However, it has been found in this study that addition of Er more than 0.40 wt.% leads to a reduction of the tensile strength. The formation of Er-containing phase (Al, Si, Mg, Er), as shown in Figure 12, seems to be responsible for the reduction of tensile properties.
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ACCEPTED MANUSCRIPT 5. Conclusions The microstructure evolution and mechanical properties due to Er additions have been investigated in an Al-7Si-0.3Mg foundry alloy. The key conclusions are as follows:
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1. Er refines the SDAS and α-Al grain. Moreover, Er modifies the eutectic Si morphology from plate toward fibrous morphology.
2. Er additions significantly decrease the nucleation frequency of eutectic Si, due to the
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formation of ErP instead of AlP.
3. The macroscopic eutectic growth mode is the propagation of a defined eutectic front from
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the mold walls opposite to the heat flux direction in this Al-7Si-0.3Mg alloy. 4. The fading phenomenon of 0.4 wt.% Er addition in grain size, SDAS, and eutectic silicon morphology was not observed.
5. The addition of 0.2 wt.% Er significantly improved the tensile strength of the casting in both
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as-cast and T6 conditions, due to the microstructure refinement. However, adding higher amounts of Er, the tensile strength slightly decreased due to the formation of Er-containing intermetallics.
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6. Acknowledgments
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Financial supports from the Thailand Research Fund through the TRF Research Grant for New Researcher (Grant No. TRG5980021), the National Research University Project of Thailand’s Office of the Higher Education Commission and King Mongkut’s University of Technology Thonburi through the “KMUTT 55th Anniversary Commemorative Fund” are acknowledged. The authors would like to express gratitude to Professor Supapan Seraphin for her valuable comments.
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[37] C.-Y. Yeh, Z.W. Lu, S. Froyen, A. Zunger, Zinc-blende-wurtzite polytypism in semiconductors, Physical review B, 46 (1992) 10086-10097. [38] M.E. Straumanis, E.Z. Aka, Lattice parameters, coefficients of thermal expansion, and atomic
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weights of purest silicon and germanium, Journal of Applied Physics, 23 (1952) 330-334.
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[39] A.L. Greer, A.M. Bunn, A. Tronche, P.V. Evans, D.J. Bristow, Modelling of inoculation of metallic melts: application to grain refinement of aluminium by Al-Ti-B, Acta Materialia, 48 (2000) 2823-2835.
[40] M. Zarif, B. McKay, P. Schumacher, Study of heterogeneous nucleation of eutectic Si in highpurity Al-Si alloys with Sr addition, Metallurgical and Materials Transactions A, 42 (2011) 16841691. [41] J.H. Li, P. Schumacher, Effect of Y addition and cooling rate on refinement of eutectic Si in Al-5 wt-%Si alloys, International Journal of Cast Metals Research, 25 (2012) 347-357.
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ACCEPTED MANUSCRIPT [42] L. Calvert, P. Villars, Pearson’s handbook of crystallographic data for intermetallic phases, ASM, Materials Park, OH, 1991. [43] K. Nogita, H. Yasuda, M. Yoshiya, S.D. McDonald, K. Uesugi, A. Takeuchi, Y. Suzuki, The
of Alloys and Compounds, 489 (2010) 415-420.
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role of trace element segregation in the eutectic modification of hypoeutectic Al-Si alloys, Journal
[44] K. Nogita, H. Yasuda, A. Takeuchi, K. Yoshida, Y. Suzuki, K. Uesugi, A.K. Dahle,
Determination of strontium segregation in modified hypoeutectic Al-Si alloy by micro X-ray
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fluorescence analysis, Scripta Materialia, 55 (2006) 787-790.
[45] C.J. Simensen, Ø. Nielsen, F. Hillion, J. Voje, NanoSIMS analysis of trace element segregation
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during the Al-Si eutectic reaction, Metallurgical and Materials Transactions A, 38 (2007) 14481451.
[46] C.J. Simensen, F. Hillion, Ø. Nielsen, K. Nogita, A.K. Dahle, Analysis of trace elements in Namodified Al-10 wt%Si using EPMA and NanoSIMS, (ed. H. Jones) Solidification processing,
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(2007) 244-248.
[47] M. Timpel, N. Wanderka, R. Schlesiger, T. Yamamoto, N. Lazarev, D. Isheim, G. Schmitz, S. Matsumura, J. Banhart, The role of strontium in modifying aluminium-silicon alloys, Acta
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Materialia, 60 (2012) 3920-3928.
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[48] J. Barrirero, J. Li, M. Engstler, N. Ghafoor, P. Schumacher, M. Odén, F. Mücklich, Cluster formation at the Si/liquid interface in Sr and Na modified Al–Si alloys, Scripta Materialia, 117 (2016) 16-19.
[49] S.C. Flood, J.D. Hunt, Modification of Al-Si eutectic alloys with Na, Metal Science, 15 (1981) 287-294. [50] C. Limmaneevichitr, W. Eidhed, Fading mechanism of grain refinement of aluminum-silicon alloy with Al-Ti-B grain refiners, Materials Science and Engineering: A, 349 (2003) 197-206.
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ACCEPTED MANUSCRIPT [51] B.M. Thall, B. Chalmers, Modification in aluminium silicon alloys, Journal of the Institute of Metals, 77 (1950) 79-97. 8. List of Table
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Table 1 Variation of primary α-Al dendrite reaction temperatures and Al-Si eutectic reaction temperature: TN, Tmin, TG and ∆TR obtained from the cooling curves in Figure 1 in ºC.
Table 2 Parameters of eutectic silicon particles of Al-7Si-0.3Mg alloys with different
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concentrations of Er addition measurement by image analyzer.
in as-cast and T6 heat-treated conditions. 9. List of Figures
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Table 3 The tensile properties of Al-7Si-0.3Mg alloys with different concentrations of Er addition
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Fig. 1 Assembly of cooling curves take from Al-7Si-0.3Mg alloys with and without Er. Fig. 2 Macrographs of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er at different cooling rates. Images are arranged with increasing Er content within a row: (a) and (e) Er
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free, (b) and (f) 0.2 wt.% Er, (c) and (g) 0.4 wt.% Er, (d) and (h) 0.6 wt.% Er, whereas the cooling rate is different in each column: (a)-(d) 1 ºC/s, (e)-(h) 3 ºC/s.
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Fig. 3 Comparative grain size analysis of Al-7Si-0.3Mg alloys with and without Er addition at different cooling rate.
Fig. 4 Dendritic microstructure of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er at different cooling rates. Images are arranged with increasing Er content within a row: (a) and (e) Er free, (b) and (f) 0.2 wt.% Er, (c) and (g) 0.4 wt.% Er, (d) and (h) 0.6 wt.% Er, whereas the cooling rate is different in each column: (a)-(d) 1 ºC/s, (e)-(h) 3 ºC/s.
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ACCEPTED MANUSCRIPT Fig. 5 Comparative secondary dendrite arm spacing analysis of Al-7Si-0.3Mg alloys with and without Er addition at different cooling rate. Fig. 6 SEM micrograph and EDX analysis at the front of dendrite of the fully sample of Er
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containing: (a) SE image and (b) EDX point analysis. Fig. 7 Optical micrographs at eutectic area of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er at different cooling rates. Images are arranged with increasing Er content within a
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row: (a) and (e) Er free, (b) and (f) 0.2 wt.% Er, (c) and (g) 0.4 wt.% Er, (d) and (h) 0.6 wt.% Er, whereas the cooling rate is different in each column: (a)-(d) 1 ºC/s, (e)-(h) 3 ºC/s.
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Fig. 8 SEM micrographs at eutectic area of fully solidified samples of Al-7Si-0.3Mg alloys with and without Er in graphite mold: (a) Er free, (b) 0.2 wt.% Er, (c) 0.40 wt.% Er, and (d) 0.6 wt.% Er. Fig. 9 Micrographs of quenched samples of Al-7Si-0.3Mg alloys at low magnification: (a) Er-free,
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and (b) 0.6 wt.% Er, and at high magnification: (c) Er-free, and (d) 0.6 wt.% Er. Fig. 10 SEM micrograph and EDX analysis of a eutectic region in quenched sample of Er containing: (a) SE image and (b) EDX point analysis.
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Kikuchi lines.
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Fig. 11 EBSD patterns (a) indexed Kikuchi pattern of the ErP phase and (b) identification of
Fig. 12 SEM micrograph and EDS analysis of an intermetallic in sample of Er-containing: (a) BSE image, EDX mapping (b) Al, (c) Si, (d) Mg and (e) Er, (f) Fe, (g) EDX point analysis at A, (h) EDX point analysis at B and (i) EDX results. Fig. 13 Macrographs and micrographs structure of Al-7Si-0.3Mg-0.4Er with various holding times of (a) 30, (b) 60, (c) 90 and (d) 120 min cast in graphite mold.
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ACCEPTED MANUSCRIPT Table 1 Al-Si eutectic reaction
TN
TMin
TG
∆TR
TN
TMin
TG
∆TR
Er-free
606.3
604.8
607.6
2.8
574.9
573.2
573.8
0.6
0.2 Er
606.5
605.4
608.0
2.6
569.0
565.7
567.7
2.0
0.4 Er
607.6
606.3
608.9
2.6
568.9
567.4
569.6
2.2
0.6 Er
607.4
605.9
608.6
2.7
567.8
565.1
567.6
2.5
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ACCEPTED MANUSCRIPT Table 2 Er Content Average length Average mean Average aspect (wt.%) of Si phases area of Si ratio of Si (µm) phases (µm2) phases
1 ºC/s
0
12.4 [1.4]
167.5 [19.2]
3.7 [0.4]
0.2
7.5 [1.1]
48.7 [6.4]
3.4 [0.4]
0.4
6.5 [0.5]
35.3 [2.7]
3.3 [0.2]
0.6
5.3 [0.4]
27.8 [4.2]
0
4.9 [0.4]
18.6 [3.8]
0.2
4.3 (0.3]
16.9 [2.0]
0.4
2.5 [0.2]
8.0 [1.6]
2.5 [0.2]
0.6
2.2 [0.2]
7.1 [1.3]
2.5 [0.1]
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4.3 [0.5]
3.4 [0.5]
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ACCEPTED MANUSCRIPT Table 3 As-cast
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Tensile strength (MPa)
Elongation (%)
Tensile strength (MPa)
Elongation (%)
Er-free
145.2 [3.9]
2.1 [0.3]
242.6 [11.9]
2.7 [1.7]
0.2 Er
159.2 [9.0]
3.2 [0.8]
280.3 [12.6]
3.3 [0.2]
0.4 Er
147.0 [6.8]
1.7 [0.2]
252.7 [11.9]
1.3 [0.2]
0.6 Er
148.0 [3.7]
1.1 [0.4]
251.1 [10.9]
1.2 [0.7]
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ACCEPTED MANUSCRIPT Highlights •
The addition of Er refines the secondary dendrite arm spacing, primary α-Al grain size and eutectic silicon particles.
•
Erbium additions cause ErP to form instead of AlP and strongly decreases the number
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Erbium modification is therefore mostly due to altered silicon nucleation.
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density of Si nucleation events.