Microstructural Evolution in AlMgSi Alloys during Solidification ... - MDPI

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Mar 10, 2017 - The length of detrimental β-Al5FeSi phases decreased only in AlSi5Fe1.0 alloy ... metallography with measurements done in ImageJ software. ..... solidification), metallurgists traditionally measure λ2 and grain size, or number.
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Microstructural Evolution in AlMgSi Alloys during Solidification under Electromagnetic Stirring Piotr Mikolajczak Institute of Materials Technology, Poznan University of Technology, Piotrowo 3, 60-965 Poznan, Poland; [email protected]; Tel.: +48-61-66-52-804 Academic Editor: Mohsen Asle Zaeem Received: 19 December 2016; Accepted: 6 March 2017; Published: 10 March 2017

Abstract: Equiaxed solidification of AlMgSi alloys with Fe and Mn was studied by electromagnetic stirring to understand the effect of forced flow. The specimens solidified with a low cooling rate, low temperature gradient, and forced convection. Stirring induced by a coil system around the specimens caused a transformation from equiaxed dendritic to rosette morphology with minor dendrites and, occasionally, spheroids. This evolution was quantitatively observed with specific surface Sv . The precipitation sequence of the phases was calculated using the CALPHAD (Computer Coupling of Phase Diagrams and Thermochemistry) technique. Melt flow decreased secondary dendrite arm spacing λ2 in the AlSi5Fe1.0 alloy, while λ2 increased slightly in Mg-containing alloys. The length of detrimental β-Al5 FeSi phases decreased only in AlSi5Fe1.0 alloy under stirring, whereas in Mg-containing alloys, changes to the β-Al5 FeSi phase were negligible; however, in all specimens, the number density increased. The modification of Mn-rich phases, spacing of eutectics, and Mg2 Si phases was analyzed. It was found that the occurrence of Mg2 Si phase regions reduced fluid flow in the late stages of solidification and, consequentially, reduced shortening of β-Al5 FeSi, diminished secondary arm-ripening caused by forced convection, and supported diffusive ripening. However, the Mg2 Si phase was found to have not disturbed stirring in the early stage of solidification, and transformation from dendrites to rosettes was unaffected. Keywords: aluminum alloys; electromagnetic stirring; dendrite arm spacing; rosettes; Mg2 Si phases; solidification

1. Introduction Aluminum alloys have widespread applications, especially in the aerospace and automotive industries. Aluminum–silicon alloys are particularly widely used (e.g., types A355, A356, A357, AK51, AK7, AG10). Increasing demands on such material properties as tensile strength, corrosion behavior, and ductility have pointed to the need for precise control of microstructure through exact casting practice, composition, and heat treatment. In the presence of forced convection, non-dendritic structures may form, with the primary α-Al phase shaped as rosettes or spheroids (globular) [1]. Such non-dendritic structures exhibit distinctive rheological properties, making semisolid metal (SSM) and thixoforming [2] unique for near net shape production of engineering parts. In this technique, semisolid slurries with globular solid particles may be obtained by mechanical stirring or magnetohydrodynamic (MHD) stirring. Rotating magnetic fields can modify the microstructure [3]—for example, by changing direction of dendrites [4]—and improve properties of castings. In this work, the effect of Fe, Mn, and Mg on AlSi alloys was particularly studied. Among many elements used, iron is considered the most deleterious in cast aluminum because the β-Al5 FeSi intermetallic phases that form are particularly detrimental to the ductility of the material [5].

Metals 2017, 7, 89; doi:10.3390/met7030089

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Manganese is the most common alloying addition applied to modify the morphology and type of intermetallic phases in Al alloys. It has been observed previously that Fe and Mn intermetallics, in combination with flow, have different effects on microstructure, and it is possible to transform needle-like Fe phases to blocky ones. Finally, the presence of Mg causes formation of a Mg2 Si phase [6]. This is of interest because recent increases in the application of light Mg alloys and their subsequent recycling may lead to higher amounts of this element in Al alloys. Here, the influence of stirring during slow cooling of bulk specimens solidifying with equiaxed microstructure has been studied as a function of chemical composition (Mg, Fe, and Mn content) based on Al-5 wt. % Si alloy. The specimens were solidified in controlled thermal conditions, including an induced rotating magnetic field (RMF). Microstructural investigation was carried out via optical metallography with measurements done in ImageJ software. Thermo-Calc [7] was used for calculation of the ternary phase diagram and property diagrams in order to determine the sequence of growth of individual precipitating phases. 2. Materials and Methods The study investigated five aluminum alloys (AlSi5Fe1.0, AlMg5Si5, AlMg5Si5Fe1.0, AlMg5Si5Mn1.0, and AlMg5Si5-Fe1.0Mn1.0) which were prepared from pure components: Al (99.999% Hydro Aluminum High Purity GmbH, Grevenbroich, DE), Si (99.9999% NewMet House, Essex, UK), Mg (99.99% NewMet House, Essex, UK), Mn (99.98% NewMet House, Essex, UK), and Fe from ferroaluminum (50 wt. % Al-50 wt. % Fe, Goodfellow Cambridge Ltd., Cambridge, UK). The melt was prepared in an electric resistance furnace using a graphite crucible (50 mm inner diameter). From the beginning of the melting process, a continuous flux of argon flushed the crucible and, after melting, the melt was degassed with argon. No modifier was added. The cylindrical specimens (38 mm diameter and 55 mm height) were melted and solidified in the graphite crucible. Both the crucible and the alloy were heated to 800–810 ◦ C and moved from the furnace into the solidification facility equipped with coils and thermal insulation (Fiberfrax Sibral, Unifrax USA, Tonawanda, NY, USA). The temperature was measured: (1) in the specimen’s center; (2) in the specimen, 4 mm from specimen–crucible surface; and (3) in the crucible, 3 mm from specimen–crucible surface. Measured cooling rates achieved: R800-liq = 0.626 (K/s), Rliq-sol = 0.112 (K/s), and Rsol-470 = 0.280 (K/s) for AlSi5Fe1.0 without stirring. Temperature gradient between the specimen center and the location 4 mm from the specimen’s surface achieved: G800-liq = 0.214 (K/mm) and Gliq-470 = 0.143 (K/mm). These measurements proved that common heating and cooling of specimen and crucible assured continuous slow cooling and simultaneous solidification in the whole specimen by using a low temperature gradient and cooling rate, leading to a microstructure with equiaxed dendrites (without RMF stirring). A rotating magnetic field with a strength of 11 mT (Tesla meter TH26, Aspan, Warsaw, Poland) was generated by electric coils powered by an autotransformer with 45 V and 10 A at a frequency of 50 Hz. The rotational speed was estimated (by camera measurement of the rotating cylinder) to be 2.1 s−1 . The solidified samples (Figure 1) were cut at a height of 10 mm from the bottom for the longitudinal section (Figure 2) and 20 mm from the bottom for the transverse cross-section (Figure 3). The metallographic microsections were prepared using a standard procedure and observed with a light optical microscope (LOM; Nikon, Optiphot-100, Nikon Corp., Tokyo, Japan). In total, 20 sections from 10 experiments were analyzed (5 alloys, each without and with fluid flow). Analysis was performed using Fiji software (ImageJ 1.51a, National Institutes of Health, Bethesda, MD, USA) on the long- and cross-sections in nine specified areas (Figures 2 and 3) using magnifications of 25× and 100×.

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Long-section Long-section Long-section Ø38 mm Ø38 mm top Ø38 mmtop top

bottom bottom bottom

Cross-section Cross-section 55 mm Cross-section 55 mm 55 mm Figure 1. The AlMg5Si5 ingot. The scheme of cutting the long- and cross-section. Figure 1.1.The of cutting cuttingthe thelonglong-and andcross-section. cross-section. Figure TheAlMg5Si5 AlMg5Si5ingot. ingot. The The scheme scheme of Figure 1. The AlMg5Si5 ingot. The scheme of cutting the long- and cross-section.

Figure 2. The scheme of microstructure parameters’ (Table 1) measurement on the long-section—the Figure 2. The scheme of microstructure parameters’ (Table 1) measurement on the long-section—the numbered areas have 1034 parameters’ × 776 μm (magnification 100×). Figure 2. 2. The scheme ofdimensions microstructure (Table 1) 1) measurement measurement onthe thelong-section—the long-section—the Figure The scheme microstructure (Table on numbered areas have of dimensions 1034 ×parameters’ 776 μm (magnification 100×). numbered areas have dimensions 1034 × 776 µm (magnification 100 × ). numbered areas have dimensions 1034 × 776 μm (magnification 100×).

Figure 3. The scheme of microstructure parameters (Table 1) measurement on the cross-section—the Figure 3. The scheme of microstructure parameters (Table 1) measurement on the cross-section—the numbered areas have dimensions 1034 × 776 μm (magnification 100×). Figure 3. The scheme microstructure parameters (Table 1) measurement on the cross-section—the numbered haveofof dimensions 1034 ×parameters 776 μm (magnification 100×). Figure 3. Theareas scheme microstructure (Table 1) measurement on the cross-section—the numbered areas have dimensions 1034 × 776 μm (magnification 100×). numbered havethe dimensions × 776 µmwere (magnification 100secondary ×). For eachareas region, following1034 parameters determined: dendrite arm spacing λ2,

For each region, the following parameters were determined: secondary dendrite arm spacing λ2, specific surface of the dendrites Sv, average length β and number density nβ of β-Al5FeSi platelets, Forsurface each region, following wereLLdetermined: secondary dendrite arm spacing λ2, specific of thethe dendrites Sv,parameters average length β and number density nβ of β-Al5FeSi platelets, and average dimensions L Mn and number density nMn of α-Alsecondary 15Si2Mn4 phases. In5FeSi addition, theλ2 , For each region, the following parameters were determined: arm spacing specific surface of the dendrites S v, average length Lβ and number density nβdendrite of β-Al platelets, and average dimensions LMn and number density nMn of α-Al15Si 2Mn4 phases. In addition, the eutectic spacing λ E for Al–Si eutectics and the spacing λ Mg2Si for Mg 2Si phases was measured. The specific surface of the dendrites S , average length L and number density n of β-Al FeSi platelets, v and average dimensions LMneutectics and number density addition, the β nMn of α-Al15Si2Mn4 phases. β 5 eutectic spacing λE for Al–Si and the spacing λMg2Si for Mg2Si phases wasInmeasured. The secondary dendrite arm spacing λ2 was measured byα-Al averaging the distance between 20 60 and average dimensions and number density nMn of Si2 Mn phases. Inwas addition, theand eutectic eutectic spacing λ E forLAl–Si eutectics and the spacing λ Mg2Si for Mg 2Si phases measured. The Mn 15 4 secondary dendrite arm spacing λ2 was measured by averaging the distance between 20 and 60 adjacent side branches along the primary dendrite stem. Specific surface of the dendrites S v was spacing λE for Al–Si eutectics and λMg2Si for Mg was measured. The secondary dendrite armalong spacing λprimary 2 spacing was measured by averaging the distance between 20secondary and 60 2 Si phases adjacent side branches thethe dendrite stem. Specific surface of the dendrites Sv was calculated from the measured perimeter and the enclosed area of the α-Al dendrites. In the adjacent side branches alongmeasured theperimeter primary dendrite Specific theand dendrites SInv was dendrite arm spacing by averaging the distance between 20 60 adjacent side 2 was calculated from the λmeasured and the stem. enclosed area surface of the of α-Al dendrites. the measurement of the about 6500 Fe-rich intermetallics, only needles with aα-Al thickness >3 μm and calculated from measured perimeter and the enclosed area of the dendrites. In the branches along the primary dendrite stem. Specific surface of the dendrites S was calculated from v measurement of about 6500 Fe-rich intermetallics, only needles with a thickness >3 μm and length/thickness ratio >5 were considered. In the only measurement of complex-shaped Mn-rich measurement of ratio about intermetallics, needles with a thickness >3 μm thelength/thickness measured perimeter and theFe-rich enclosed area In of the dendrites. the measurement of and about >56500 were considered. the α-Al measurement ofIn complex-shaped Mn-rich intermetallics, we considered the overall dimensions of each precipitate. The eutectic spacing λE was length/thickness ratio >5 were considered. In the measurement of complex-shaped Mn-rich 6500 Fe-rich intermetallics, only needles a thickness >3 precipitate. µm and length/thickness ratioλ>5 were intermetallics, we considered the overallwith dimensions of each The eutectic spacing E was measured by averaging the distance LE between adjacent eutectic plates; the same was done for intermetallics, wemeasurement considered the of eacheutectic precipitate. Thewe eutectic λEoverall was considered. In of overall complex-shaped Mn-rich intermetallics, considered the measured bythe averaging the distance LE dimensions between adjacent plates; the samespacing was done for Mg2Si phases. measured byeach averaging the distance LE between adjacent plates; the same was done forL dimensions of precipitate. The eutectic spacing λE waseutectic measured by averaging the distance Mg 2Si phases. E Mg 2Si phases. between adjacent eutectic plates; the same was done for Mg2 Si phases.

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3. Results 3. Results 3. Results Microstructure on long- and cross-sections of solidified specimens was investigated in 2D 3. Results Microstructure on long- and cross-sections of solidified specimens was investigated in 2D micrographs using LOM. Theand measured parameters are collected in Table The ternary in phase Microstructure on longcross-sections of solidified specimens was1. investigated 2D micrographs using LOM. The measured parameters are collected in Table 1. The ternary phase diagram, Scheil solidification, andcross-sections property diagrams calculated in Thermo-Calc Microstructure on longof solidified specimens was investigated in the 2D micrographs using LOM. Theand measured parameters arewere collected in Table 1. The ternaryfor phase diagram, Scheil solidification, and property diagrams were calculated in Thermo-Calc for the studied alloys. micrographs usingsolidification, LOM. The measured parameters are collected Table 1. The phase diagram, diagram, Scheil and property diagrams were in calculated in ternary Thermo-Calc for the studied alloys. Scheil solidification, and property diagrams were calculated in Thermo-Calc for the studied alloys. studied alloys. 3.1. Microstructure 3.1. Microstructure 3.1. Microstructure 3.1. Microstructure Figure 4 shows, in AlSi5Fe1.0 alloy micrographs, the typical structures obtained in experiments Figure 4 shows, shows, in in AlSi5Fe1.0 AlSi5Fe1.0 alloy alloy micrographs, micrographs, the the typical typical structures structures obtained obtained in in experiments experiments Figure for solidification without and with Solidification without flow is characterized by Figure 44 shows, in AlSi5Fe1.0 alloy stirring. micrographs, the typical structures obtained in experiments for solidification solidificationwithout without and with stirring. Solidification without flow is characterized by for andand with stirring. Solidification without flow is characterized by well-formed well-formed α-Al without dendrites, while for melt stirring, α-Al formed asflow rosettes and, rarely, by as for solidification with stirring. Solidification without is characterized well-formed α-Al dendrites, while for melt stirring, α-Al formed as rosettes and, rarely, as α-Al dendrites, while for melt stirring, α-Al formed as rosettes and, rarely, as incompletely formed incompletely α-Al formed dendriteswhile or asfor spheroids (globular forms). Figure 5 shows and, α-Al rarely, dendrites well-formed dendrites, melt stirring, α-Al formed as rosettes as incompletely formed dendrites or forms). as spheroids (globular forms). Figure(white), 5 shows(α-Al)-Si α-Al dendrites dendrites or asformed spheroids (globular Figure 5 shows eutectic (white), (α-Al)-Si eutectic (grey), andasβ-Al 5FeSi phases (darkα-Al grey)dendrites inFigure the form of needles spread over incompletely dendrites or spheroids (globular forms). 5 shows α-Al dendrites (white), (α-Al)-Si eutectic (grey), and β-Al 5FeSi phases (dark grey) in the form of needles spread over (grey), and β-Al FeSi phases (dark grey) in the form of needles spread over the entire sample and 5 the entire sample and only slightly visible, very small β needles. The micrographs (Figures 6 and 7) (white), (α-Al)-Si eutectic (grey), and β-Al5FeSi phases (dark grey) in the form of needles spread over the entire sample andvery onlysmall slightly visible, very small β needles. The micrographs (Figures 6 and 7) only slightly visible, β needles. The micrographs (Figures 6 and 7) for AlMg5Si5Fe1.0 for entire AlMg5Si5Fe1.0 additionally Mgvery 2Si (very phases between white α-Al the sample and only slightlyshow visible, small dark) β needles. Thearranged micrographs (Figures 6 and 7) for AlMg5Si5Fe1.0 additionally showphases Mg2Si (very dark) phases arranged between (or white α-Al additionally show Mg dark) between α-Al dendrites rosettes), 2 Si (very dendrites (or rosettes), (α-Al)–Si eutectic (grey), and β-Alwhite 5FeSi needles (dark white grey). For for AlMg5Si5Fe1.0 additionally show Mg2Siarranged (very dark) phases arranged between α-Al dendriteseutectic (or rosettes), (α-Al)–Si eutectic (grey), and β-Al 5FeSi needles (dark grey). For (α-Al)–Si (grey),the and β-Al5 FeSi needles (dark grey). AlMg5Si5Mn1.0 alloys, the Fe-rich β AlMg5Si5Mn1.0 alloys, Fe-rich β needles were replaced by Mn-rich well-developed dendrites (or rosettes), (α-Al)–Si eutectic (grey), and For β-Al 5FeSi phases needleswith (dark grey). For AlMg5Si5Mn1.0 alloys, the Fe-rich β needles were replaced by Mn-rich phases8a) with well-developed needles were replaced by Mn-rich phases with well-developed shapes (Figure or a more compact shapes (Figure 8a) or a more compact form (Figure 8b). For all solidified AlMg5Si5Mn1.0 alloys, the Fe-rich β needles were replaced by alloys Mn-rich phases with electromagnetic well-developed shapes (Figure 8a) or aall more compact form (Figure 8b). For all alloys solidified with electromagnetic form (Figure 8b). alloys solidified electromagnetic the dendrites changed into4 stirring, the dendrites into rosettes, minor dendrites, orstirring, (occasionally) spheroids (Figures shapes (Figure 8a)For or achanged more compact formwith (Figure 8b). For all alloys solidified with electromagnetic stirring, the dendrites changed into rosettes, minor dendrites, orand (occasionally) spheroids (Figures 4 rosettes, minor dendrites, or (occasionally) spheroids (Figures 4 6). and 6). the dendrites changed into rosettes, minor dendrites, or (occasionally) spheroids (Figures 4 stirring, and 6). and 6).

(a) (b) (a) (b) (a) (b)without and (b) with Figure 4. The microstructures of the AlSi5Fe1.0 specimen solidified: (a) Figure The microstructures microstructures of of the the AlSi5Fe1.0 (a) without without and and (b) Figure 4. 4. The AlSi5Fe1.0 specimen specimen solidified: solidified: (a) (b) with with

electromagnetic stirring. Light Light optical magnification Figure 4. The microstructures of themicroscope AlSi5Fe1.0(LOM), specimen solidified: 25×. (a) electromagnetic ×. without and (b) with electromagnetic stirring. stirring. Light optical optical microscope microscope (LOM), (LOM), magnification magnification 25 25×. electromagnetic stirring. Light optical microscope (LOM), magnification 25×.

Figure 5. The microstructures of the AlSi5Fe1.0 specimen solidified with electromagnetic stirring. Figure 5. 5. The microstructures of the AlSi5Fe1.0 specimen solidified with electromagnetic stirring. Figure The microstructures microstructures of the the AlSi5Fe1.0 specimen solidified with with electromagnetic electromagnetic stirring. stirring. LOM, magnification 100×. Clearly visible β-Al5FeSi phases. solidified Figure 5. The of AlSi5Fe1.0 specimen LOM, magnification 100 × . Clearly visible β-Al 5FeSi phases. LOM, magnification 100 × . Clearly visible β-Al FeSi phases. 5 LOM, magnification 100×. Clearly visible β-Al5FeSi phases.

(a) (b) (a) (b) (a) (b) Figure 6. The microstructures of the AlMg5Si5Fe1.0 specimen solidified: (a) without and (b) with Figure 6. The microstructures of the AlMg5Si5Fe1.0 specimen solidified: (a) without and (b) with electromagnetic stirring. LOM, of magnification 25×. Figure 6. The the AlMg5Si5Fe1.0 specimen with Figure 6. The microstructures microstructures the AlMg5Si5Fe1.0 specimen solidified: solidified: (a) (a) without without and and (b) (b) with electromagnetic stirring. LOM, of magnification 25×. electromagnetic stirring. LOM, magnification 25 × . electromagnetic stirring. LOM, magnification 25×.

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Figure 7. 7. The microstructures of the AlMg5Si5Fe1.0 specimen solidified with with electromagnetic electromagnetic stirring. stirring. Figure The microstructures microstructures of of the AlMg5Si5Fe1.0 AlMg5Si5Fe1.0 specimen specimen solidified solidified Figure 7. The the with electromagnetic stirring. LOM, magnification 100 × . Highly visible β-Al 5FeSi phases. LOM, Highlyvisible visibleβ-Al β-Al5FeSi phases. 5 FeSiphases. LOM, magnification magnification 100 100× ×.. Highly

(a) (a)

(b) (b)

Figure 8. The microstructures of the AlMg5Si5Mn1.0 specimen solidified with electromagnetic Figure 8.8.The Themicrostructures microstructures of the AlMg5Si5Mn1.0 specimen solidifiedelectromagnetic with electromagnetic Figure of (RMF). the AlMg5Si5Mn1.0 specimen solidified stirring stirring rotating magnetic field LOM, magnification 100. Highlywith visible Mn-rich phases: (a) stirring rotating magnetic field (RMF). LOM, magnification 100. Highly visible Mn-rich phases: (a) rotating magnetic fieldon(RMF). LOM, 100.between Highly dendrites. visible Mn-rich phases: (a) α-Al α-Al dendrites grown Mn-phase; (b)magnification Mn phase grown α-Al dendrites grown on Mn-phase; (b)phase Mn phase between dendrites. dendrites grown on Mn-phase; (b) Mn growngrown between dendrites.

3.2. Parameters Characterizing Microstructure 3.2. Parameters Characterizing Microstructure 3.2. Parameters Characterizing Microstructure The microstructure evolution caused by electromagnetic stirring is characterized by parameters The microstructure evolution caused by electromagnetic stirring is characterized by parameters The evolution electromagnetic characterized byand parameters measuredmicrostructure and counted in specifiedcaused areas by (Figures 1–3): nine stirring areas onisthe cross-section nine on measured and counted in specified areas (Figures 1–3): nine areas on the cross-section and nine on measured and counted in specified areas (Figures areasand on the and nine on the the long section. The results are almost equal in1–3): all 18nine figures, no cross-section trend was found across the long section. The results are almost equal in all 18 figures, and no trend was found across the long section.Such The results are almost equal inaallcomprehensive 18 figures, and no trend was acrossspecimen, the specimens. specimens. methodology provides overview of found the whole and specimens. Such methodology provides a comprehensive overview of the whole specimen, and Such methodology a comprehensive overview of the whole specimen, credible results are credible results areprovides presented in Table 1. For AlSi5Fe1.0, induced fluid flowand caused a decrease of credible results are presented in Table 1. For AlSi5Fe1.0, induced fluid flow caused a decrease of presented 1. For AlSi5Fe1.0, induced fluid flow caused a decrease of about 12% in secondary about 12%ininTable secondary dendrite arm spacing λ2 (from 87 to 77 µm), while for the other alloys, λ2 about 12% in secondary dendrite arm spacing λ2 (from 87 to 77 µm), while for the other alloys, λ2 dendrite arm spacing 87 to 77from µm),60while theClearly other alloys, seems unchanged andλreaches to 68for µm. visible λin Figuresunchanged 4 and 6 isand the 2 (from values 2 seems seems unchanged and reaches values from 60 to 68 µm. Clearly visible in Figures 4 and 6 is the reaches values 60 to to 68rosette µm. Clearly Figures 6 is the from dendritic evolution from from dendritic shape,visible whichinresults in 4a and decrease in evolution the specific surface Sv of evolution from dendritic to rosette shape, which results in a decrease in the specific surface −1Sv of to rosette shape, which in a decrease the specific surface Sv offound α-Al primary phase all α-Al primary phase forresults all alloys. For the in AlSi5Fe1.0 alloy, Sv was to be 0.023 µmfor for α-Al primary phase for all alloys. For the −1AlSi5Fe1.0 alloy, Sv was found to be 0.023 µm−1 for − 1 alloys. For thewithout AlSi5Fe1.0 alloy,and Sv was found be 0.023 µm forfor solidification without stirring anda solidification stirring 0.019 µm towith flow, while, other alloys, stirring caused solidification without stirring −1and 0.019 µm−1 with flow, while, for other alloys, stirring caused a 1 with 1 to −1. Stirring 0.019 µm− flow, while, stirring causedcaused a decrease from about 0.036 µm−time decrease from about 0.036 µmfortoother aboutalloys, 0.026 µm a decrease in solidification decrease from about 0.036 µm−1 to about 0.026 µm−1. Stirring caused a decrease in solidification time − 1 about µmalloys. . Stirring caused a decrease in solidification time for all studied alloys. for all 0.026 studied for all studied alloys. Fe-rich intermetallics were characterized characterized by by average average length lengthLLββ and by by number number density densitynnββ. It is Fe-rich intermetallics were characterized by average length Lβ and by number density nβ. It is clear (Table flow caused shortening of βofphases for AlSi5Fe1.0 alloy, alloy, where where averageaverage length (Table1)1)that thatfluid fluid flow caused shortening β phases for AlSi5Fe1.0 clear (Table 1) that fluid flow caused shortening of β phases for AlSi5Fe1.0 alloy, where average L decreased by about 20%, from 71 to 57 µm. The flow effect is much smaller for AlMg5Si5Fe1.0, length L β decreased by about 20%, from 71 to 57 µm. The flow effect is much smaller for β length Lβ decreased by about 20%, from 71 to 57 µm. The flow effect is much smaller for about 5%, and almost AlMg5Si5Fe1.0Mn1.0. For all alloys, electromagnetic AlMg5Si5Fe1.0, aboutnegligible 5%, andforalmost negligible for AlMg5Si5Fe1.0Mn1.0. For all stirring alloys, AlMg5Si5Fe1.0, about 5%, and almost negligible for AlMg5Si5Fe1.0Mn1.0. For all alloys, induced a higherstirring numberinduced densityaofhigher β-Al5 number FeSi intermetallics. electromagnetic density of β-Al5FeSi intermetallics. electromagnetic stirring induced a higher number density of β-Al5FeSi intermetallics. Mn-rich byby measurement of the average overall dimension LMn Mn-rich intermetallics intermetallicswere werecharacterized characterized measurement of the average overall dimension Mn-rich intermetallics were characterized by measurement of the average overall dimension and densitydensity nMn of mostly precipitates.precipitates. Fluid flow decreased the average overall LMn number and number nMn ofcomplex-shaped mostly complex-shaped Fluid flow decreased the LMn and number density nMn of mostly complex-shaped precipitates. Fluid flow decreased the 2. dimension LMn by about 9%,Lfrom to 273, increased number density from 20 to density 27 mm−from average overall dimension Mn by299 about 9%,and from 299 to 273, and increased number 20 average overall dimension LMn by about 9%, from 299 to 273, and increased number density from 20 −2 eutectic, the influence of stirring is unclear; there is no direct modification to the to 27For mm(α-Al)–Si . to 27 mm−2. eutectic λE .eutectic, For magnesium-rich Si phases, the spacing of λisMg2Si seemsmodification to decrease weakly Forspacing (α-Al)–Si the influenceMg of2stirring is unclear; there no direct to the For (α-Al)–Si eutectic, the influence of stirring is unclear; there is no direct modification to the under forced convection. spacing measured in the the range from 5of to λ 18Mg2Si µm.seems to decrease eutectic spacing λE. For The magnesium-rich Mg2Siwas phases, spacing eutectic spacing λE. For magnesium-rich Mg2Si phases, the spacing of λMg2Si seems to decrease These measurements were performed onmeasured a large number weakly under forced convection. The spacing was in of thegrains rangeand fromintermetallics 5 to 18 µm. in order weakly under forced convection. The spacing measured was in the range from 5 to 18 µm. to provide reliable results. In Table 1, numbers of grains andand dendrite arms inspected These highly measurements were performed onthe a large number of grains intermetallics in orderare to These measurements were performed on a large number of grains and intermetallics in order to provide highly reliable results. In Table 1, the numbers of grains and dendrite arms inspected are provide highly reliable results. In Table 1, the numbers of grains and dendrite arms inspected are included in parentheses, while the standard deviations are listed in square brackets. For example, for included in parentheses, while the standard deviations are listed in square brackets. For example, for

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included in parentheses, while the standard deviations are listed in square brackets. For example, for AlSi5Fe1.0 FeSi AlSi5Fe1.0 solidified solidified without without stirring stirring we we measured measured 60 60 grains, grains, 652 652 dendrite dendrite arms, arms, and and 1571 1571 β-Al β-Al55FeSi intermetallic phases; 12% variation variation in from intermetallic phases; standard standard deviation deviationfor forλλ22 was was 6.7 6.7 µm. µm. Stirring Stirringcaused caused− −12% in λ λ22,, from 87 to 77 µm, and a decrease in solidification time from 468 to 393 s. 87 to 77 µm, and a decrease in solidification time from 468 to 393 s. Table 1. Microstructure parameters measured on inspected micrographs. Table 1. Microstructure parameters measured on inspected micrographs. Microstructure MicrostructureParameters Parameters Fe-phases (β-Al5FeSi) Mn-phases Fe-phases (β-Al5 FeSi) Mn-phases nβ nMn Lβ (µm) LMn (µm) −2 (mm−2n)β (mm nMn) λSDAS (µm) Lβ (µm) L (µm) Sv (µm−1 ) Mn 71 [4.1] (mm−2 ) (mm−2 ) 87 [6.7] (60/652) 0.023 [0.002] 109 87 [6.7] (60/652) 0.023 [0.002] (1571) 71 [4.1] (1571) 109 77 [5.7] (68/631) 0.019 [0.001] 57 [4.1] 160 77 [5.7] (68/631) 0.019 [0.001] 57 [4.1] (2306) 160 (−12%) (−17%) (2306) (−20%) (47%) (−12%) (−17%) (−20%) (47%) 67 [5.1] 0.036 [0.004] 67 [5.1] (134/1135) 0.036 [0.004] (134/1135) 68 [4.0] 0.024 [0.002] 0.024 [0.002] (109/880) 68 [4.0] (109/880) (−33%) (−33%) 60 [4.5] 79 [4.9] 0.034 [0.003] 72 60 [4.5] (153/1399) 0.034 [0.003] (1043) 79 [4.9] (1043) 72 (153/1399) 62 [4.0] 0.026 [0.002] 75 [5.4] 82 0.026 [0.002] 75 [5.4] (1178) 82 (150/864) (1178) (−5%) (14%) 62 [4.0] (150/864) (−24%) (−24%) (−5%) (14%) 65 [5.0] 0.036 [0.003] 299 [13.3] 20 65 [5.0] (136/1054) 0.036 [0.003] 299 [13.3] 20 (136/1054) 65 [4.9] 0.028 [0.002] 273 [15.1] 27 0.028 [0.002] 273 [15.1] (117/599) (−9%) 65 [4.9] (117/599) (−24%) 27 (−24%) (−9%) 62 [4.5] 0.035 [0.002] 44 [2.2] (128) 9 (109/678) 62 [4.5] (109/678) 0.035 [0.002] 44 [2.2] (128) 9 62 [5.2] 0.028 [0.002] 43 [2.4] (229) 16 0.028 [0.002] (−2%) 43 [2.4] (229) (77%)16 (116/382) (−20%) 62 [5.2] (116/382) (−20%) (−2%) (77%)

Aluminium RMF (mT) Dendrites Dendrites Aluminium RMF(s)} (mT) Alloys {Solid. Time Alloys {Solid. Time (s)} λSDAS (µm) Sv (µm−1)

AlSi5Fe1 AlSi5Fe1

AlMg5Si5 AlMg5Si5

AlMg5Si5Fe1 AlMg5Si5Fe1

AlMg5Si5Mn1 AlMg5Si5Mn1

AlMg5Si5Fe1Mn1 AlMg5Si5Fe1Mn1

0{468} 0{468} 11{393} 11{393} 0{584} 0{584} 11{495} 11{495} 0{537} 0{537} 11{477} 11{477} 0{686} 0{686} 11{638} 11{638} 0{723} 0{723} 11{627} 11{627}

AlSi Mg2Si AlSi eutectics Mg2 Si eutectics λMg2Si λEut (µm) λEut λMg2Si (µm) (µm) (µm) 8.5 [0.3] 8.5 [0.3] 9.5 [0.3] 9.5 [0.3]

-

-

9.7 [0.5] 9.8 [0.3] 9.7 [0.5] (5–20) 9.8 [0.3] (5–20) 10.1 [0.5] 8.8 [0.3] 10.1 [0.5] (4.8–18.6) 8.8 [0.3] (4.8–18.6) 11.7 [0.6] 11.7 [0.6]11.7 [0.6] 11.7 (5.3–19.3) [0.6] (5.3–19.3) 12.7 [0.4] 6.8 [0.2] 12.7 [0.4] (6.3–18.2) 6.8 [0.2] (6.3–18.2) 11.6 [0.7] 7.5 [0.3] 11.6 [0.7] (5.6–19.0) 7.5 [0.3] (5.6–19.0) 12.0 [0.7] 8.7 [0.2] 12.0 [0.7] (4.3–20.6) 8.7 [0.2] (4.3–20.6) 11.3 [0.7] 8.1 [0.3] 11.3 [0.7] (5.1–18.4) 8.1 [0.3] (5.1–18.4) 12.1 [0.7] 9.1 [0.2] 12.1 [0.7] (5.8–19.1) 9.1 [0.2] (5.8–19.1)

(1) Curly brackets {…} present solidification time [s]; (2) brackets […] present standard deviation; (1) Curly brackets { . . . } present solidification time [s]; (2) brackets [ . . . ] present standard deviation; (3) parentheses (3) (…/…) of present numbersandofdendrite grainsarms inspected dendrite ( .arms counted; (4) ( . . . parentheses / . . . ) present numbers grains inspected counted; and (4) parentheses . . %) present percent variation of the(…%) parameters under electromagnetic stirring; parentheses (under . . . – . . .electromagnetic ) present the rangestirring; of measured parentheses present percent variation of the(5)parameters (5) eutectic spacings. parentheses (…–…) present the range of measured eutectic spacings.

3.3. Precipitation Precipitation Sequence Sequence 3.3. On the the basis of the phase diagram diagram [7,8] [7,8] and and property diagram (Figure (Figure 9), 9), one On basis of the ternary ternary Al–Fe–Si Al–Fe–Si phase property diagram one ◦ C (L → α-Al + L); then, the liquid can state that for the AlSi5Fe1.0 alloy: first, α-Al will form at 627.4 can state that for the AlSi5Fe1.0 alloy: first, α-Al will form at 627.4 °C (L → α-Al + L); then, the liquid enriches in in Si Si and and Fe Fe to to aa concentration concentration of of 7.77 7.77 %Si %Si and and 1.63 1.63 %Fe %Fe until until it it reaches reaches the the eutectic eutectic reaction reaction enriches ◦ at 608.4 C. Next, FeSi ++ L L starting starting from from at 608.4 °C. Next, solidification solidification follows follows the the eutectic eutectic groove groove LL → → α-Al α-Al ++ β-Al β-Al55FeSi ◦ C and ending at 575 ◦ C with composition 12.6 %Si and 0.9 %Fe, where the final eutectic reaction 608.4 608.4 °C and ending at 575 °C with composition 12.6 %Si and 0.9 %Fe, where the final eutectic L → α-AlL+ → β-Al Si2β-Al + Si9Fe occurs. At this point (temperature 575 ◦ C), β-Al FeSi mass fraction reaches 9 Fe2+ reaction α-Al 2Si2 + Si occurs. At this point (temperature5 575 °C), β-Al5FeSi mass f β = 3.7%, α-Al mass fraction f α = 93.31%, mass fractionmass reaches f Eut =reaches 2.98%. fraction reaches fβ = 3.7%, α-Alreaches mass fraction reachesand fα =eutectics 93.31%, and eutectics fraction ◦ C: f = 3.7%, The phase fractions continue to evolve in the solid until, finally, at the temperature 20 β 20 °C: fEut = 2.98%. The phase fractions continue to evolve in the solid until, finally, at the temperature ffβα==3.7%, 91.63%, and f = 4.67%. Eut fα = 91.63%, and fEut = 4.67%.

Figure Figure 9. 9. Property Property diagram diagram for for AlSi5Fe1.0 AlSi5Fe1.0 alloy. alloy.

Analyzing the solidification path on the ternary Al–Mg–Si phase diagram (Figure 10) and property diagram (Figure 11), one can state that for the AlMg5Si5 alloy: first, α-Al forms at 608.0 °C (L → α-Al + L); then, the sample enriches in Si to a concentration of 7.46 %Si and 6.99 %Mg, where it

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Analyzing the solidification path on the ternary Al–Mg–Si phase diagram (Figure 10) and property diagram 11), oneatcan state°Cthat forα-Al the AlMg5Si5 alloy: first, α-Al atNext, 608.0 ◦ C reaches the reaction 581.7 and mass fraction reaches fα = forms 37.3%. Metals 2017,eutectic 7,(Figure 89 7 of 16 (L →solidification α-Al + L); then, thethe sample enriches Si α-Al to a concentration of 7.46 %Si and°C6.99 it follows eutectic groove in L→ + Mg2Si + L starting from 581.7 and%Mg, endingwhere at eutectic reaction 581.7 °C and where α-Al mass fraction reaches fα = L 37.3%. Next, ◦C 558.6 °Ceutectic with the composition %Siat and 4.5α-Al %Mg, the final eutectic → α-Al + Mg2Si reaches thereaches reaction at12.5 581.7 and mass fraction reaches f αreaction = 37.3%. Next, solidification solidification follows the eutectic groove L → α-Al + Mg2Si + L starting from °C and ending at ◦ ◦ C 581.7 + Sithe occurs. follows eutectic groove L →12.5 α-Al Mg4.5 + L starting 581.7 reaction andLending 558.6 C with 2 Si%Mg, 558.6 °C with composition %Si+and where the from final eutectic → α-Al +atMg 2Si composition %Si and 4.5 %Mg, where the final eutectic reaction L → α-Al + Mg2 Si + Si occurs. + Si12.5 occurs.

10. Ternary phase diagram—Al–Si–Mg system. Liquidus projection with red arrows marking Figure 10.Figure Ternary phase diagram—Al–Si–Mg system. Liquidus projection with red arrows marking Figurethe 10. Ternary phase diagram—Al–Si–Mg system. with red arrows marking solidification path (Scheil–Gulliver solidification) forLiquidus AlMg5Si5projection alloy. the solidification path (Scheil–Gulliver for AlMg5Si5 AlMg5Si5alloy. alloy. the solidification path (Scheil–Gulliversolidification) solidification) for

Figure 11. Property diagram for AlMg5Si5 alloy.

For AlMg5Si5Fe1.0 alloy (Figure 12): first, α-Al starts to precipitate at 603.8 °C (L → α-Al + L) Figure 11. Property diagram AlMg5Si5 alloy. and continues until 592.5 °C 11. where the α-Al mass for fraction reaches fα = 19.16%. At 592.5 °C, Figure Property diagram for AlMg5Si5 alloy. β-Al8Fe2Si2 starts to form according to the reaction L → α-Al + β-Al8Fe2Si2 and continues until For AlMg5Si5Fe1.0 (Figure 12):fα first, α-Alfβ-Al8Fe2Si2 starts to= precipitate at 603.8continues °C (L → with α-Al + L) 584.4 °C, where massalloy fraction reaches = 30.63%, 0.73%. Solidification ◦ C (L → α-Al + L) and For 12):580.9 first, α-Al starts to precipitate and AlMg5Si5Fe1.0 continues until 592.5 the α-Al mass fraction reaches = 19.16%. At 592.5 reaction L → α-Alalloy + β-Al(Figure 9°C Fe2Siwhere 2 until °C, where mass fraction reaches fat αfα = 603.8 34.75% and fβ-Al9Fe2Si2 = °C, ◦ C, 1.36%. At 580.9◦toC °C, the magnesium phase Mgreaction 2Si starts to form according to8L α-Al +592.5 β-Al 9Fe 2Si2β-Al + until β-Al8Fe 2Si2 starts form according tomass the Lreaches → α-Al Fe→ 2Si 2At and continues continues until 592.5 where the α-Al fraction f α+ =β-Al 19.16%. 8 Fe2 Si2 Mg continuing °C wherefmass fraction fα = = 73.16%, fβ-Al9Fe2Si2 = 3.19%, continues and until fMg2Si = with ◦ C, °C,2Si, where massuntil fraction αL = 30.63%, fβ-Al8Fe2Si2 0.73%. Solidification starts584.4 to form according to 568.9 the reaches reaction → α-Alreaches + β-Al Fe Si and continues 584.4 8 2 2 5.10%. Finally, the reaction L → α-Al + Mg 2Si + Al18Fe2Mg7Si10 + Si commences at 568.9 °C and reaction L → α-Alreaches + β-Al9Fe 2Si2 until 580.9 °C, where mass fraction reaches fα = 34.75% and fβ-Al9Fe2Si2 = where mass fraction fα = 30.63%, f β-Al8Fe2Si2 = 0.73%. Solidification continues with reaction solidification at 565.5 °C. 1.36%.finishes At 580.9 °C, the magnesium Mg2Si starts to form according to L → α-Al + β-Al9Fe2Si2 + ◦ C,phase L → α-Al + β-Al Fe Si until 580.9 where mass (Figure fraction13)reaches 34.75% andatf β-Al9Fe2Si2 The9 AlMg5Si5Mn1.0 alloy starts to solidify with thef α Al= 15Si 2Mn4 phase 651.6 °C = 1.36%. 2 2 Mg2◦Si, continuing until 568.9 °C where mass fraction reaches fα = 73.16%, fβ-Al9Fe2Si2 = 3.19%, and fMg2Si = according to L → Al15Si 2Mn4 +Mg L, continuing until 606.0 °C where mass fraction reaches f Al15Si2Mn4 =2 + Mg2 Si, At 580.9 C, the magnesium phase Si starts to form according to L → α-Al + β-Al 2 9 Fe2 Si 5.10%.1.83%. Finally, the reaction L →starts α-Al +form Mg2at Si606.0 + Al°C 18Fe2Mg7Si10 + Si commences at 568.9 °C and The second phase, α-Al, to according to L → Al 15Si2Mn4 + α-Al + L, and ◦ continuing until 568.9 atC565.5 where mass fraction reaches f α = 73.16%, f β-Al9Fe2Si2 = 3.19%, and finishes solidification °C. fraction reaches fα = 35.32% and at temperature 582.6 °C, mass fAl15Si2Mn4 = 2.68%. The third phase, ◦C f Mg2Si = 5.10%. Finally, the reaction L →toα-Al + Mgreaction Fe Mg +2Mn Si +commences at 568.9 2 Si + Al 7 Si 10 The AlMg5Si5Mn1.0 solidify 13) the Al 15α-Al Si 4Mg phase at and 651.6 °C Mg2Si, starts to form atalloy 582.6 starts °C according to the(Figure L 18 →with Al215Si 2Mn 4 + 2Si + L, ◦ C. and finishes solidification solidification at565.5 558.3 °C. according to L →finishes Al15Siat 2Mn 4 + L, continuing until 606.0 °C where mass fraction reaches fAl15Si2Mn4 = The AlMg5Si5Mn1.0 alloy starts to solidify (Figure 13) with the Al Si Mn phase at 651.6 ◦ C

1.83%. The second phase, α-Al, starts to form at 606.0 °C according to L → Al 1515Si22Mn44+ α-Al + L, and

◦ C where at temperature reaches fα = 35.32% and fAl15Si2Mn4 = 2.68%. Thefraction third phase, according to L →582.6 Al15°C, Si2mass Mn4 fraction + L, continuing until 606.0 mass reaches ◦ Mg 2 Si, starts to form at 582.6 °C according to the reaction L → Al 15 Si 2 Mn 4 + α-Al + Mg 2 Si + L, and f Al15Si2Mn4 = 1.83%. The second phase, α-Al, starts to form at 606.0 C according to L → Al15 Si2 Mn4 solidification at 558.3 582.6 °C. ◦ C, mass fraction reaches f = 35.32% and f + α-Al + L, and atfinishes temperature α Al15Si2Mn4 = 2.68%. The third phase, Mg2 Si, starts to form at 582.6 ◦ C according to the reaction L → Al15 Si2 Mn4 + α-Al + Mg2 Si + L, and solidification finishes at 558.3 ◦ C.

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Figure 12. Property diagram for AlMg5Si5Fe1.0 alloy. Figure 12. Property diagram for AlMg5Si5Fe1.0 alloy. Figure 12. 12. Property Property diagram diagram for for AlMg5Si5Fe1.0 AlMg5Si5Fe1.0 alloy. alloy. Figure

Figure 13. Property diagram for AlMg5Si5Mn1.0 alloy. Figure 13. Property diagram for AlMg5Si5Mn1.0 alloy. Figure Figure 13. 13. Property Property diagram diagram for for AlMg5Si5Mn1.0 AlMg5Si5Mn1.0 alloy. alloy.

The AlMg5Si5Fe1.0Mn1.0 alloy starts to solidify (Figure 14) with the Al15Si2Mn4 phase at The AlMg5Si5Fe1.0Mn1.0 AlMg5Si5Fe1.0Mn1.0 alloy alloy starts starts to to solidify solidify (Figure (Figure 14) 14) with with the the Al Al15 15Si Si22Mn Mn44 phase phase ◦at at 667.1The °C AlMg5Si5Fe1.0Mn1.0 according to L → Alalloy 15Si2Mn 4 + L, continuing until fraction reaches The starts to solidify (Figure 14)604.2 with °C thewhere Al15 Si2mass Mn4 phase at 667.1 C 667.1 °C °C according according to to L L→ → Al Al15 15Si Si22Mn Mn44 ++ L, L, continuing continuing until 604.2 604.2 °C °C where where mass mass fraction fraction reaches reaches 667.1 until ◦ fAl15Si2Mn4 =to3.59%. phase, α-Al,until starts to form at 604.2 according tof L → Al15Si=2Mn 4 + according L → AlThe Mn4 + L, continuing 604.2 C where mass°C fraction reaches 3.59%. 15 Sisecond 2second Al15Si2Mn4 = = 3.59%. 3.59%. The The phase, α-Al, α-Al, starts starts◦ to to form form at at 604.2 604.2 °C °C according according to to L LAl15Si2Mn4 → Al Al15 15Si Si22Mn Mn44 ++ ffα-Al Al15Si2Mn4 second phase, → + L, and at theα-Al, temperature mass fraction reaches The second phase, starts to582.4 form °C at 604.2 C according to fLα =→33.16% Al15 Si2and Mn4fAl15Si2Mn4 + α-Al =+ 5.12%. L, andAt at α-Al ++ L, L, and at at the the temperature temperature 582.4 °C °C mass mass fraction fraction reaches reaches ffαα == 33.16% 33.16% and and ffAl15Si2Mn4 Al15Si2Mn4 = = 5.12%. 5.12%. At At α-Al 582.4 ◦ C mass 582.4 °C, and the third phase, magnesium-rich Mgf 2αSi,= begins according to theAtreaction → the temperature 582.4 fraction reaches 33.16% to andform f Al15Si2Mn4 = 5.12%. 582.4 ◦ C,L the 582.4 °C, °C, the the third third phase, magnesium-rich magnesium-rich Mg Mg22Si, Si, begins begins to to form form according according to to the the reaction reaction L L→ → 582.4 Al15Siphase, 2Mn4 +magnesium-rich α-Al + phase, Mg2Si + L, at the temperature 566.8 to °Cthe mass fraction α = 474.47%, third Mgand begins to form according reaction L →reaches Al15 Si2fMn + α-Al 2 Si, at Al 15 Si 2 Mn 4 + α-Al + Mg 2 Si + L, and the temperature 566.8 °C mass fraction reaches f α = 74.47%, Al15Si2Mn=4 6.19% + α-Al + Mg 2Si + L, and at the ◦temperature 566.8 °C mass fraction reaches fα = 74.47%, fMg2Si = 5.60%. At566.8 566.8C°C, the fraction Al18Fe2Mg 7Si10 intermetallic starts = to6.19% form +fAl15Si2Mn4 Mg2 Si + L, andand at the temperature mass reaches f α = 74.47%,phase f Al15Si2Mn4 Al15Si2Mn4 = = 6.19% 6.19% and and ffMg2Si Mg2Si = = 5.60%. 5.60%. At At 566.8 566.8 °C, °C, the the Al Al18 18Fe Fe22Mg Mg77Si Si10 10 intermetallic intermetallic phase starts to to form form ffaccording Al15Si2Mn4 phase starts ◦ to the reaction L → Al 15 Si 2 Mn 4 + α-Al + Mg 2 Si + Al 18 Fe 2 Mg 7 Si 10 + L and solidification and f Mg2Si = 5.60%. At 566.8 C, the Al18 Fe2 Mg7 Si10 intermetallic phase starts to form according to the according to the the reaction reaction L L→ → Al Al15 15Si Si22Mn Mn44 ++ α-Al α-Al ++ Mg Mg22Si Si ++ Al Al18 18Fe Fe22Mg Mg77Si Si10 10 + +L L and and solidification solidification according to finishes at 560.7 °C. reaction L→ Al15 Si2 Mn4 + α-Al + Mg2 Si + Al18 Fe2 Mg7 Si10 + L and solidification finishes at 560.7 ◦ C. finishes at 560.7 °C. finishes at 560.7 °C.

Figure 14. Property diagram for AlMg5Si5Fe1.0Mn1.0 alloy. Figure 14. 14. Property Property diagram diagram for for AlMg5Si5Fe1.0Mn1.0 AlMg5Si5Fe1.0Mn1.0 alloy. alloy. Figure Figure 14. Property diagram for AlMg5Si5Fe1.0Mn1.0 alloy.

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4. Discussion The modification of the microstructure caused by melt stirring requires discussion, especially the transformation from a dendritic structure to rosettes with minor dendrites and, occasionally, spheroids. Here, the measured parameters (dendrite arm spacing, specific surface, length and number density of intermetallics and eutectic spacing) will be analyzed in comparison to literature data, and the various effects of stirring on the studied alloys and growing phases will be discussed. The non-dendritic morphologies are mostly explained by dendrite fragmentation in presence of forced convection [1]. Mentioned by Flemings [1], the mechanism responsible for fragmentation may be: (a) dendrite arm fracture; (b) remelting of the dendrite root; and/or (c) recrystallization caused by fluid flow inducing mechanical stress. 4.1. Rosettes The rosette morphology and its growth were studied by Mullis [9] using the cellular automaton method. Mullis observed that forced convection induces rotation of the dendrite tip caused by thermal and solutal advection. The size of the crystal was defined by the magnitude of the bending parameter, and eventually such a rosette might form without external mechanical interaction coming from flow. Birol [10] conducted experiments with metal alloys flowing along a cooled plate, and proved that during flow, crystals have the possibility to collide with each other and coalesce to grow as agglomerates. Therefore, in the final microstructure may be observed fully shaped dendrites, α-Al crystals formed as rosettes, and some globular grains. Niroumand and Xia [11], on microstructure of AlCu10 solidified by stirring, observed many round crystals shaped as clusters. With the help of classical LOM and in 3D, they found out that globular crystals of clusters are connected to each other in three dimensions, even though on a 2D micrograph they can look separate. This means that dendrite fragmentation and subsequent agglomeration influences microstructure formation only weakly, and suggests that rosette-shaped clusters are ripened arms of deformed dendritic crystals. 4.2. Spheroids Ji et al. [12] theoretically analyzed the stability of the solid–liquid interface during solidification under forced convection. The numerical analysis indicated that the growth morphology may change from equiaxed dendrite to spheroid via rosette when shear rate and turbulence increases significantly. Das et al. [13], using Monte Carlo simulations, showed that a tendency for dendritic growth was reduced and globularization of the primary phase in the melt is caused by the rotation of the solid particle and elimination of constitutional undercooling through reduction of the thermal/solutal diffusion layers at the solid–liquid interface. In the experiments on rheomoulding, Ji et al. [12] produced a spherical morphology, instead of a rosette or dendrite, by high shear rates. Birol [14] obtained globular structures with forced convection and internal cooling, and proposed that stirring caused the uniform chemical composition near the solid–liquid interface. The results of Li et al. [15] on succinonitrile (SCN)-5% water also supports the hypothesis that nondendritic microstructure comes from natural nucleation and globular growth. Martinez and Flemings’ [16] study on AlCu4.5 alloy found that intensive convection may lead to growth in spheroidal form just below the liquidus temperature. In the current experiment, rosettes seem to form as an effect of rotation of the dendrite tip during growth, and are also ripened arms of deformed dendritic crystals. This explanation is especially supported by the fact that many dendrites were found in addition to overwhelming rosettes. The spheroids seem to be part of the dendrites, because the formation of spheroids requires intensive stirring, causing a very high shear rate.

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4.3. Dendrites In columnar solidification, the microstructure is described with primary dendrite arm spacing λ1 [17–19] and by secondary dendrite arm spacing λ2 . For equiaxed grain morphologies resulting from free growth (equiaxed solidification), metallurgists traditionally measure λ2 and grain size, or number of grains or distance between grains [17]. Recently, the additional measure of specific surface Sv of dendrites was proposed and applied by Voorhees et al. [20]. Secondary dendrite arms start to grow very close to the dendrite tip, initially as perturbations, which develop into cell-like structures and further grow as separate arms located parallel to each other with similar sizes varying with time. The growth by natural thermal convection and solutal fluctuations leads to arms of various sizes, where smaller arms dissolve and the larger overgrow smaller. The distance between secondary arms λ2 is determined by the coarsening process [21,22]. Many mathematical models [19,23,24] were developed based on the concept that dendrite coarsening is diffusion controlled, but mostly they calculate λ2 with a simple power law as a function of local solidification time t: λ2 = c1 · tn1 (1) where n1 = 0.33 for diffusive mass transport and 0.48 for convective regime [25]. c1 is a coefficient containing material constants (diffusion coefficient, concentration, etc.) and is given different forms by Kattamis and Flemings [23], Voorhees and Glicksman [26], and by Mortensen [27]. Bouchard and Kirkaldy [19,28] proposed a formula based on local cooling rate: λ2 = c2 · R−n2

(2)

where c2 = coefficient and n2 = 0.33 [19] or is in the range 0.22–0.33 for AlSi alloys [28] with various Si content. Mullis [29] found that flow from the tip of the secondary arm towards the root will enhance the ripening rate, while flow in the opposite direction will reduce the ripening rate. For flow aligned along primary trunk, all secondary arms will experience a transverse flow, which will enhance ripening. Because of the four-fold symmetry of dendrites, the flow effect on secondary spacing λ2 could be small (increase in ripening rate). Diepers et al. [30] developed a model for Ostwald ripening, which holds that in the dependence of secondary spacing λ2 on solidification time, the exponent n1 changes from 0.33 for diffusive ripening to 0.5 for flow-governed dendrite ripening because of the flow increase caused by coarsening. The simulation results are in agreement with considerations of Ratke and Thieringer [31] on convective ripening theory and experimental results of Steinbach and Ratke [25] on directional solidification of A357 alloy (AlSi7Mg0.6). Steinbach and Ratke [25] found that λ2 increases continually with increasing magnetic induction (increasing flow velocities) (e.g., by t = 1250 s from 80 µm to 140 µm). According to power law expressions of solidification time and by the c1 coefficient from [25], here the change in n1 exponent is from 0.36 in a solute-controlled system with natural convection to 0.48 by flow caused by electromagnetic stirring (6 mT). In the current equiaxed solidification experiments, with increasing flow, secondary spacing λ2 decreased from 87 to 77 µm (−12%) for AlSi5Fe1.0 alloy, but remained almost unchanged for the other alloys, which contained Mg and therefore Mg2 Si phases. Stirring decreased solidification time (Table 1), and according to the Equation (1) for λ2 , this should mean a decrease in secondary spacing λ2 . Lower secondary spacing was measured for only AlSi5Fe1.0 alloy, while for other alloys it remained almost unchanged. For the AlSiFe1.0 alloy solidified without stirring, λ2 = 87 µm, but with flow, based on measured solidification time, calculated secondary spacing (λ2 = 82 µm) is still bigger than measured spacing (λ2 = 77 µm). In order to reach the measured value of 77 µm, the exponent n1 in Equation (1) should be lower, even lower than n1 for diffusive ripening (without stirring) [25]. These results would require a decrease from 0.33 to 0.318; this is opposed to literature [25,30], suggesting an increase in n1 caused by convective ripening. For other alloys

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(containing Mg), according to Equation (1), shorter solidification time results in a smaller secondary spacing λ2 , but only slightly smaller. In order to reach, by calculation, the measured λ2 values, the exponent by stirring should be slightly higher, increasing from 0.33 for diffusive mass transport to the range 0.334–0.342 for convective ripening, and that is much less than the range 0.47–0.50 found in literature [25,30] for directional solidification. This behavior of exponents suggests a lack of forced convection by electromagnetic stirring for Mg-containing alloys, or its significant reduction during coarsening of secondary arms. Instead of n2 = 0.33 in Equation (2), the exponent would be 0.342 for AlSi5Fe1.0 alloy and 0.326–0.318 for Mg-containing alloys. This is in contrast to [25,30], where secondary spacing λ2 increased considerably under flow during directional solidification Much of research devoted to dendritic structure was focused on measuring secondary dendrite arm spacing λ2 , but this measure does not provide information on the complexity of the dendritic structure. Marsh and Glicksman [32] found out that, despite drastic morphological changes from dendritic structure to spheroidal during the coarsening processes, the solidification time t dependence of specific interfacial area Sv was always described as: SV ∼ t1/3

(3)

Kasperovich and Genau [33] studied mushy zone coarsening for AlCu30 alloy, and Sv was found to be in the range 0.04–0.22 µm−1 for holding times from 20 to 500 min. Sv [33] decreased from 0.077 to 0.035 µm−1 (for holding time of 200 min) with increasing forced convection generated by 6 mT RMF. For all alloys studied here, Sv decreased under fluid flow, meaning that the dendrites or rosettes are rounder with stirring. Based on Equation (3) and the solidification time measured, the calculated decrease in Sv is smaller than measured. For AlSi5Fe1.0 alloy with stirring, the measured Sv is 0.019 µm−1 , while the calculated value is 0.022 µm−1 , and similarly for AlMg5Si5, measured Sv is 0.024 µm−1 and calculated Sv is 0.0341 µm−1 . The difference between measured (without stirring) and calculated (with stirring) values of Sv seems to present well the effect of stirring on microstructure. 4.4. Eutectics Eutectic morphologies [17,18] are characterized by the simultaneous growth of two (or more) phases from liquid, where the exchange of solute between two simultaneous phases occurs via transport in the liquid phase, and transport may be strongly influenced by forced convection. The eutectic spacing λE was determined by Jackson and Hunt [18,34]: λE = c3 · V −0.5

(4)

where c3 = coefficient and V solidification front velocity. Steinbach and Ratke, in directional solidification of A357 alloy, found that fluid flow increased eutectic spacing (e.g., for solidification velocity 90 µm/s, spacing increased from 3 to 7 µm), and also reported the reduction in λE with increasing solidification velocity according to the well-known Jackson and Hunt relation (4). In directional solidification of AlSi5/7/9Fe0.2/0.5/1.0 alloys [35], the eutectic spacing did not show any clear correlation with fluid flow, which was in agreement with earlier results by Sous [36]. As an example, eutectic solidification for AlSi5Fe1.0 [8] occurs at the final eutectic reaction at 575 ◦ C. During directional solidification with a temperature gradient of 3 K/mm, the mushy zone is about 18 mm wide [8], while the eutectic zone is about 1–3 mm wide and the flow deep between dendrites might be reduced. However, in the current equiaxed solidification with a temperature gradient of 0.143–0.214 K/mm and cooling rate of 0.112–0.626 K/s, equiaxed dendrites grew freely in the mush, moving in the liquid with similarly growing eutectic phases, so one cannot expect that convection is diminished by dendrites. From Table 1, it is clear that for AlSi5Fe1.0, forced convection increases the eutectic spacing λE by about 9%. For the other alloys, the change in spacing is similar in magnitude, but both increases and decreases were observed. For alloys with either Fe or Mg, we saw

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an increase, while for alloys containing both Fe and Mg, flow reduces the spacing. It seems that some common influence of both elements causes such an effect on growing phases. The unclear effect of stirring on the eutectic spacing may come from the small amount of eutectic and the resulting small number of measurements. The mass fraction for eutectic reached 4.67 wt. % in AlSi5Fe1.0 and only about 1.5–2.1 wt. % in Mg-containing alloys, whilst Mg2 Si reached about 7.8 wt. % [7]. The Mg2 Si phases also seem to be affected by convection that causes about a 5%–10% increase in spacing λMg2Si . There are no literature data concerning fluid flow effects on Mg2 Si phases. 4.5. Intermetallics The shortening of β phases observed in this work is consistent with results by Nafisi et al. [37], where, for specimens solidified in a sand mold, stirring decreased average length from Lβ = 9–10 µm to 7–8 µm, and respectively in a copper mold decreased weakly in the range Lβ = 4.5–5 µm. Fang et al. [38], for LM25 alloy (AlSi7Mg0.2–0.6Fe0.5), found a shortening of β phases (from 75 to 15 µm), and for LM24 alloy (AlSi8Cu3Fe1.3), β-Al5 FeSi with lengths 95–110 µm were completely eliminated. Steinbach et al. [39] observed that flow causes the formation of a eutectic center region in directionally solidified AlSi7Fe1.0 alloy and growth of about 280 µm long Fe platelets, while the same system without stirring precipitated shorter β phases of about 160 µm long. In [40], forced convection in directional solidification of AlSi5/7/9Fe0.2/0.5/1.0 alloys caused about 20% shortening of β phases in the dendritic region of the specimen and 9% increase in the eutectic reach center. The histograms displayed that the most common are β phases with lengths from 5 to 40 µm, and that convection led to a higher number of short phases, causing a smaller average length Lβ . Nafisi et al. [37], for AlSi6.8Fe0.8 alloy, showed an increase under stirring in the number density for sand mold from nβ = 600–1200 mm−2 to nβ = 800–2600 mm−2 , and copper mold from nβ = 5000–13,000 mm−2 to nβ = 5000–14,000 mm−2 . The data in [40] clearly showed an increase in number density nβ , stronger in the eutectic center (42%) than in the outer part (17%) of the specimens. The analysis of the stirring effect in directional solidification of AlSiFe alloys [40] has pointed to shortening of β phases as a complicated effect of forced convection, solute segregation, dendrites, and intermetallic morphology. Fragmentation and partial dissolution of β-Al5 FeSi acting as nucleation sites were explained as a cause of higher number density nβ and a decrease in β platelet length Lβ . 4.6. Solidification by Stirring For AlSi5Fe1.0 alloy (Table 1), stirring decreased secondary spacing λ2 from 87 to 77 µm (−12%), changed Sv from 0.023 to 0.019 µm−1 , shortened β phases from 71 to 57 µm (−20%), increased its number density from 109 to 160 mm−2 , and increased eutectic spacing from 8.5 to 9.5 µm. According to the property diagram (Figure 9), α-Al started to form first at 627.4 ◦ C, next was β-Al5 FeSi at 608.4 ◦ C, and, finally, eutectics precipitated at 575 ◦ C. Flow determined growth of α-Al, causing rosettes formation, minor dendrites, and, occasionally, spheroids, and changed secondary spacing λ2 . β phases started to grow when the solid mass fraction of α-Al reached 40%, which means the shortening of β occurred between well-formed dendrites (at least precipitated in 40%). The occurrence of such a large amount of α-Al phase probably decelerated flow and might support mechanical interactions between α-Al dendrites and moving β. Eutectics precipitated at the end and the flow was probably strongly reduced by dendrites and β, so the increase in eutectic spacing λE was influenced by local small flows between solid phases. For AlMg5Si5 alloy (Table 1), stirring produced a negligible change in measured λ2 from 67 to 68 µm, decreased Sv from 0.036 to 0.024 µm−1 , and decreased eutectic spacing from 9.8 to 8.8 µm. According to Figures 10 and 11, α-Al started to form first at 608.0 ◦ C, next was the Mg2 Si phase at 581.7 ◦ C, and, finally, eutectics precipitated at 558.6 ◦ C. The growth of α-Al seems to be undisturbed by Mg2 Si until 581.7 ◦ C, and until the fraction of α-Al f α = 37.28%, mostly rosettes were formed. The arm-ripening was influenced by fluid flow for AlSi5Fe1.0 alloy, but for AlMg5Si5 stirring was probably diminished by Mg2 Si phases, and convection changed secondary dendrite arm spacing λ2 only slightly.

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For AlMg5Si5Fe1.0 alloy (Table 1), forced convection modified λ2 only from 60 to 62 µm, Sv from 0.034 to 0.026 µm−1 , decreased eutectic spacing from 11.7 to 6.8 µm, and changed Mg2 Si spacing from 11.7 to 12.7 µm. According to the property diagram (Figure 12), α-Al started to form first at 603.8 ◦ C, next was β-Al5 FeSi phase at 592.5 ◦ C, next was Mg2 Si at 580.9 ◦ C, and, finally, eutectics precipitated at 565.5 ◦ C. The growth of α-Al seems to be undisturbed by β-Al5 FeSi until 592.5 ◦ C and by Mg2 Si until 580.8 ◦ C, and rosettes formed but λ2 changed only weakly. A 5% shortening of the β-Al5 FeSi phase and increase in number density nβ (14%) seems to be affected by growing Mg2 Si, in comparison to AlSi5Fe1.0 alloy. Initially, β grew between α-Al, but further in the presence of Mg2 Si, which seems to have decreased forced convection and reduced the possible modification of β-Al5 FeSi. The flows in almost-solidified alloys are probably very small, but still have potential to influence the Mg2 Si and eutectic spacing. For AlMg5Si5Mn1.0 alloy (Table 1), when stirring is applied, there is no change in measured λ2 , a decrease in Sv from 0.036 to 0.028 µm−1 , an increase in eutectic spacing from 7.5 to 8.7 µm, and an increase in Mg2 Si spacing from 11.6 to 12.0 µm. For Mn–rich phases, flow decreased the average overall dimension by about 9% from 299 to 273 µm and increased number density nMn from 20 to 27 mm−2 . According to the property diagram (Figure 13), the first thing to form was the Mn-rich phase Al15 Si2 Mn4 at 651.6 ◦ C, next was α-Al at 606.0 ◦ C, third was Mg2 Si at 582.6◦ C, and, finally, eutectics precipitated at 558.3 ◦ C. About 60% (Figure 13) of Mn-rich phases grew freely in liquid melt and, under forced convection, formed as smaller structures. Similar to previous Mg-containing alloys, calculated λ2 changed only weakly under reduced forced convection. Growth of initial Mn-rich phases influenced flow very slightly and allowed for rosette formation. Mn-rich phases precipitated gradually during solidification in liquid melt as complicated structures, being overgrown by dendrites (Figure 8a) and later grew between dendrites and Mg2 Si (Figure 8b). For AlMg5Si5Fe1.0Mn1.0 alloy (Table 1), stirring produced no change in measured λ2 , a decrease in Sv from 0.035 to 0.028 µm−1 , an increase in eutectic spacing from 8.1 to 9.1 µm, and in Mg2 Si spacing from 11.3 to 12.1 µm. Forced convection produced the growth of α-Al rosettes, β phases with unmodified length Lβ , but with increased number density nβ . Mn-rich phases growing early during solidification, before or with α-Al, have not influenced melt flow, or the influence is small. Convection was therefore not disturbed in the early stage of solidification, causing formation of rosettes. Flow disturbed by Mg2 Si during the late stage of solidifications influenced the ripening process slightly and resulted in only a small difference between the secondary spacing λ2 calculated from (1) and the measured value. Mg2 Si phases, reducing flow, clearly limited shortening of β-Al5 FeSi phases, but did not stop the increase in number density nβ . Flow decreased secondary dendrite arm spacing λ2 in AlSi5Fe1.0 alloy (without Mg) and only very slightly in Mg-containing alloys. Mg2 Si diminished flow and convective ripening while supporting diffusive mass transport. α-Al grew as rosettes with modified Sv in all studied alloys. The decrease in Sv became weaker with growing complexity of alloy (growing number of elements), and for AlMg5Si5, Sv decreased by 33%, whilst decreasing by 20% for AlMg5Si5Fe1.0Mn1.0 alloy. The current study confirmed the results [40] for directional solidification of AlSi5Fe1.0 alloy, where forced convection decreased average length Lβ of β phases by about 20% and increased number density nβ by about 17% under comparable rotational speed of 2 s−1 . Completely new is that the flow effect was observed in equiaxed growth by low temperature gradient and low cooling rate, where there is no possibility of phases remelting. This suggests shortening of β as an effect of mechanical fragmentation or mechanism determined by solute distribution changed under flow. For AlSi5Fe1.0 alloy, flow caused a 20% decrease of average length Lβ of β phases and 47% increase in number density nβ , whilst for AlMg5Si5Fe1.0, Lβ decreased only 5% and nβ increased 14%. Flow effect on β is weaker in Mg-containing alloys, as a result of reducing the flow by Mg2 Si. For AlMg5Si5Fe1.0Mn1.0, Lβ decreased only 2% and nβ increased 77%. By measured length Lβ , number density nβ , and non-inspected thickness of β, a larger number of shorter phases by the same

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Fe quantity were included in β-Al5 FeSi, which means that β must be thinner. This mechanism is unclear and needs more experimental and analytical investigations. 5. Conclusions 1. Electromagnetic stirring caused a transformation in microstructure from equiaxed dendritic to rosettes with minor dendrites as an effect of rotation of dendrite tips and ripening arms of deformed dendrites. Occasionally occurring spheroids seem to be part of deformed dendritic crystals. 2. Stirring caused a decrease in solidification time for all alloys. Contrary to what has previously been reported in the literature, flow decreased secondary dendrite arm spacing λ2 for AlSi5Fe1.0 alloy. Even considering the lower solidification time, the exponents in the power law equation for λ2 changed values: n1 from 0.33 to 0.318 and n2 from 0.33 to 0.342. For Mg-containing alloys, measured secondary spacing λ2 was almost unchanged, but considering lower solidification time and lower calculated λ2 , the exponents should change: n1 from 0.33 to 0.334–0.342, and n2 from 0.33 to 0.326–0.318, less than in previous works. 3. Specific surface Sv decreased under forced flow and clearly signaled the modification of microstructure caused by stirring. 4. Forced convection decreased the length of β-Al5 FeSi (20%) and increased number density (47%) in AlSi5Fe1.0 alloy in equiaxed solidification, in accordance with literature data for directional solidification. In Mg-containing alloys, the changes were much smaller. 5. Stirring decreased length of Mn-rich phases (9%) and increased number density (35%) in AlMg5Si5Mn1.0 alloy. 6. Melt flow changed eutectic spacing λE depending on alloy composition, while λMg2Si increased weakly for all alloys. 7. Mg2 Si phases reduced fluid flow generated by RMF and consequentially reduced shortening of β-Al5 FeSi phases, diminished secondary arm ripening caused by forced convection, and supported diffusive ripening. Mg2 Si did not disturb transformation from dendrites to rosettes under flow. 8. Shortening of β-Al5 FeSi phases caused by stirring occurred in equiaxed solidification without remelting, probably by mechanical fragmentation, modified solute distribution, and additional nucleation sites. 9. Stirring application and efficiency in microstructure modification depends on chemical composition, precipitating phases (e.g., Mg2 Si), and growth sequence of phases in alloys. Acknowledgments: The research leading to these results has received funding from the People Programme (Marie Curie Actions) of the European Union’s Seventh Framework Programme (FP7/2007-2013) under REA grant agreement No. PCIG13-GA-2013-613906. Thanks are due to Mark Morton from Hydro Aluminum High Purity GmbH (Grevenbroich, DE) for the aluminum. Conflicts of Interest: Funding sponsor is European Union’s Seventh Framework Programme (FP7/2007-2013). The funding sponsors had no role in the design of the study; in the collection, analysis, or interpretation of data; in the writing of the manuscript, and in the decision to publish the results.

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