Effect of (Mn + Cr) addition on the microstructure and thermal stability of spray-formed hypereutectic Al–Si alloys

Materials Science and Engineering A 527 (2009) 85–92

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Effect of (Mn + Cr) addition on the microstructure and thermal stability of spray-formed hypereutectic Al–Si alloys L.G. Hou a , H. Cui b , Y.H. Cai a , J.S. Zhang a,∗ a b

State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China School of Materials Science and Engineering, University of Science and Technology Beijing, Beijing 100083, China

a r t i c l e

i n f o

Article history: Received 26 December 2008 Received in revised form 17 July 2009 Accepted 22 July 2009

Keywords: Spray Forming technology Hypereutectic Al–Si alloys Microstructure Thermal stability Transformation

a b s t r a c t Microstructures and thermal stability of hypereutectic Al–Si alloys with or without (Mn + Cr) addition, prepared via Spray Forming technique, are studied and compared with traditional cast alloys with same composition, using scanning electron microscopy with energy diffraction spectrum, X-ray diffraction, transmission electron microscopy and differential scanning calorimeter. The results show that the Febearing and primary silicon phases in SF-3C alloy can be refined to less than 10 ␮m, especially in SF-MC21 alloy the Fe-bearing phase is refined into uniformly distributed ␣-Al(Fe,Mn,Cr)Si phase particles with sizes smaller than 5–6 ␮m, contributing to the decrease/elimination of the deleterious effect of needlelike Fe-bearing phases. The results of different heat treatments show SF-MC21 alloy possesses excellent thermal stability than SF-3C alloy which is unstable below 750 K for the coarsening of ␤-Al5 FeSi phase and formation of Al7 Cu2 Fe phase. The study indicates that both the existence of thermodynamically stable ␣-Al(Fe,Mn,Cr)Si particles and the increase of solidus temperature of SF-3C alloy induced by adding (2Mn + 1Cr) elements contribute to the high thermal stability of SF-MC21 alloy. Contemporarily, combined the phase reactions or transformation occurred during the melting and solidification processes of both spray-formed hypereutectic Al–Si alloys, the microstructure formation of spray-formed alloys is discussed. © 2009 Published by Elsevier B.V.

1. Introduction With rapid development of new vehicles with preeminent properties, reducing the weight of new vehicles will undoubtedly become the most direct and beneficial way to reduce energy consumption [1]. Lightweight materials with low density, such as Al/Mg/Ti alloy, composites, etc., may become the main direction of research in materials science to reduce vehicles weight and energy consumption [2–4]. Al–Si alloys as the main cast Al alloys have been mainly applied in engine blocks, cylinder heads and wheels. Particularly, the traditional materials, cast irons and steels, used for heatand wear-resistant parts (e.g. pistons, cylinder liners, rotors, etc. [5–11]), can be substituted by hypereutectic Al–Si (HEAS) alloys for its good castability, low density, low coefficient of thermal expansion (CTE) and excellent wear resistance [12,13]. Traditional cast (TC) HEAS alloys contain coarse primary Si and needle-like Fe-rich phases, both inducing fracture and failure for stress concentration effects [10,14,15]. Microstructural optimization, especially the coarse primary Si and needle-like Fe-bearing phases, becomes an important field of study and several strategies

∗ Corresponding author. Tel.: +86 10 62334717; fax: +86 10 62332508. E-mail address: [email protected] (J.S. Zhang). 0921-5093/$ – see front matter © 2009 Published by Elsevier B.V. doi:10.1016/j.msea.2009.07.041

have been proposed in order to achieve high strength and good ductility. Firstly, adding some modifiers into alloy melts to increase the heterogeneous nucleation and depress the growth of primary Si phase so as to refine the primary Si phase, e.g., Na [16,17], Sr [18], Nd [19], Ca [20], P [21–23], Sb [24], RE [21,25] elements, especially P and RE elements are more effective [21]. Secondly, adding some trace elements such as Mn [14,26–29], Cr [14,29–31], Be [32], Sr [33], Ca [20], Co [14,31], K [34], Li [35], named “neutralizer”, can eliminate the negative effects of needle-like Fe-bearing phase in cast Al alloys [14,26,27,36,37]. These “neutralizers” have been already applied in cast 3000, 6000 series Al alloys [38–41] and can be effectively used in hypoeutectic and eutectic Al–Si alloys with low Fe content (<1 wt.%). But the mechanical properties of high-Fe HEAS alloys, manufactured by TC process, still cannot be improved fundamentally even though the morphology of Febearing phase can be modified by adding “neutralizer”. Thirdly, the increase of cooling rate during solidification can increase the number of nucleus and refine the microstructures. Rapid solidification technique, e.g., Rapid Solidification Powder Metallurgy (RS/PM) [42,43], Semi-Solid Process (SSP) [44–46], Spray Forming (SF) [47–57], etc., have been studied and applied in practice, particularly, the SF with high cooling rate (103 –104 K/s) can refine the microstructures and eliminate the macrosegregation, accompanying uniformly distributed second phase particles [47]. Presently,

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SF process has become an important method for the preparation of HEAS alloys with outstanding properties [49–58]. With further densification process, the as-deposited HEAS alloys could be fully compacted and excellent properties can be obtained, such as high fracture toughness and high elevated temperature strength, good fatigue property and machinability, and applied in practice successfully in Germany, Japan and others. Yang et al.’s research [58] showed that ternary Fe-bearing phases can be optimized via the composite effects of Mn addition and SF process. In addition, some researches showed that Cr can modify the morphology of needle-like Fe-bearing phase as Mn did in Al–Si alloys [14,31–33]. So, adding Cr to HEAS alloys containing Fe might optimize the microstructure. However, the effects of Cr or (Mn + Cr) addition on the microstructures and properties of HEAS alloys containing Fe have not been investigated and reported to the present, therefore, the composite effects of (Mn + Cr) addition and SF process have been investigated presently aiming to obtain excellent microstructures. 2. Experimental procedure 2.1. Materials The HEAS alloys with/without (2Mn + 1Cr, wt.%) addition had been prepared by TC and SF processes with basal composition: Al–25Si–5Fe–3Cu (wt.%) (Signed as 3C, also the alloy with (2Mn + 1Cr) addition was signed as MC21). The raw materials included pure Al (99.7% purity), Al–40Si and Al–50Cu master alloys and commercial Fe- (75 wt.% Fe), Mn- (75 wt.% Mn), Cr-additives (60 wt.% Cr). The others were refiners for Al alloys except for pure Fe, Mn and Cr elements in the three additives. 2.2. Traditional cast process Both 3C and MC21 alloys were prepared by TC process. Firstly, after completely melting of the pure Al, Al–40Si and Al–50Cu master alloys in an intermediate frequency induction furnace, some commercial Fe-additives were added into the Al–Si–Cu alloy melts and held for 20–30 min among 1073–1123 K for the fully melt-

Table 1 The parameters of SF process. Atomization gas

N2

Atomization pressure/MPa Deposition distance/mm Melting temperature/K Diameter of Tube/mm Spray angle/◦ Eccentricity/mm

0.6–0.8 390–420 1123–1203 3 22 16–20

ing of Fe-additives. For alloy with (2Mn + 1Cr) addition, Fe-, Mnand Cr-additives were orderly added into the melts and same heat preservation treatment as done to 3C alloy was used. And the temperature of alloy melts should be reduced to ∼1073 K and the alloy melts were subsequently poured into graphite mould with dimensions Ø30 mm × 100 mm for the ingots. For the liquidus temperature of Al–Si binary alloy with 25 wt.% Si was about 1053 K [14], the temperature of heat preservation treatment should be 50–70 K high than liquidus temperature for the entire melting of all the alloying elements. 2.3. Spray Forming process The melting process was same with Section 2.2, but the degree of superheat during melting should be 100–150 K. And the melts should be held for 10–15 min among 1153–1203 K to eliminate any un-melted alloying elements or high temperature compounds, which could reduce the fluidity of alloy melts and affect the atomization process. The atomization and deposition processes were finished in selfmade SF facility and the parameters of SF process were shown in Table 1. An annular convergent–divergent nozzle had been utilized to get a spray of micro-sized droplets through a melt delivery tube from the alloy melts. The melts were atomized into droplets by Nitrogen gas and the droplets were deposited onto a rotating and withdrawing collector copper substrate below the atomizer to form a rapid solidified cylinder preform with dimensions Ø130 mm × 180 mm. In order to reduce oxidation of the spray, the atomization chamber had been filled with inert gas.

Fig. 1. Microstructures of spray-formed alloys (a and c) and their over-sprayed powders (b and e). (a and b) 3C, (c and e) MC21, (d) high magnification of (c).

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All the alloy melts during the TC and SF processes were not treated by degassing, refinement or modification process. 2.4. Microstructures analysis and phase identification The specimens were cut from cast and spray-formed alloys, and polished with standard metallographic procedures by using 2.5 ␮m diamond polishing paste. The microstructures were investigated by field emission scanning electron microscope (FE-SEM, ZEISS SUPRA55) with energy dispersive spectroscopy (EDS) as well as over-sprayed powders, and the composition of various phases reported were identified by Electron Probe Microanalyzer (EPMA, JEOL JXA-8100). 5–6 measurements were taken to obtain an average value. The phase components were determined by PHILIPS APD-10 X-ray Diffractometer (XRD) which traces were obtained with monochromatic Cu K␣ radiation with a 0.02◦ step size. The specimen for transmission electron microscopy (TEM, HITACHI H800 TEM, 200 kV) was cut away from the as-deposited preform, mechanically polished, and ion milled in a dual ion-beam miller at a current of 0.5 mA. 2.5. Thermal stability The present HEAS alloys would be applied in high temperature and wear conditions, and the thermal stability at high temperature should be highlighted. Then spray-formed HEAS alloys had been treated with different times at different temperatures and the corresponding microstructures and phases had been studied by SEM, XRD, and EPMA. The parameters of heat treatments showed as follows:

Fig. 2. XRD spectrums of spray-formed alloys (a and c) and their over-sprayed powders (b and d). (a and b) 3C, (c and d) MC21.

SF-3C alloy: (700 K, 750 K, 803 K) × 5 h; SF-MC21 alloy: (753 K, 783 K, 813 K) × 10 h, 813 K × 3 0 h. All the temperatures above were selected based on the DSC curves (shown in the text). Also, the reactions or transformations between phases in sprayformed HEAS alloys during remelting and solidification processes were studied by NETZSCH STA 409 C/CD DSC equipment. The load receptor for the specimen was Al2 O3 pan and the heating/cooling rate was 10 K/min. The weight of each specimen was 5–6 mg. During the experimental procedure, the specimen chamber was vacuum and protected by Ar gas with flow rate 20 mL/min. 3. Results and discussion 3.1. Microstructural evolution induced by (Mn + Cr) addition Fig. 1(a) shows that SF-3C alloy contains white short-rod phase and bright white phase besides Al matrix and grey Si particles. The XRD result in Fig. 2(a) shows ␤-Al5 FeSi, ␦-Al4 FeSi2 and Al2 Cu phases are present. Based on the EPMA results, it can be deduced that the white short-rod phase might be ␤-Al5 FeSi/␦-Al4 FeSi2 phase and the bright white phase is Al2 Cu phase. It is difficult to differentiate these two Fe-bearing phases each other for the exceeding similarity on their sizes and morphologies. But TC-3C alloy contains coarse needle-like compounds (hundreds of micrometers) and fine white phase except ␣-Al matrix and coarse primary Si phase (Fig. 3(a)). According to EPMA (Table 2) and XRD (Fig. 4(a)) results, the white phase should be Al2 Cu phase and the coarse needle-like compounds might be ␦-Al4 FeSi2 phase (its stoichiometric is Al3.46 Fe1.04 Si2.51 which is close to ␦-Al4 FeSi2 .) but not ␤-Al5 FeSi phase, because no diffraction peaks of ␤-Al5 FeSi phase are detected. The sizes and morphologies of different phases in SF- and TC-3C alloys are gathered in Table 4. Compared to TC-3C alloy (Fig. 3(a)), both primary Si

Fig. 3. The microstructures of cast alloys, (a) 3C (BSE), (b) MC21 (BSE), and (c) MC21 (SE).

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Table 2 Composition of some compounds in cast alloys using EPMA (at.%). Elements

Al

Si

Fe

Mn

Cr

Ref.

Phase

Needle-like phase (Fig. 3(a))

49.43 49.0 70.25

35.85 35.0 13.52

14.91 16.0 8.51

– – 4.35

– – 3.46

Present [59] Present

␣-Al15 (Fe,Mn,Cr)3 Si2

Phase

Skeletal phase (Fig. 3(b) and (c))

␦-Al4 FeSi2

Table 3 The composition of compounds in spray-formed alloys by EPMA (at.%). Elements

Al

Si

Fe

Mn

Cr

Ref.

Needle-like phase (Fig. 1(e))

49.06 49.0 67.49 72.93

36.60 35.0 18.60 12.09

10.57 16.0 12.73 8.65

3.21 – 1.07 3.74

0.58 – 0.06 1.75

Present [59] Present Present

Plate phase (Fig. 1(d)) Particulate phase (Fig. 1(c) and (d))

Fig. 4. XRD spectrums of cast alloys: (a) 3C and (b) MC21.

and Fe-bearing phases can be refined (<10 ␮m) via SF process, contributing to improving mechanical properties. Figs. 1(a) and 2(b) show ␦-Al4 FeSi2 phase are also present in over-sprayed powders without ␤-Al5 FeSi phase. Fig. 1(c) shows the coarse Fe-bearing phase in TC-3C alloy has been further refined into granular ␣ Fe-bearing phase (∼5–6 ␮m), even nano-sized (Fig. 1(d)) with (2Mn + 1Cr) addition while the sizes and morphology of primary Si phase are same with SF-3C alloy, shown in Table 4. But a few plate phases (∼20 ␮m size) appear (Fig. 1(c)). Based on XRD (Fig. 2(c)) and EPMA results (Table 3), the particles might be ␣-Al(Fe,Mn,Cr)Si phase and its stoichiometric

␦-Al4 (Fe,Mn,Cr)Si2 ␤-Al5 (Fe,Mn,Cr)Si ␣-Al15 (Fe,Mn,Cr)3 Si2

is close to ␣-Al15 (Fe,Mn,Cr)3 Si2 while the short plate phase is ␤Al5 (Fe,Mn,Cr)Si phase. Also a number of ␣-Al(Fe,Mn,Cr)Si particles appear in over-sprayed powders accompanying many needle-like Fe-bearing phases (Fig. 1(e)), which might be ␦-Al4 (Fe,Mn,Cr)Si2 /␤Al5 (Fe,Mn,Cr)Si phase deduced from XRD (Fig. 2(d)) and EPMA (Table 3) results. It indicates that the main Fe-bearing phases in over-sprayed powders, e.g., ␣-Al(Fe,Mn,Cr)Si, ␦-Al4 (Fe,Mn,Cr)Si2 or ␤-Al5 (Fe,Mn,Cr)Si phase (Table 4), can be substituted by ␣-Al(Fe,Mn,Cr)Si phase in as-deposited preform. The transition elements of Fe, Mn and Cr can exchange with each other [14] so that they can exist synchronously in Fe-bearing intermetallic compounds, such as ␣-Al(Fe,Mn,Cr)Si, ␦-Al4 (Fe,Mn,Cr)Si2 phases. Fig. 5 shows the TEM bright field image of particulate ␣-Al(Fe,Mn,Cr)Si phase and SAD patterns with [0 2 3] and [0 1 1] diffraction direction. It indicates that partial ␣-Al(Fe,Mn,Cr)Si phase is less than 1 ␮m that is consistent with Fig. 1(d), and the SAD patterns show that the ␣-Al(Fe,Mn,Cr)Si phase is cubic structure, which is similar to ␣-Al(Fe,Mn)Si phase with BCC structure [60]. Meanwhile, coarse needle-like ␦-Al4 FeSi2 phase in TC-3C alloy can be replaced by skeletal Fe-bearing phase (the stoichiometric: Al14.05 (Fe,Mn,Cr)3.26 Si2.71 , approximated to ␣-Al15 (Fe,Mn,Cr)3 Si2 deduced by EPMA (Table 2)) except for much less ␤-Al5 FeSi phase detected by XRD (Fig. 4(b)) in TC-MC21 alloy. The sizes and morphologies of primary Si and needle-like Febearing phases in 3C alloy have been improved extensively using SF process, and the needle-like Fe-bearing phase has been further refined into granular ␣ Fe-bearing phase (<5–6 ␮m) through (2Mn + 1Cr) addition (Table 4). Large number of spherical ␣Al(Fe,Mn,Cr)Si phase particles with uniform distribution in ␣-Al matrix can hinder the sliding of dislocation in the interface between

Table 4 Size and morphology of different phases in both two alloys with different processing type. PT

Al–25Si–5Fe–3Cu Phase

Morphology

Size/␮m

Phase

Morphology

Size/␮m

Cast

␤-Si ␪-Al2 Cu ␦-Al4 FeSi2 a ␤-Al5 FeSi

Plate/block Netlike Needle-plate –

>100 <20 100–800 –

␤-Si ␪-Al2 Cu ␤-Al5 (Fe,Mn,Cr)Si ␣-Al15 (Fe,Mn,Cr)3 Si2 a

Petaliform Netlike Platelike Skeletal

<100 <5 <50 <100

SF

␤-Si ␪-Al2 Cu ␤-Al5 FeSia ␦-Al4 FeSi2

Granular Granular Shot-rod Shot-rod

<10 <5 <10 <10

␤-Si ␪-Al2 Cu ␤-Al5 (Fe,Mn,Cr)Si ␣-Al15 (Fe,Mn,Cr)3 Si2 a

Blocky Granular Short plate Particulate

<10 <5 10–20 <5

␤-Si ␪-Al2 Cu ␦-Al4 FeSi2 a ␤-A l5 FeSi

Block Netlike Acicular –

2–20 <15 <50 –

␤-Si ␪-Al2 Cu ␤-Al5 (Fe,Mn,Cr)Si/or ␦-Al4 (Fe,Mn,Cr)Si2 ␣-Al15 (Fe,Mn,Cr)3 Si2 a

Granular Granular Needle-like Needle-like Particulate

10–20 <5 <40 <40 <5

OP

Al–25Si–5Fe–3Cu–2Mn–1Cr

PT = processing type, SF = Spray Forming, OP = over-sprayed powders. a It means this phase is the main Fe-bearing phase in the alloy with different processing type.

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Fig. 5. TEM bright field image of particulate ␣-Fe phase and selected area diffraction pattern of this phase with [0 2 3] and [0 1 1] directions.

␣-Al(Fe,Mn,Cr)Si phase particles and ␣-Al matrix, which could be propitious to the improvement of strength and uniform deformation of present alloy, avoiding fast fracture and failure originated from breakable coarse plate Fe-bearing phase or the interface between ␣-Al matrix and Fe-bearing phase. Meanwhile, the phase/microstructure formation of present SF3C and SF-MC21 alloys can be explained via analyze the DSC cooling process, combining the SF process. During the solidification process of HEAS alloys with low cooling rate (traditional cast), ␦-Al4 FeSi2 phase can precipitate after primary Si phase when the Si content is up to 25 wt.% [61], and grow at the consumption of Fe and Si atoms supersaturated in ␣-Al matrix along with temperature decreasing. The starting precipitation temperature of primary Si and ␦-Al4 FeSi2 phases are 1076 and 997 K (exothermic peak and in Fig. 6(a)), respectively. Once the nucleation of ␦-Al4 FeSi2 phase can be occurred, both Fe and Si elements would be transferred to the edge of ␦-Al4 FeSi2 phase because of the decreasing solubility of these two elements in ␣-Al matrix [14], promoting further growth of ␦-Al4 FeSi2 phase with one fast growth direction and two slow directions, which contributes the needle-like morphology. Subsequently, ␦-Al4 FeSi2 phase will be transformed into ␤-Al5 FeSi phase by a peritectic reaction [61] which can be accomplished via the diffusion of Fe and Si atoms from ␦-Al4 FeSi2 phase to ␣-Al matrix. However, low diffusion coefficients of Fe and Si atoms in ␣-Al matrix (DFe = 7.923 × 10−14 m2 /s and DSi = 1.524 × 10−12 m2 /s at 883 K, calculated according to Ref. [62]) and incessant temperature reduction during solidification process both limit the diffusion of Fe and Si atoms, especially, the plentiful diffusion of Si atom from ␦-Al4 FeSi2 phase with high concentration to ␣-Al matrix. As a result, ␦-Al4 FeSi2 → ␤Al5 FeSi transformation cannot be carried out so that the ␦-Al4 FeSi2 phase becomes the dominant Fe-bearing phase in TC-3C alloy with scarcely any ␤-Al5 FeSi phase, and the corresponding exothermic peak has not been detected at present DSC experiment. However, an exothermic peak among 923–875 K (Fig. 6(b)) is characterized during the cooling process of MC21 alloy. Fig. 2(b) shows ␣-Al(Fe,Mn,Cr)Si phase is the predominant phase except Al matrix and primary silicon phase and it can be concluded that this peak (among 923–875 K) is related to the formation of ␣-Al(Fe,Mn,Cr)Si phase. But a small exothermic peak among 1009–995 K in Fig. 6(b) is noted and Huang et al. [63] considered it is about the precipitation of ␦-Al4 (Fe,Mn,Cr)Si2 phase, which can transform into ␣-Al(Fe,Mn,Cr)Si phase later. If this is true, ␣-Al(Fe,Mn,Cr)Si phase also can be transformed into ␦-Al4 (Fe,Mn,Cr)Si2 phase during heating process, but the endothermic peak about this transformation is not performed in Fig. 6(b). Therefore, a further investigation should be focused on the formation mechanism of ␣-Al(Fe,Mn,Cr)Si phase. Also, the ternary and quaternary eutectic microstructures can be formed during cooling

of 3C (exothermic peaks and in Fig. 6(a)) and MC21 (exothermic peaks of (837–809 K) and (789–773 K) in Fig. 6(b)) alloys. During the atomization process, the atomized droplets undergo higher cooling rates than the average cooling rate during SF process (103 –104 K/s [48]), promoting partial atomized droplets to solidify completely after the precipitation of primary Si or ␦-Al4 FeSi2 phase and depressing the formation of ␤-Al5 FeSi phase from residual liquid or ␦ → ␤ transformation. But the cooling rate within the as-deposited body will be greatly decreased for the mush layer presented on the surface of as-deposited preform [47,64] (The calculated and experimental results both show the cooling rate might be less than 5 K/s [65,66].) and recalescence will be occurred within mush layer so that the liquid and latent heat on the surface of asdeposited preform will be increased, contributing to the formation

Fig. 6. DSC curves of SF-3C (a) and SF-MC21 (b) alloys.

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Fig. 7. Microstructure (a) and XRD spectrum (b) of spray-formed Al–25Si–5Fe–3Cu alloy solution-treated at 888 K for 2 h, (c) EDS spectrum of the grey block phase in (a).

of ␤-Al5 FeSi phase via eutectic reactions or transformation from metastable ␦-Al4 FeSi2 phase. Therefore, the short-rod ␤-Al5 FeSi phase (<10 ␮m) will dominate the Fe-bearing phases in SF-3C alloy. In the same way, the appearance of mushy layer, in which the original microstructures will undergo a process approximated to an elevated temperature annealing [65], will promote the spherulization of the irregular, fine ␣-Al(Fe,Mn,Cr)Si particles [67] in SF-MC21 alloy. If Huang et al.’s results [63] is true, the dominant phase in over-sprayed powders should be ␦-Al4 (Fe,Mn,Cr)Si2 phase, not ␣Al(Fe,Mn,Cr)Si phase, because of the much high cooling rate during atomization. But the fact is that lots of ␣-Al(Fe,Mn,Cr)Si particles are present in over-sprayed powders (Figs. 1(e) and 2(d)). So it is considered that ␣-Al(Fe,Mn,Cr)Si phase can precipitate likely from liquids and further studies should be attached importantly.

3.2. Thermal stability of present HEAS alloys For application in engine pistons, cylinder liners and other parts, which might endure the combinatorial effects of high temperature and repeating mechanical action, high thermal stability of present HEAS alloys should be desired so as to avoid the dimensional change of the parts that can beget the mechanical failure and influence the working order. It is known that if the coarsening and structure transformation cannot be appeared during a heat treatment, the alloy could be thermal stable below the heat treatment temperature. Therefore, DSC experiments and different heat treatments have been carried out to investigate thermal stability as well as corresponding phase transformation of present HEAS alloys.

Fig. 8. Microstructures of SF-3C alloy heat-treated for 5 h at 700 K (a), 750 K (b), 803 K (c), respectively, (d) is the high magnification of (c).

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According to ternary Al–Fe–Si and quaternary Al–Cu–Fe–Si [14,68] systems, the endothermic peaks of , and in Fig. 6(a) during heating SF-3C alloy are related to the melting of quaternary eutectic, ternary eutectic and primary silicon phase via the following reactions [14]:

About the endothermic peak (800–815 K) in Fig. 6(a), our research [69] shows it could be related to the formation and transformation of Al7 Cu2 Fe phase via reactions: L → Al + Al7 Cu2 Fe L + Al7 Cu2 Fe → Al + ␤-Al5 FeSi + Al2 Cu At the ending of peak , ␤-Al5 FeSi phase would be the main Fe-bearing phase. But after being heat-treated at 888 K for 2 h, the main Fe-bearing phase becomes blocky phase (Fig. 7(a)). The XRD (Fig. 7(b)) and EDS (Fig. 7(c)) results show the blocky phase is ␦Al4 FeSi2 phase. As a result, the peak should be related to the formation of ␦-Al4 FeSi2 phase via the reaction [51]: L + ␤-Al5 FeSi → ␦-Al4 FeSi2 + Si. With increasing temperature, ␦-Al4 FeSi2 phase becomes the only Fe-bearing phase before its starting melting temperature, 982 K (endothermic peak ). Based on above DSC results, SF-3C alloy has been heat-treated for 5 h at 700, 750 and 803 K, respectively. The SEM images are shown in Fig. 8. Compared to SF-3C alloy in Fig. 1(a), the sizes of ␤-Al5 FeSi phase (grey block phase in Fig. 8, composition (at.%): Al 67.86, Si 18.02, Fe 14.11) increase to >20 ␮m with temperatures increasing (Fig. 8). Fig. 9 shows the diffraction peaks of ␤-Al5 FeSi phases are strengthened with temperature increasing while a few Al7 Cu2 Fe phase appears at 750 K. But after heat-treated at 803K for 5 h, the coarsened ␤-Al5 FeSi phase has almost been surrounded by white phase, Al7 Cu2 Fe, containing 71.96 Al, 8.34 Fe and 15.20 Cu (at.%) (Fig. 8(c) and (d)). Also the results (not shown here) of

Fig. 9. XRD spectrums of SF-3C alloy heat-treated for 5 h at 700 K (a), 750 K (b), and 803 K (c).

isothermal-treatment experiments at 803 K show the size and volume fraction of Al7 Cu2 Fe phase increase with times increasing. It can be considered that the ␤-Al5 FeSi phase will be coarsened at low temperature (e.g., 700 and 750 K) with times increasing, but new Al7 Cu2 Fe phase could be formed easily at high temperature besides the coarsening of ␤-Al5 FeSi phase, which could be the beneficial heterogeneous nucleation site for Al7 Cu2 Fe phase. Meanwhile, isothermal treatments (e.g., at 750, 803 K) can promote the formation of Al7 Cu2 Fe phase because it can coexist with Al in equilibrium state in ternary Al–Cu–Fe system and tend to form with high Cu and Fe contents [14] as well as the sufficient diffusion of Fe and Cu atoms advancing the formation of Al7 Cu2 Fe phase. Compared to SF-3C alloy, besides the melting endothermic peaks of ternary eutectic microstructures (Al + Si + ␣Al(Fe,Mn,Cr)Si) and primary silicon phase in SF-MC21 alloy and in Fig. 6(b), which are similar to the peaks (peaks and in Fig. 6(a)), the only peak in Fig. 6(b) is different to peak in Fig. 6(a). But previous results show ␣-Al(Fe,Mn,Cr)Si phase is the leading Fe-bearing phase besides very tiny amount of

Fig. 10. SEM images of SF-MC21 alloy with different heat treatment, (a) 753 K × 10 h, (b) 783 K × 10 h, (c) 813 K × 10 h, and (d) 813 K × 30 h.

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␤-Al5 (Fe,Mn,Cr)Si phase in SF-MC21 alloy. As a result, the peak in Fig. 6(b) should be related to the melting of ␣-Al(Fe,Mn,Cr)Si phase. And it shows ␣-Al(Fe,Mn,Cr)Si phase could be much stable for its high starting melting temperature (943 K). Also it is found that the melting peak of quaternary eutectic disappears for (Mn + Cr) addition so as to increase the solidus temperature of MC21 alloy. Whereafter the SF-MC21 alloy has been heat-treated for 10 h at 753, 783, and 813 K, respectively. The corresponding microstructures (Fig. 10(a)–(c)) show there are scarcely any changes between heat-treated and untreated SF-MC21 alloy, for instance, the sizes of granular ␣-Al(Fe,Mn,Cr)Si phase are still less than 5–6 ␮m and less plate ␤-Al5 (Fe,Mn,Cr)Si phase (composition (at.%): Al 67.84, Si 18.72, Fe 12.35, Mn 1.06, Cr 0.03) also is present with no size changes. Prolonging treatment times to 30 h at 813 K dose not induce size changes of granular ␣-Al(Fe,Mn,Cr)Si phase (Fig. 10(d)). It shows that the granular ␣-Al(Fe,Mn,Cr)Si phase has excellent thermodynamically stability without any coarsening or transformation. It is concluded that SF-MC21 alloy possesses better thermal stability than SF-3C alloy and would satisfy the microstructure request for the alloys used in pistons, cylinder liners and other parts. 4. Conclusion (1) Both primary Si and Fe-bearing phases have been refined (<10 ␮m) in SF-3C alloy using SF technique and (2Mn + 1Cr) combined addition induces the formation of uniformly distributed, fine ␣-Al(Fe,Mn,Cr)Si phase particles (∼5–6 ␮m), as the dominant Fe-bearing phase in SF-MC21 alloy, contributing to the improvement of properties. Also, the needle-like ␦-Al4 FeSi2 phase in TC-3C alloy can be substituted by skeletal ␣-Al(Fe,Mn,Cr)Si phase with (2Mn + 1Cr) addition. (2) The SF-3C alloy possesses low thermal stability for the coarsening of ␤-Al5 FeSi phase and formation of Al7 Cu2 Fe phase during heat treatments and it can be considered that SF-3C alloy cannot be applied in heat-resistant condition. Contrarily, SF-MC21 alloy has excellent thermal stability because of the presence of fine ␣-Al(Fe,Mn,Cr)Si phase particles, which are thermodynamically stable without any coarsening or transformation during the heat treatment, even at high temperature (813 K) for a long time, contributing to improve the heat-resistance. (3) Besides the quaternary and ternary eutectic reactions and the melting of primary Si phase, a new reaction about Al7 Cu2 Fe phase and ␤-Al5 FeSi → ␦-Al4 FeSi2 peritectic transformation occur at 803 and 880 K, respectively, in SF-3C alloy during the heating process. No endothermic peaks appeared below 800 K after (2Mn + 1Cr) addition. During solidification process, ␦ → ␤ peritectic transformation cannot be found and ␦-Al4 FeSi2 phase becomes the main Fe-bearing phase except for tiny ␤-Al5 FeSi phase generated by ternary eutectic reaction in TC-3C alloy. In addition, a further investigation should focus on the formation of ␣-Al(Fe,Mn,Cr)Si phase with high melting temperature (943 K). Acknowledgment The authors would like to acknowledge the financial support of the State Key Development Program for Basic Research of China (Grant No. 2006CB605204). References [1] K. Ito, H. Kobayashi, Adv. Eng. Mater. 8 (9) (2006) 828–835. [2] W.S. Miller, L. Zhuang, J. Bottema, et al., Mater. Sci. Eng. A280 (2000) 37–49.

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