稀有金属(英文版) 2020,39(10),1210-1221

Microstructure and properties evolution of Al-17Si-2Fe alloys with addition of quasicrystal Al-Mn-Ti master alloy

Jing Zhang Zhou-Yi Pang Ling Wang Chi-Chi Sun Ning Liu Hong-Mei Chen Zhao-Xia Cheng Ke Li

School of Metallurgy and Materials Engineering,Jiangsu University of Science and Technology

Zhangjiagang Industrial Technology Research Institute,Jiangsu University of Science and Technology

School of Materials Science and Engineering,Yingkou Institute of Technology

School of Materials Science and Engineering,Jiangsu University of Science and Technology

作者简介：*Jing Zhang,e-mail:sduzhangjing@126.com;

收稿日期：12 November 2019

基金：financially supported by the National Natural Science Foundation of China (No.51201071);the National Natural Science Foundation of Jiangsu Provence(BK20161270);Jiangsu Overseas Visiting Scholar Program for University Prominent Young&Middle-aged Teachers and Presidents(2018);

Microstructure and properties evolution of Al-17Si-2Fe alloys with addition of quasicrystal Al-Mn-Ti master alloy

Jing Zhang Zhou-Yi Pang Ling Wang Chi-Chi Sun Ning Liu Hong-Mei Chen Zhao-Xia Cheng Ke Li

School of Metallurgy and Materials Engineering,Jiangsu University of Science and Technology

Zhangjiagang Industrial Technology Research Institute,Jiangsu University of Science and Technology

School of Materials Science and Engineering,Yingkou Institute of Technology

School of Materials Science and Engineering,Jiangsu University of Science and Technology

Abstract：

The Fe-rich intermetallic compounds in Al-17Si-2 Fe were modified via Al-Mn-Ti quasicrystal master alloy.The effect of master alloy content on the Fe-rich phase morphology was studied by scanning electron microscope(SEM) and thermodynamic calculation.Results show that the microstructure of the Al-Mn-Ti master alloy consists of binary quasicrystal matrix and ternary AlMnTi secondary phase.The evolutive tendency of Fe-rich intermetallic compounds with content of quasicrystal Al-Mn-Ti master alloy increasing can be described as follows:long needle-shaped β phase for Al-17 Si-2 Fe alloy,long plate-shaped ternary δ phase for 3 wt% master alloy addition,Chinese-script and polyhedral α phases for 4 wt% master alloy addition and finer plate-shaped quaternary δ phase with a phases for 5 wt% master alloy addition.The ultimate tensile strength of the Al-17 Si-2 Fe alloy with4 wt% master alloy addition(a mass ratio of w_Mn/w_Fe ≈0.7) increases by 23.8% and the friction coefficient decreases from 0.45 to 0.35 compared with those of Mnfree alloy.α-Fe phases have less negative effect on the matrix compared with the long needle-shaped β phase and the plate-shaped δ phase.

Keyword：

Fe-rich intermetallic compounds; Al-Mn-Ti quasicrystal alloy; Microstructure; Mechanical properties; Friction coefficient;

Received： 12 November 2019

1 Introduction

Hypereutectic Al-Si alloys are attractive materials for many commercial applications due to their low density,low thermal expansion coefficient and high wear resistance [ 1, 2, 3] .They are ideal piston materials.With modern requirement of increasing engine speed and power,the mechanical properties are expected to be improved.Adding higher Fe amount to hypereutectic Al-Si alloys can improve the thermal stability of the alloys.However,iron is always treated as impurity in Al-Si alloys because Fe can combine with Al and Si to form coarse needle-shapedβ-Al₅FeSi phase,which is detrimental to the mechanical properties due to its platelet morphology,leading to a stress concentration and crack initiation [ 4, 5] .Besides,the needle compounds tend to impede the fluidity of molten metals which is easy to lead to casting defects such as gas holes and shrinkage cavities [ 6] .Therefore,the modification of needle-shapedβ-Al₅FeSi phase is the key to improving the properties of hypereutectic Al-Si-Fe alloys.

It has been reported that the morphology of Fe-rich intermetallic compounds can be modified by rapid solidification process,ultrasonic vibration and elements addition.Spray processing and hot extrusion were used to prepare an Al-18%Si-5%Fe-1.5%Cu alloy,in which onlyδ-Al₄FeSi₂phase was found in atomized powder particles [ 7] .After ultrasonic vibration treatment,the Fe-containing compounds in the Al-17Si-xFe alloys were modified toδ-Al₄FeSi₂ particles with average grain size ranging from 26to 37μm,and only a small amount ofβ-Al₅FeSi phases remained [ 8] .Elements such as Mn,Sc,Be,Sr have been already proved to change the Fe-rich phase from needleshaped to nodular or Chinese script form multi-component compounds [ 9, 10, 11] .Mn is the most effective and widely studied element among the above elements modification method.Mn has been used to replace the needle-shapedβ-phase withα-Al(Mn,Fe)Si which has granular or Chinese script morphology [ 12, 13] .The fundamental mechanism is that Mn atoms disperse to displace part of Fe atoms in the form of ternary compounds.In other words,the efficiency depends on the diffusion of free Mn atoms.Mn is introduced as master alloy traditionally in which the state of Mn is compound with high melting point (usually above 900℃).Therefore,the modification efficiency can be improved if the Mn diffusion rate increases.As known,the Al-Mn alloys are a good option to obtain quasicrystal structure in a relatively high cooling velocity [ 14, 15] .Furthermore,quasicrystal is a metastable state which means that the diffusion energy barrier of Mn atoms is much smaller than that in normal master alloy.Furthermore,Al-Mn-Ti quasicrystal master alloy has been proved to result in positive effect on the primary Si.Primary silicon changed from coarse five-petal asteroid shape to small particles,and the average size of primary silicon decreased from 300 to 30μm when quasicrystal master alloy addition was 4 wt% [ 16] .

Accordingly,Al-Mn-Ti quasicrystal master alloy was prepared to modify the morphology of Fe-rich phase in Al-17Si-2Fe alloy in present work.The refinement effect of the master alloy on primary Si was also identified.The effect of master alloy on the microstructure,phase evolution,mechanical properties and friction coefficient was studied.

2 Experimental

The Al-Mn-Ti quasicrystal master alloys were prepared with commercial pure Al (99.99%),pure Mn (99.99%) and pure Ti (99.99%).The chemical compositions of the AlMn-Ti master alloy are 36 wt%Mn,7 wt%Ti,and balance Al.The alloys were melted in a graphite cup (wall thickness of 6 mm) using a high-frequency furnace.The power of the high-frequency induction furnace was gradually increased from 0 to 5 kW in 15 min,and then the melt was held for 20 min for homogeneous dissolution before being poured into a rapid solidification copper mold (preheated to200℃),as shown in Fig.1.The master alloy at Position A was chosen to be used in the following processing.

Fig.1 A cross section of copper mold for pouring Al-Mn-Ti master alloys (mm)

The base alloys in present study are ternary Al-17Si-2Fe alloys.0%,3 wt%,4 wt%and 5 wt%Al-Mn-Ti quasicrystal master alloys were introduced to Al-17Si-2Fe alloys as shown in Table 1.The base alloys were first melted in a resistance furnace;then master alloys were added at 800℃for holding 15 min.The liquid was then poured into a steel mold (preheated to 250℃,as shown in Fig.2) at 770℃.

To characterize the microstructure and determine the constitution of phases,samples were obtained from the central ingots by wire-electrode transverse cutting and then mechanically ground with several grades of sandpaper and polished.Microstructural observations were carried out using a Zeiss Axio Scope Al optical microscope (OM) and a JSM-6510LA scanning electron microscope (SEM)operated at 15 kV,equipped with an energy-dispersive spectrometer (EDS).The samples for OM and SEM were etched in a solution of hydrochloric acid for 15-30 s.

A Thermo-calc software package was used to investigate the chemical reaction and mechanism of phase transition in combination with the phase diagrams of the Al-17Si-2Fe-yMn (y=1.08/1.44/1.8,corresponding to 3wt%/4 wt%/5 wt%master alloys addition,respectively)alloys.The phase diagrams were calculated using the latest Al database for the temperature range of 300-800℃.

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Table 1 Chemical compositions of investigated Al-17Si-2Fe alloys(wt%)

Fig.2 Diagram of casting composite (mm)

The tensile samples were manufactured according to the GB/T228.2-2015 (ISO6892-2:2011) standard.The tensile tests were performed on a Model 5582 Universal Tester with an extension rate of 1.0 mm.min^-1 at 20℃.At least three tensile samples were tested for each chemical composition.

The wear tests were performed at room temperature using a pin-on-disk type sliding wear apparatus (HT-1000).The Al-Si-Fe alloys were processed to the sample with a diameter of 5 mm and a length of 12 mm.304 stainless steel disc with a hardness of HV 450 and a diameter of 60mm was used as the counterpart during the wear test.The test load was 15 N,and the friction radius was 20 mm.The plate rotation speed was 300 r·min^-1 with a test time of 16min and sliding distance of 600 m for each sample.

3 Results and discussion

3.1 Microstructure of master alloy

Figure 3 shows the microstructure of Al-Mn-Ti master alloy.Star-shaped secondary phases with round branches edge can be observed distributing uniformly on the matrix(Fig.3a,b).The fracture of matrix (Fig.3c) shows decagonal prism which corresponds to the typical characteristic of quasicrystal alloy.Furthermore,the cracks can be easily found along the decagonal prism.It means that the matrix is quite brittle,which is another great feature of quasicrystal alloy.EDS results show that the matrix is AlMn binary compounds while the secondary phase has Al,Mn,Ti three elements.As known,Al-Mn alloy has high tendency to form decagonal quasicrystal structure under high cooling rate.The cooling rate to form Al-Mn quasicrystal phases is between the range of 1×10⁴ and1×10⁷ K·s^-1 [ 14] .Other research results proved that Ti is useful to make the Al-Mn quasicrystal structure more stable [ 15] .The present cooling rate of the casting at the position A (Fig.1) was 3.15×10⁵ K·s^-1.Therefore,decagonal prism matrix should be AIMn binary quasicrystal phase.Accordingly,the microstructure of the AlMn-Ti master alloy consists of binary quasicrystal matrix and ternary AlMnTi secondary phase.

3.2 Microstructural evolution of Fe-rich intermetallic compounds

Figure 4 shows typical microstructure of as-cast Al-17Si-2Fe alloy,which is mainly composed of primary Si,eutectic (Al+Si) and coarse needle-shaped Fe-rich intermetallic compounds.It is worth noting that some of these thin needle-shaped phases with only 1-μm width bend and some of them have branches.According to EDS results,these intermetallic phases are supposed to beβ-Al₅FeSi.

With 3 wt%quasicrystal Al-Mn-Ti master alloy addition,different morphologies of Fe-rich compounds are formed,as shown in Fig.5.Coarse plate-shaped phases present a length of 150μm and a width of 40μm.EDS results reveal that the compounds are depleted in Mn,which is supposed to beδ-Al₄FeS_hphase.There are higher Si content and lower Al content in theδthan in theβphase shown in Fig.4.Besides,the needle-like phases are still likely to form.However,the size and amount are remarkably reduced.

Figure 6 shows the microstructure of as-cast Al-17Si-2Fe alloy with 4 wt%master alloy.Needle-shapedβFerich phases convert completely into two different morphologies:finer Chinese-script and coarser polyhedral compounds.EDS results show that the later has a higher Mn content.These Fe-rich compounds can be classified asα-Al₁₅(Fe,Mn)3Si₂ phases.

As the master alloy gradually increases to 5 wt%,plateshapedδFe-rich phase presents again accompanied by twoαFe-rich phases,as shown in Fig.7.However,the size seems much smaller in Alloy 3,compared with theδFerich phase in Alloy 1.The length reduces from 150 to 100μm.Furthermore,the former has a higher Mn content.Besides,fine needle-shape Ti-rich compounds can be found in the top left corner in Fig.7a.Figure 8 gives the morphology and elemental composition of this phase.These compounds with average length of 25μm are supposed to be AISiTi ternary phases.In contrast,the Ti distributes uniformly on the matrix for 3 wt%and 4 wt%master alloy additions.No obvious compounds can be detected according EDS mapping results (Figs.5b,6b).

Combining the above microstructure (from Figs.4,5,6b,7,8),the evolutive tendency of Fe-rich intermetallic compounds with content of quasicrystal Al-Mn-Ti master alloy increasing can be described as follows:long needleshapedβphase→long plate-shaped ternaryδphase→Chinese-script and polyhedralαphases finer plate-shaped quaternaryδphase.

Fig.3 Microstructure and EDS analysis:a OM observation,b SEM observation result,c decagonal matrix,d cracks along decagonal matrix,and corresponding EDS analysis of e grey matrix and f bright secondary phase in b

The intermetallic phases differ in morphology,structure and composition.The needle-shapedβFe-rich phase has been studied widely.Bacaicoa et al.'s study on microcomputed tomography shows that [ 17, 18] theβ-Al₅FeSi phases in the 2D micrographs are seen as long needles present a platelet morphology in the 3D reconstructions.Although there are many crystallographic structures ofβphase reported in the literature including a tetragonal or B-centered orthorhombic,the most wildly accepted theory is a monoclinic structure with a=0.6161 nm,b=0.6157nm,c=2.0813 nm,andβ=90.42° [ 19] .

Since the morphology ofβFe-rich phase is detrimental to mechanical properties,modification to proper forms is the key to resolving this issue.Mn,Cr and Co are usually added to reduce the harmful effects of theβphase by replacing it with the less-detrimentalαphase.Besides,αphase has several kinds of morphologies such as Chinesescrip,polyhedral,bulk and star-like shape,depending on the cooling rate or the alloy element distribution.In present study,two typical Chinese-scrip and polyhedral oaAl₁₅(Fe,Mn)₃Si₂ phases appear after 4 wt%master alloy addition.3D morphology of the later can be seen in Fig.9a,b,showing the branch lateral growth pattern.Theαphases have been identified as cubic with a lattice parameter a=1.25 nm [ 20, 21] or a hexagonal structure with lattice parameters around a=1.240 nm,c=2.623 nm [ 22, 23] .

Fig.4 Microstructure of as-cast Alloy 0 (Al-17Si-2Fe):a SEM image,b EDS elemental distribution maps,and c EDS elemental composition result of intermetallic compound marked as A in a

δphase is an intermediate transition phase or metastable phase during solidification.The typical 3D morphology is plate with or without lateral growth defects,as shown in Fig.9c,d.Twoδphases with different elemental compositions can be found in present study:ternary and quaternary,corresponding to lower and higher Mn content,respectively.It has been reported [ 24] that theδphase has a tetragonal PdGa5-type structure with lattice parameters of a=0.614 nm and c=0.957 nm.Day and Hellawell [ 25] re-determined the symmetry and proposed an orthorhombic structure with a pseudo-tetragonal cell(a=0.6061 nm,b=0.6061 nm,c=0.9525 nm and space group Pbcn).Considering the length of two plate-shapedδphases,it can be deduced that more Mn solid-solute into the compounds can reduce the c value,resulting in finer structure.

Besides,the master alloy has a slight refinement on the primary Si.The size of primary Si decreases from 23μm

for Al-17Si-2Fe alloy to 16,15 and 15 with 3 wt%,4 wt%and 5 wt%master alloy additions,respectively,while in Yang's results [ 16] ,primary silicon changed from coarse five-petal asteroid shape to small particles,and the average size of primary silicon decreased from 300 to 30μm when quasicrystal master alloy addition is 4 wt%.This is mainly because the original primary Si is already in relatively small bulk-shaped form in present Al-17Si-2Fe alloy.No obvious refinement could be obtained as Yang's research.

3.3 Thermodynamic assessment

The varieties of Fe-rich intermetallic phases are explained in the above section.However,the precipitation thermodynamics needs further discussion.Figure 10 shows the calculated phase diagrams of Al-17Si-2Fe alloys with different Mn contents.

According to Fig.10a,the solidification sequence of various phases in Al-17Si-2Fe alloy without Mn addition is as follows:

The first precipitate phase is Si followed by theδphases.Following is the peritectic reaction at approximately600℃,δphase converting toβphase.The precipitation temperature range forδphase is relatively narrow.Therefore,the thermo-stable intermetallic compound isβphase for Al-17Si-2Fe alloy at room temperature which is in agreement with the micro structure in Fig.4.

Fig.5 Microstructure of as-cast Alloy 1 (Al-17Si-2Fe with 3 wt%master alloy):a SEM image,b EDS elemental distribution maps,and c EDS elemental composition result of intermetallic compound marked as A in a

Table 2 reveals the precipitate sequence of different phases in Al-17Si-2Fe alloys with different Mn contents.The final Fe-rich intermetallic phases areα-Fe+β-Fe for Alloy 1 andα-Fe for Alloy 2/3,respectively.It does not comply with the observed solidification micros truc ture shown in Figs.5,7,except for the 4 wt%master alloy shown in Fig.6.This discrepancy can be attributed to the kinetic effects and alloying elements influence during solidification,which are not considered by the thermodynamic calculations.

For Alloy 1 with 3 wt%master alloy,the plate-shapedδphases take place of part needle-shapedβphase.However,Mn content is too low to be detected by EDS.It means that Mn addition has stronger effect on the solidification sequence rather than the phase structure evolution induced by Mn solid solution.The precipitated temperature of 8phases increases due to Mn addition,which widens the existing temperature range to keep some metastableδphases through solidification processing,as shown in Fig.5.The mechanism is similar to the ultrasonic vibration treatment on Al-17Si-2Fe alloys studied by Lin et al. [ 8, 26] .Acoustic streaming and cavitation of USV homogenize the solute field and temperature field and increase the start-freezing temperature ofδ-Al₄FeSi₂phase,thereby promoting the formation of fineδ-Al₄FeSi₂particles.

The microstructure of Al-17Si-2Fe with 4 wt%master alloy is consistent with the thermal-cal result,finalα-Fe intermetallic compounds.α-Fe phase is supposed to be the thermodynamic stable phase in terms of Fig.l0b.Narayanan et al. [ 27] considered that Mn addition broadened theα-Fe precipitated temperature range.Kim et al. [ 28] put forward the solidified temperature rained as Mn content increased.Accordingly,needle-shapedβFe-rich phases convert completely into two finer Chinese-script and coarser polyhedral oa-Fe compounds.

As to Alloy 3 with 5 wt%master alloy,finer quaternary plate-shapedδcompounds appear along with ot phases.Lin et al. [ 8, 26] point that the peritectic reaction is that the primary phase reacts with liquid directly and peritectic phase is then formed (L+δ→α) in Al-17Si-2Fe from experimental aspect.Besides,the influence of Ti cannot be ignored in Alloy 3.The content of Ti is high enough to form AISiTi compound,so does to affect the solidification dynamics of Fe-rich phase.Ti can lead to constitutional supercooling which makes local higher cooling rate.An agreement has been reached by many researchers that the cooling rate plays a significant role in the Fe-rich intermetallic phase formation.In Becker et aFs study [ 29] ,higher cooling rate can result in solidification path changing fromαphase toα+δin the same composition.Therefore,higher Ti addition promotes certain portion ofδFe-rich phases precipitating in Alloy 3.

Fig.6 Microstructure of as-cast Alloy 2 (Al-17Si-2Fe with 4 wt%master alloy):a SEM image,b EDS elemental distribution maps;c and d EDS elemental composition result of the intermetallic compound marked as A and B in a,respectively

3.4 Mechanical and wear resistance properties

The ultimate tensile strengths (UTS) of four alloys in the as-cast condition are summarized in Fig.11a.As Mn addition increases,the UTS of Al-17Si-2Fe alloy improves first and then has a slight decrease.Alloy 2,Al-17Si-2Fe alloy with 4 wt%master alloy (corresponding to a mass ration of wMn/wFe≈0.7),has the best general mechanical properties.The UTS increases by 23.8%compared with that of the Mn-free alloy.

Fig.7 Microstructure of as-cast Alloy 3 (Al-17Si-2Fe with 5 wt%master alloy):a SEM image,b EDS elemental distribution maps,and c EDS elemental composition result of intermetallic compound marked as A in a

Fig.8 Microstructure of Ti-rich compounds in as-cast Alloy 3 (Al-17Si-2Fe with 5 wt%master alloy):a SEM image and b EDS elemental composition result

Figure 11b shows friction coefficient of Al-17Si-2Fe alloys with different contents of master alloys.The average friction coefficient of Mn-free alloy is approximately 0.45.After 3 wt%,4 wt%and 5 wt%master alloy additions,the friction coefficient decreases to 0.38,0.35 and 0.38,respectively.Additionally,the profile of Al-17Si-2Fe alloy with 4 wt%master alloy is the most stable as testing time increases,which means that the alloy with a mass ratio of w_Mn/w_Fe≈0.7 has the best wear resistance properties.

Fig.9 SEM images of a,b polyhedral ct phases and c,d plate-shaped quaternaryδFe-rich

Fig.10 Calculated phase diagrams of Al-Fe-Si alloys:a vertical section phase diagram of Al-17Si-xFe alloys with different Fe contents;b vertical section phase diagram of Al-17Si-2Fe alloys with different Mn contents

Since the morphology and size of primary Si have no big difference for the Al-17Si-2Fe with different master alloy additions,the main improvement of properties is attributed to the morphology change of Fe-rich intermetallic compounds.It is known that the long needle-shapedβphase is detrimental to the mechanical properties.It can split the matrix under tension or wear processing which provides for the racks generation and multiplication.Besides,the needle compounds tend to impede the fluidity of molten metals which is easy to lead to casting defects such as gas holes and shrinkage cavities [ 15, 17] .Two types of shrinkage pores were analyzed in the studied alloy,one small with regular shape,and the other one large and elongated with dendritic arms.These issues make Fe usually be treated as impurity element.In contrast,α-Fe compounds have a modified morphology,which have little effect on the formation of surface and sub-surface microcracks [ 30] .Besides,oa-Fe compounds form a rough interface with the matrix and their better bonding with matrix reduces the possibility of crack formation in the interface of intermetallic compounds with the matrix during tension and wear processing.Therefore,the modified Chinese-script and polyhedral oa-Fe compounds have light effect on the matrix,resulting in better mechanical and wear resistance properties in Alloy 2 with 4 wt%master alloy addition.However,plate-shapedδFe-rich phases have a larger length-width ratio than oa-Fe compounds,which is the reason for the slight decrease in UTS and increase in friction coefficient when master alloy addition reaches up to 5 wt%.

With regard to the optimal xMn:xFe,a ratio of 1:2 is suggested when Fe content is below 0.3 wt% [ 31] .Hwang et aFs [ 32] research shows that,with Mn/Fe ratio of 1.2 in the Al-7 wt%Si-3.8 wt%Cu-0.5 wt%Fe alloy,the platelike intermetallic phase is completely converted to the Chinese script phase,resulting in improved tensile property.However,it has also been reported by Seifeddine et al. [ 33] that the plates cannot be completely suppressed even at x_Mn:x_Fe=2:1 in Al-9 wt%Si alloy.It reveals that Si also plays a role in Fe-rich intermetallic compounds evolution.Taylor et al. [ 34] proposes that the content of Ferich phases also rises with Si content increasing.It means that the optimal Mn/Fe ratio for Al-17Si-2Fe is far more than 1:2.Mascre [ 35] proposed a general formula of the type after carrying out metallographic studies on sand and permanent mould casting of the same alloy containing0.55-1.20 wt%Fe and up to 1.30 wt%Mn:

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Table 2 Solidification sequence of various phased in Al-17Si-2Fe alloys with different Mn contents

According to this equation,at least 3 wt%Mn (w_Mn/w_Fe≈1.5) should be introduced to Al-17Si-2Fe to obtain complete theα-Fe microstructure.Anyway,the alloy with x_Mn/x_Fe≈0.7 can be improved to the best mechanical properties in present study.This is can be attributed to the Mn existing state in the master alloy.In traditional Al-Mn master alloy,Al₄Mn compound is the common state for Mn.The melting temperature is approximately 916℃.The casting temperature is usually (750±50)℃for Al-Si alloys.In other words,sufficient holding time is necessary for Mn dissolving into the molten metal from Al₄Mn compound to modify the morphology of Fe-rich phase.The diffusion coefficient is known as:

where D₀ is the diffusion constant,Q is the diffusion active energy,R is the gas constant,and T is the temperature.In present work,the Al-Mn-Ti master alloy is quasicrystal with binary and ternary phases.As known,quasicrystal is a metastable state [ 36, 37] .The energy barrier of Mn atoms,namely the diffusion active energy (Q),is much smaller in quasicrystal than in the normal Al-Mn master alloy under the same holding temperature and time.Therefore,the optimal Mn/Fe ratio is less than the theoretical calculation result for Al-17Si-2Fe.

Fig.11 a UTS and b friction coefficients of as-cast Al-17Si-2Fe alloys with different Mn content

4 Conclusion

The microstructure of the Al-Mn-Ti master alloy prepared in present work using copper mold consists of binary quasicrystal matrix and ternary AIMnTi secondary phase.The evolutive tendency of Fe-rich intermetallic compounds with content of quasicrystal Al-Mn-Ti master alloy increasing from 0 to 5 wt%can be described as follows:long needle-shapedβphase→long plate-shaped ternaryδphase→Chinese-script and polyhedralαphases→finer plate-shaped quaternaryδphase.Higher Ti content for 5wt%master alloy results in large undercooling value which leads to the non-equilibrium solidification,theαphases converting toδphase.Besides,the master alloy has a slight refinement on the primary Si.The size of primary Si decreases from 23μm for Al-17Si-2Fe alloy to 16,15 and15μm for 3 wt%,4 wt%and 5 wt%master alloy additions,respectively.

The alloy adding 4 wt%master alloy (w_Mn/w_Fe≈0.7)with modified Chinese-script and polyhedralα-Fe phases has the best mechanical properties.The UTS increases by23.8%and the friction coefficient decreases from 0.45 to0.35 compared with those of Mn-free alloy.α-Fe phases have light effect on the matrix compared with the long needle-shapedβphase and the plate-shapedδphase.The energy barrier of Mn diffusion is much smaller in quasicrystal than the normal Al-Mn master alloy under same holding temperature and time.Therefore,the optimal Mn/Fe ratio is less than the theoretical calculation result for Al-17Si-2Fe.

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