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Rare Metals2020年第6期

Semisolid microstructural evolution of (CNTs+Si_p)/AZ91D powder compacts prepared from powders by cold pressing and remelting

Pu-Bo Li Wan-Ting Tan Mang-Mang Gao Kuan-Guan Liu

Ningxia Key Laboratory of Photovoltaic Materials,Ningxia University

作者简介：*Pu-Bo Li,e-mail:lipubo@nxu.edu.cn;

收稿日期：8 May 2019

基金：financially supported by the Ningxia Natural Science Foundation of China(No.2018AAC03031);the Basic Scientific Fund of Ningxia university(No.NGY2018009);the Natural Science Foundation of Ningxia University(No. ZR1702);

Semisolid microstructural evolution of (CNTs+Si_p)/AZ91D powder compacts prepared from powders by cold pressing and remelting

Pu-Bo Li Wan-Ting Tan Mang-Mang Gao Kuan-Guan Liu

Ningxia Key Laboratory of Photovoltaic Materials,Ningxia University

Abstract：

The evolution of semisolid microstructure during partial remelting of(CNTs+Si_p)/AZ91 D powder compacts prepared by cold pressing was studied.The results indicate that rapid grain coarsening is driven by the dissolution of eutectic β phase material during the initial heating period of 0-10 min,so the AZ91 D powders with fine equiaxed grains surrounded by intergranular eutectic phases evolve into compact particles.As the heating time proceeds,α-Mg particles were gradually separated by liquid due to the phase transformations of α-Mg+β→L and α-Mg→L.The primary particles coarsened rather slowly and the as-received Mg powder evolved into nearly spheroidal particles surrounded by liquid phase after partial remelting.The in situ synthesized Mg₂ Si_p were distributed homogeneously around the CNTs with maintained structural integrity during partial remelting.Moreover,this microstructural evolution was accompanied by densification through pore filling.An ideal semisolid ingot suitable for thixoforming can therefore be obtained by partially remelting a(CNTs+Si_p)/Mg powder compact.

Keyword：

Semisolid microstructural evolution; Densification; Phase transformation; Coarsening behavior;

Received： 8 May 2019

1 Introduction

Magnesium (Mg) and its alloys are the lightest structural metal materials with higher specific strength and damping capacity than aluminum and iron [ 1, 2] .However,the relatively low mechanical strength and creep resistance of Mg alloys especially at high-temperature limit the application of Mg in the transportation and automobile industries.Mg-based composites,especially with nanomaterial reinforcements,could be an effective and promising way to overcome these intrinsic drawbacks [ 3] .

Mg2Si nanoparticles (Mg2Si_p) have a high melting temperature (1085℃),low thermal expansion coefficient(7.5×10^-6 K^-1),and high elastic modulus (120 GPa),which makes them a desirable candidate for a reinforcement additive in Mg₂Si_p/Al or Mg₂Si_p/Mg matrix composites [ 4] .Moreover,Mg₂Si_p can be in situ synthesized through the reaction between Si and Mg.However,the strengthening efficiency of metal matrix composites(MMCs) reinforced with conventional ceramic particles is low and the material lacks ductility and functionality.Conventional MMCs are therefore ill-suited to the increasingly comprehensive performance requirements [ 5] .Carbon nanotubes (CNTs) have excellent features like lightweight,high strength (～100 GPa),and high Young’s modulus (～1 TPa) [ 6, 7] .These features make them ideal nanofiber reinforcements for structural and functional composite materials.In the past decade,CNTsreinforced MMCs have been studied intensively [ 8, 9] .Zhang et al.reported that the small addition of 0.5 wt%CNTs into the Mg matrix increased the yield strength (YS)and elongation of the composite by 16.4%and 14.2%,respectively,yielding a material with a good combination of strength and ductility [ 9] .However,the improved mechanical properties of the CNTs-reinforced MMCs were not commensurate with the superior strength of the CNTs,partly because of the agglomeration of CNTs in the matrix.Hybrid additives with low content of one-dimensional CNTs and other zero-dimensional nanoparticles will significantly enhance the performance of the matrix alloy compared with single-CNTs reinforcement due to the more-uniform dispersion of CNTs and the associated increase in strengthening efficiency [ 10, 11, 12] .The YS(247 MPa) of Al matrix composites reinforced with a combination of 0.5 vol%CNTs and 0.5 vol%SiC nanoparticles (SiC_p) increased by 14%and 57%over that of 1.0 vol%CNTs/Al (217 MPa) and 1.0 vol%SiC_p/Al(157 MPa) composites,respectively [ 13] .This increase in strength is attributed to the SiC_p that promotes the uniform dispersion of CNTs like a ball-milling medium and inhibits excessive interfacial reactions.Though this literature is extensive,most references are focused on hybrid CNT and nanoparticles reinforced Al matrix composites,while investigations about hybrid-reinforced Mg matrix composites are quite scarce.

Hybrid composites are mainly fabricated with powder metallurgy (PM).PM makes it difficult to fabricate large components with complex shapes.Powder thixoforming technology,proposed in our previous studies,may overcome these drawbacks by combining the ball-milling and cold-compact procedures of PM with the partial remelting and forming processes of thixoforming [ 14] .The microstructural evolution of the material before thixoforming has important effects on the mechanical properties of the final thixoformed composites.Many studies have considered the preparation of nondendritic semisolid ingots,using methods like mechanical stirring and straininduced melt activation [ 15, 16, 17] .Moreover,the evolution processes during partial remelting of alloys with different initial microstructures (including developed dendrites,plastic deformation microstructure,and fine grains) have been studied extensively.Chen et al. [ 16] reported that the semisolid microstructural evolution of the AZ91D Mg alloy refined with SiC particles can be pided into four stages:initial coarsening,microstructural separation,spheroidization,and finally coarsening of the primary particles.The coarsening rate of the extruded AZ80M alloy during partial remelting increased at first and then decreased as the liquid fraction increased,which was confirmed with a modified liquid film migration model [ 15] .The grain size in existing investigations is larger than that of atomized alloy powders (several micrometers),while a clear understanding of the evolution mechanism of these fine grains during partial remelting has not yet been achieved.Our previous work found that an ideal semisolid ingot can be obtained when green compacts prepared from2024Al powders with a particle size of 20μm were partially remelted at 625℃for 30 min and an as-received Al powder evolved into a spherical primary particle because the Al₂O₃ film surrounding the Al powder surface decreased the coarsening rate [ 18] .However,to the best of our knowledge,no investigations have considered the microstructure and mechanical properties of (CNTs+Mg₂Si_p)/Mg hybrid composites synthesized by powder thixoforming.We need to determine whether AZ91D powders can evolve into semisolid ingots suitable for thixoforming through partial remelting processes.In addition,the homogeneity of the semisolid microstructure and the coarsening behavior of primary particles are significantly influenced by the introduction of reinforcements in semisolid slurries,which further affects the mechanical properties of the composites [ 19] .However,no investigations have considered how hybrid additives of CNTs and Mg₂Si_p affect the microstructural evolution of the(CNTs+Mg₂Si_p)/Mg powder compacts.Therefore,this work investigated the semisolid microstructural evolution of the (CNT+Si_p)/Mg powder compact during partial remelting.Particular attention was given to semisolid microstructural evolution,coarsening behavior of the primary particles,interfacial reactions between CNTs-Si_phybrids and the Mg matrix,and microstructure densification.These behaviors allow us to clarify the thixoformed microstructure and better understand the mechanical properties of (CNTs+Mg₂Si_p)/Mg hybrid composites synthesized by powder thixoforming.

2 Experimental

The gas-atomized AZ91D powder shown in Fig.1a has a chemical composition of Mg-9.08Al-0.65Zn-0.23Mn(wt%) and an average particle size of 35μm (Tangshan Weihao Magnesium Powder Co.,Ltd.,China).Reinforcement additives were prepared from CNTs powders with diameter of 30-50 nm and length of 10μm and spherical Si nanoparticles (Si_p) with average size of 50 nm (Chengdu Institute of Organic Chemistry Co.Ltd.,China).These materials were used to fabricate Mg₂Si_p/CNTs hybrid reinforced Mg composites (Fig.1b-d).

To achieve homogeneous dispersion of the CNTs and Si_p in the Mg matrix,two procedures were used.The first was a wet method using ultrasonication and mechanical stirring.CNTs,Si_p,and Mg powders were dispersed separately in ethanol by ultrasonication for 30 min to make the uniformly dispersed suspensions.These suspensions were then blended and stirred magnetically for 1 h.The mixtures were filtered and vacuum-dried for 6 h at 70℃to obtain the (CNTs+Si_p)/Mg composite powders.Second,a dry process using QM-QX2 planetary ball-milling machine was applied.The mixed powders were sealed into a 250-ml stainless jar together with stainless milling balls.Ball milling was applied for 4 h,with the ball-to-powder weight ratio of 5:1,and rotation speed of 200 r·min^-1.The obtained (CNTs+Si_p)/Mg composite powders were consolidated under 200 MPa into a powder compact with dimensions ofΦ10 mm×10 mm.Finally,the powder compact was heated at the semisolid temperature of 585℃for different time (0-60 min) in a vacuum furnace and was then immediately quenched in cold water.To examine the temperature variation of the ingot during partial remelting,a thermocouple was inserted into a hole in the center of the ingot.For the sake of comparison,reference samples made from Mg alloy powders with no additives were also fabricated using the same processes.

Fig.1 SEM images of a AZ91D powders,b,c CNTs and d Si_p

Metallographic specimens were cut from the center of each ingot,and a cross section of one of specimens was finished and polished with standard metallographic techniques.A field-emission scanning electron microscope(SEM,JSM-7500FG) was used to characterize the semisolid microstructures and morphologies of the powders.In addition,microstructural changes in the specimens were further characterized with the energy-dispersive spectrometer (EDS) on the SEM.X-ray diffraction (XRD,SmartLab,Rigaku) using Cu Ka radiation at a scan rate of2 (°)·min^-1 was used to verify the phase constituents of the(CNTs+Si_p)/Mg powder compact heated at 585℃for different durations to infer the phase transformations that took place during partial remelting.The carbonaceous structures were characterized with Raman spectroscopy(DXR0304040404,Thermo Fisher,USA) using 514.5-nm incident laser light.To quantitatively measure the liquid fraction and primary particle size,SEM images were analyzed with Image-Pro Plus software (Media Cybernetics Company,Silver Spring,MD,USA).At least three typical SEM images with magnification of 600 time were examined.The relative density of each sample was determined using Archimedes’principle [ 20] .

3 Results and discussion

3.1 Microstructure of Mg powders

SEM observations were carried out on cross-sectioned asatomized AZ91D powders (Fig.2).The Mg powder consisted of some fine equiaxedα-Mg grains together with a quasi-continuous network of microscale phase along the grain boundaries.These microstructure features were attributed to the metastable microstructure that results from rapid solidification during the atomization process [ 21] .Figure 3 shows XRD results of AZ91D powders,which confirm the existence of p-Mg₁₇Al₁₂ phase along withα-Mg phase.In the final stage of solidification during powder atomization,the eutecticα-Mg phase directly grew on the surfaces of existingα-Mg grains and only the eutecticβphase remained in the intergranular regions,forming a porced eutectic structure.Thus,fineα-Mg grains within the Mg powder were separated by the eutecticβphase.In other words,the as-received microstructure of AZ91D alloy powders is characterized by a composition gradient with local regions oversaturated with alloying elements while other areas had relatively low content of alloying elements.The sizes of theβphase andα-Mg grains were1.5 and 2.0μm,respectively,much smaller than that of the grains in a conventional metal-mold casting.The rapid solidification of the micro structure depends on the size of the atomized powder.Finer powder gives a finer microstructure due to the enhanced undercooling and faster solidification rate.

Fig.2 a SEM image and b enlarged areas of as-atomized AZ91D alloy powders

Fig.3 XRD pattern of AZ91D powders

3.2 Microstructural evolution of (CNTs+Sip)/Mg powder compacts during partial remelting

Figure 4 shows micrographs indicating the microstructural evolution of the (CNTs+Si_p)/Mg powder compacts heated at 585℃for different time.After 5-min heating,the quasi-continuous network of the eutectic phase became a discontinuous distribution along ambiguous grain boundaries (Fig.4a).Simultaneously,the inset image in Fig.4a clearly shows that some white particles appeared within the Mg particles.The amount of the eutectic phases obviously decreased and only a few sparse particles appeared in local regions as the heating time increased(marked by the arrows in Fig.4b).Moreover,the amount of these spherical particles decreased while their size increased.The powders with fine equiaxedα-Mg grains and intergranular eutectic phases then evolved into compact particles with no grain boundaries.In addition,many pores appeared among the Mg powders (marked by the arrows in Fig.4a,b).

According to the binary Mg-Al phase diagram,the AZ91D alloy should experience a singleα-Mg interval during the phase transformation ofβ→αduring heating(considering only Mg and Al) [ 16] .Theβphase dissolved toward the surroundingα-Mg phase,decreasing the amount ofβphase and leading to the coalescence ofα-Mg grains(Fig.4a).The dissolution of the Al-richβphase intoα-Mg phase increased the solubility of Al inα-Mg (Table 1).The heating of the specimens is quite rapid during the initial stage of heating,and the eutectic reaction temperature was exceeded in a short time (Fig.5).The residual eutectics have no enough time to dissolve completely into theα-Mg melted through the reverse eutectic reaction,α+β→L,forming entrapped liquid pools within the powder(Fig.4b).Correspondingly,the spherical phase within the Mg particles was the liquid pools.The melting of the Mgrichαphase surrounding the Al-richβphase decreased the solubility of Al in the liquid,and thus,the Al content in the liquid was lower than that in the eutecticβphase (Table 1).Moreover,the solutes (such as Al) concentrated in the formed liquid tend to dissolve toward the surroundingα-Mg as the heating time goes on.On the other hand,the liquid pools are driven to coalesce because the solid/liquid interfacial energy decreases with coalescence.As a consequence,the amount of liquid pools decreased while their size increased.Therefore,the dominant phenomenon during the initial heating period of 0-10 min was the coarsening ofα-Mg grains through the rapid dissolution of the eutecticβphase.

Once the heating time exceeded 10 min,liquid phase formed around the particles (marked by arrows in Fig.4c)and is more clearly shown in the inset image in Fig.4c.The amount of liquid phase increased gradually as the heating time increased by comparing Fig.4c,d.XRD patterns of the partially remelted samples reveal that the intensity ofβphase peaks decreased at first,and then,its intensity increased gradually after 10-min heating (Fig.6).The progressive increase in theβphase content reached a maximum at about 30 min and held constant after that.As the heating time went on further,theα-Mg particle edges partially remelted through theα→L reaction as the increased temperature increased the amount of liquid,so that neighboring particles became separated by the liquid phase.The more the liquid phase in the sample was,the more the solidifiedβphase will be after quenching in water.The extent of theβphase was gradually enhanced after 10-min heating.Theαphase melting further reduced the solubility of Al in the liquid phase,while the Al content in the primaryα-Mg particles increased.Figure 7 shows the variations in the liquid fraction during partial remelting.The liquid fraction remained nearly constant when the heating time exceeded 30 min,indicating that the whole system reached its final solid/liquid equilibrium state.Correspondingly,the solubility of Al in the liquid and the amount ofβphase also remained almost constant.

Fig.4 SEM images of (CNTs+Si_p)/Mg powder compacts heated at 585℃for different durations and then water-quenched.a 5 min,b 10 min,c 20 min and d 30 min

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Table 1 Chemical compositions of different structures of (CNTs+Si_p)/Mg powder compacts during partial remelting at 585℃

Fig.5 Variation of temperature in specimen with heating time

Fig.6 XRD patterns of (CNTs+Si_p)/Mg powder compacts heated for different durations

Fig.7 Variation of primary particle size and liquid fraction with heating time

Compared to that of the (CNTs+Si_p)/Mg mixed powders,the intensity of the Si diffraction peaks decreased for the heated samples,until this peak disappeared completely after 60-min heating.Correspondingly,the peaks indicating the Mg2Si phase appeared as the heating time increased (Fig.6).The C peaks corresponding to CNTs without carbides,such as Al₄C₃,appeared during partial remelting.The Raman spectra in Fig.8 indicate the change of carbon structures.The Raman characteristics of the CNTs show the G band at 1563.65 cm^-1,D band at1334.162 cm^-1 and 2D band at 2672.518 cm^-1.The D band reflects that defects in the CNTs and the G band are caused by in-plane tangential stretching of the C-C bonds [ 10] .The ratio of the intensities of the D and G bands (I_D/I_G) is usually used to examine the defects in the carbon structure.As more defects arose,theI_D/I_G values increased correspondingly.TheI_D/I_G ratio increased from 0.801 to1.006 after ball milling and increased smoothly to 1.147after the (CNTs+Si_p)/Mg powder compact was heated for60 min.These changes indicate that the ball-milling and partial remelting procedures damaged the structure of CNTs.Compared with the popularly used high-energy ballmilling method,the blending process in this work caused less structural damage because of its lower energy input [ 22] .On the other hand,the G band of CNTs in the partially remelted (CNTs+Si_p)/Mg compact shifted slightly to a higher wave number than that of the as-received CNTs.This change was related to residual stress and the CNTs/Mg interfacial bonding condition [ 23, 24] .Therefore,in situ synthesized Mg₂Si_p were successfully obtained due to the reaction between Si_p and Mg matrix during partial remelting and the corresponding hybrid reinforcements of the Mg matrix composites were CNTs and Mg₂Si_p.Moreover,no obvious structural damage to the CNTs was induced and a relatively intact carbon structure remained after partial remelting.

Fig.8 Raman spectra of as-received CNTs,(CNTs+Si_p)/Mg mixed powders and (CNTs+Si_p)/Mg powder compact heated for 60 min

Figure 9 shows EDS results from (CNTs+Si_p)/Mg powder compacts heated for different time.The distribution of Al,C,and Si was homogeneous when the heating time was 10 min (Fig.9a).As discussed above,the main phenomenon was the rapid coalescence ofα-Mg grains due to the dissolution of the eutectic phase during the period from 0 to 10 min.These findings indicate that the added nanoreinforcements were homogeneous if no intergranular liquid phase formed.As the heating time went on,Al was seriously segregated from other materials in the samples,confirming that liquid phase surrounding the primary particles had formed (Fig.9b).However,C and Si were also distributed homogeneously in the liquid phase,demonstrating that CNTs and Mg₂Si_p remained homogeneous during partial remelting.To further analyze the distribution of hybrid reinforcement during partial remelting,the fracture surfaces of the (CNTs+Si_p)/Mg powder compact were examined,as shown in Fig.10.The Mg particles remained spherical after the ball milling and cold pressing(Fig.10a).The CNTs were distributed uniformly without clumping on the surface of the Mg particles and some CNTs were embedded into the Mg matrix.Moreover,Si_palways distributed around the CNTs.As the heating time increased,the reinforcements were still homogeneously distributed (Fig.10b),consistent with the EDS results shown in Fig.9.Therefore,the dispersion homogeneity of CNTs and Mg2Si_p was maintained after partial remelting,which is indispensable for the preparation of high-performance composites.

Fig.9 EDS surface scanning analysis of (CNTs+Si_p)/Mg powder compact heated for a 10 min and b 30 min

Fig.10 SEM fractured images of (CNTs+Si_p)/Mg powder compacts heated at 585℃for a 0 min and b 40 min

Quantitative measurements indicate that the particle coarsening rate (0.33μm·min^-1) of the (CNTs+Si_p)/Mg powder compacts was low during the period from 0 to10 min,consistent with the SEM results shown in Fig.4that the size of the Mg particles almost remained unchanged.Theα-Mg particles were compactly contacted together with each other at many sites,and the distances between them were quite short during this period.To minimize the interface energy,the contacted particles should coarsen through coalescence.MgO layers are known to inevitably exist on the surfaces of Mg alloy powders,which hinders the merging of neighboring Mg particles.The specific area of the nanoreinforcements is much larger than that of micro-reinforcements at the same concentration,so the CNTs and Si_p reinforcements more effectively separate neighboringα-Mg particles from each other.This separation decreases the contact area between particles and inhibits grain growth through grain boundary migration.As a result,the coarsening rate of the particles was slow during the initial heating stage.Figure 11 shows micrographs of the AZ91D powder compacts heated at 585℃for different time.After 5-min heating,the primary particles of the AZ91D powder compact (43.5μm) were larger than those of the hybrid powder compact (36.8μm),demonstrating that the dispersed reinforcements indeed inhibited grain growth.The primary particle size did not increase sharply after heating for 10 min,and the average coarsening rate reached 0.06μm·min^-1 (Figs.4,7).Moreover,primary particles of the hybrid compact(39.8μm) were also smaller than those of the AZ91D compact (46.4μm)(Fig.11).The coarsening rate was rather slow compared with that of conventional as-casting materials during the heating process [ 25] .As the heating time further increases,the particles were increasingly separated by liquid phase due to the higher temperature,which increased the distance between the particles.Correspondingly,the coarsening of the particles transitioned to a mild regime of Ostwald ripening.Ostwald ripening is a process by which a semisolid system lowers its energy by decreasing the solid/liquid interfacial area through the dissolution of small particles and the growth of large particles.This process is governed by the diffusivity of atoms in the system [ 26] .If the Ostwald ripening regime is in effect,the variation of particle size with heating time is expressed as:

Fig.11 SEM images of AZ91D powder compacts heated at 585℃for a 5 min and b 30 min

where D_t and D₀ are the particle size at time t and the beginning of heating,respectively,and K is a constant.Figure 12 shows the calculated results,taking 10 min as the starting time,t=0.That is to say,the value at 10 min corresponds to that at 0 min in Fig.12.The calculated values are far from the linear fitted to the observed values,implying that Ostwald ripening was not the dominant coarsening behavior occurred during 10-40 min.MgO layers on the Mg alloy powder surface discouraged atom diffusion.On the other hand,the addition of reinforcements decreased the effective diffusion coefficient of the solute atoms within the liquid phase [ 27] .The effect of Ostwald ripening was thereby impeded to some extent,and thus,the linear fit does not match the calculated values well.Therefore,the primary particles were slightly coarsened due to the dominance of the pinning effect of dispersed reinforcements on grain-boundary migration and the impeded Ostwald ripening regime.We conclude that an asreceived Mg powder nearly evolved into a spheroidal particle surrounded by liquid phase after partial remelting.

Fig.12 Cube of primary particle size with heating time (taking10 min as starting time)

Pores are responsible for the microstructure,strength and other properties of press-formed ingots,so pore evolution during partial remelting should be analyzed in detail.As shown in Figs.4,10,many intergranular pores appeared between the powders.The powders bonded together only mechanically in the green compact,so many pores inevitably formed.As the heating time went on,the number of the intergranular pores decreased clearly (Fig.10b).These variations are clearly demonstrated in the quantitative examination results shown in Table 2.The relative density of the powder compact increased from 55.1%before partial remelting to an approximately constant value of 79.6%after heating for 40 min.The densification rate increased gradually over the heating period.Pore filling is known to be an essential process for densification and is obviously determined by the amount of liquid in the powder compact and the balance between the ambient pressure,the liquid pressure,and the pressure exerted by the liquid meniscus due to tension at the gas-liquid interface [ 28, 29] .As the liquid fraction increased,liquid flow was increasingly sufficient to fill in the interstitial spaces among the primary particles,eliminating the small internal preexisting pores and increasing the relative density.The curvature of the liquid meniscus decreased as the pores were filled by the enhanced liquid,so the surface tension increased,resulting in an imbalance in the liquid pressure [ 30] .The liquid pressure was determined by the curvature of the smallest remaining pores,and smaller curvature corresponded to lower liquid pressure.Thus,the liquid pressure decreased as the small pores were eliminated.At some critical point,the pressure due to the contribution of interfacial energy balanced the liquid pressure,which impeded the densification process if the grain size remained constant.As discussed above,the primary particle size increased gradually with the heating time.As the primary particles grew,the interfacial energy made less contribution to densification,so that the relative density increased gradually over the heating time.However,microstructure densification remained relatively constant after the heating time exceeded 30 min.This plateau could be related to the material reaching a final solid/liquid equilibrium with slower coarsening after 30-min heating.Therefore,microstructure densification also occurred during partial remelting.

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Table 2 Variation in relative density with heating time

4 Conclusion

The microstructural evolution during partial remelting of(CNTs+Si_p)/Mg powder compacts prepared by cold pressing was investigated in this study.The evolution processes can be pided into three stages:rapid grain coarsening through the dissolution of eutectic phases,formation of the liquid phase due to partial remelting of the Mg powders,and the final slight coarsening.An as-received Mg powder in the green compact nearly evolved into a spheroidal particle surrounded by liquid phase after partial remelting.The in situ synthesized Mg2Si_p were distributed homogeneously around the CNTs with wellmaintained structural integrity during partial remelting.The relative density of the semisolid ingot increased during the heating period between 0 and 30 min and then basically remained constant.Partial remelting of the (CNTs+Si_p)/Mg powder compact at 585℃for 30 min yields an ideal semisolid ingot with fine and spheroidal a-Mg particles suspended uniformly in the liquid phase.