Preparation of semi-solid billet of magnesium alloy and its thixoforming
镁合金英文综述Review of mechanical behavior and microstructure of Magnesium alloy
Review of mechanical behavior and microstructure ofMagnesium alloyAbstract: Magnesium alloys are introduced in this article. The mechanical behavior and microstructure of Magnesium alloy are discussed in the review. The characteristics of Magnesium alloy are researched by researchers. The mainly deformation mechanisms of Magnesium alloy are slip and twinning which determined by the grain structure of magnesium. There is a great relationship between mechanical properties and microstructure of Magnesium alloy. And there are many ways to improve the mechanical properties of Magnesium alloy by grain refinement. Superheating, carbon inoculation, the elfinal process, control of impurity level, zr addition, other element additions, rapid solidification and physical grain refining are illustrated in this review, and all those can be used to refine the grain of Magnesium alloy.Key words: Magnesium alloy; Microstructure; Deformation; Strength; Grain refinement 1.IntroductionMagnesium alloys have been received a great attention as light-weight structure materials because of specific strength, high stiffness, good damping capacity andeasy-recycling and so on[1]. Magnesium is the lightest structural metal with a densityof only 1.738 g/cm3 at 20℃[2]. For engineering applications, magnesium is usually strengthened by alloying mechanism; it can be alloyed with other alloying elements such as aluminum, zinc, manganese, zirconium and rare earth.[3]Contain of various ingredients of magnesium alloy are largely studied by scientific researchers. Magnesium alloys containing rare earth elements are known to have high specific strength, good creep and corrosion resistance up to 523K. The addition of SiC ceramic particles strengthens the metal matrix compo site resulting in better wear and creep resistance while maintaining good machinability [4]. Kawamuraet al.[5] have developed a RS P/M Mg-1Zn-2Y alloy, and this alloy shows excellent mechanical properties. Liu et al.[6] investigated the thixoformability in alloys based on the Al–Si–Cu and Al–Si–Cu–Mg systems using MTDATA thermo-dynamic and phase equilibrium software combined with the MTAL database. Criteria for thixoformability are identified and a range of alloy compositions based on Al–Si–Cu andAl–Si–Cu–Mg evaluated in relation to these criteria. Birol[7] studied the thixoformability of AA6082 aluminum alloy reheated from the as-cast and extruded states, respectively. The thixoformability of the as-cast alloy was inferior with respectto that from the extruded material. Camacho et al.[8] studied the wrought alloy compositions amenable to semi-solid processing, using a commercial software package MTDATA, NPL alloy solution database MTSOL and SGTE substance database. Commercial thixoforming is generally based on conventionalaluminum-based casting alloys such as A356 and A357, which provide high fluidity and good castability[9].2. DeformationFor magnesium alloys, slip and twinning, are well known to be two major orientation-dependent deformation mechanisms. Both basal slip (with a 1/3〈112＿0〉Burgers vector) and non-basal slip (e.g. first-order {101＿0} prism slip and {101＿1} pyramidal slip) systems have been reported extensively. Moreover, as all the slip systems mentioned above cannot produce plastic deformation parallel to thec-direction, twinning usually plays an important role in the plasticity of these materials[10].Twinning and slip in hexagonal close-packed structures have been extensively studied using molecular dynamics. Barrett et al.[11] utilized MD simulations to explore slip and twin nucleation mechanisms and their sensitivity to Schmid and non-Schmid stresses by first loading a defect-free crystal having full periodic boundary conditions under various uniaxial loading directions.The deformation temperature is an important factor to Magnesium. Li et al.[12] investigated the effects of deformation temperature on microstructure and mechanical properties of AZ80 magnesium alloy. The mechanical properties and microstructure were carried out in Gleeble-1500 thermal simulation experiment and optical microscope. The extrusion deformation, dynamic recrystallization had taken place in all the deformation samples, grains were thinner than before deformation. The reasonable deformation process can make the dynamic recrystallization organization of grain smaller and obtain higher strength. The best deformation temperature was about 360 degrees C to 390 degrees C to AZ80 magnesium alloy [12]. Xu et al. discussed the effect of microstructure and texture on the mechanical properties of the as-extruded Mg–Zn–Y–Zr alloy specimens at room temperature [13].Lu et al. [10] focuses on the monotonic and cyclic behavior of a high-pressure die cast magnesium AM60B alloy. The mechanical results are discussed with respect to the microstructure in terms of clusters of pores, grain size and theorientation-dependent activation of different deformation mechanisms.Shi et al. [14] investigated the compression of a semi-solid Zn-Al alloy disc as it is often used as a filler metal to braze aluminum alloys and their composites. Three different size discs were used with height-to-diameter ratios (hid) of 0.6, 0.3 and 0.1. Stress-strain curves were obtained during disc compressions. The maximum stress obtained during the compressions increased with a decrease in disc size (hid). Hagihara et al. [15] investigated the influence of a change in the stacking sequence of the close-packed plane in a Mg12ZnY long-period stacking ordered (LPSO) phase on its mechanical properties. A 14H-typed LPSO-phase crystal was fabricated by annealing a directionally solidified (DS) crystal with a 18R-typed LPSO-structure at 525 degrees C for 3 days, and the temperature dependence and orientation dependence of the yield stress were examined via compression tests. (0001)< 11 (2) over bar0 > basal slip was identified as a dominant deformation mode, and deformation kink bands were formed under compression in the case of suppression of basal slip motion. The deformation mechanism of the 14H-typed LPSO-phase is almost similar to that of the 18R-typed LPSO-phase, even though a slight differencewas observed at temperatures above 300 degrees C.3. StrengthThere many factors that affect strength of magnesium alloy. Accordingly, considerable approaches have been explored to attain higher strength and ductility on Magnesium alloys in recent years. Grain refining is an effective procedure for achieving high strength at RT together with possible superplastic forming capabilities at elevated temperatures for many face-centered cubic (fcc) metals[16].Severe plastic deformation will effectively result in significant grain refinement in many metals and this may be achieved using procedures such as equal-channel angular pressing (ECAP), accumulative rolling bonding (ARB) and high pressure torsion[17].In general, the strength and ductility of materials processed by ECAP cannot be readily enhanced simultaneously. Recently, several strategies were proposed to achieve the relatively high uniform elongation in those strong UFG fcc metals [18]. Wang et al. pro-posed that the UFG metals and alloys can exhibit a combination of high strength and good ductility by designing their grain size to form a bimodal microstructure in pure copper. Zhao et al. [19] reported that a large fraction of equilibrium HAGBs and low dis-location density could improve the toughness and the uniform elongation of UFG materials by imparting excessive processing plastic strain. Lu et al.[10] suggested that the strength and uniform elongation could be simultaneously ameliorated in the UFG Cu by inducing nanotwins. In all theabove-mentioned toughening strategies, the UFG microstructures can provide the high strength; while different microstructures (duplex grain sizes, equilibrium HAGBs and nano-twins and stacking faults) contribute to the ductility by improving thestrain-hardening ability[20].The microstructure, texture and tensile properties of this magnesium alloy before and after ECAP were systematically investigated at RT and the relationship between microstructure and mechanical properties was elucidated through detailed analysis[21].A particle-strengthened magnesium alloy, Mg–12Gd–3Y–0.5Zr, has been processed successfully by ECAP. The microstructural evolution was studied systematically by TEM. The relation-ship between strength and elongation was discussed in termsof its BU microstructure. The following conclusions may be drawn: (1) After four-passes ECAP, a BU microstructure, containing the matrix grains and the second-phase particles with the average sizes of 490 nm and 290 nm, respectively, was obtained.(2) The tensile strength and elongation of the Mg alloy with the BU microstructure can be simultaneously increased at RT.(3) The tensile strength increment of the BU microstructure can be mainly attributed to the combined influences of matrix grain refinement and enhanced dispersion strengthening. While the tensile ductility increment of the BU microstructure is closely related to the formation of profuse microscale shear bands.(4) The fracture mechanisms are attributed to debonding the inter-face between particle and matrix grain for the samples E0 with eutectic-phase particles at grain boundary and the linkage of microscale shear bands for samples E4 with BU microstructure[21].4. MicrostructureAlthough Magnesium alloys have outstanding strength/weight ratio, the important disadvantages of Magnesium alloys are low strength and low ductility compared with the other competitive structural materials such as Al and steel. It is well known that a finer grain size may contribute synchronously to the strength and ductility [3]. A good mechanical behavior of Magnesium alloy can be obtained through changing its microstructure. The mainly method can get excellent mechanical behavior is grain-refining method.Fine grain microstructure favors uniform deformation and improves isotropic mechanical properties of the materials with hexagonal close-packed (hcp) structure [22]. It is also well-known that, the microstructure prior to forging or extrusion, i.e. the solidified structures of an ingot, has a significant impact on the subsequent forging properties [22].Many techniques are available to achieve grain refinement. Among them, the post-solidification techniques involve deformation processing and severe plastic deformation techniques. Among the solidification processing techniques, rapid quenching, particle inoculation (chemically assisted), and use of physical means have shown promise.Nowadays, there have been various grain-refining methods developed for changing the microstructure of Magnesium alloy, such as superheating, carbon inoculation, the elfinal process, control of impurity level, zr addition, other element additions, rapid solidification and physical grain refining.4.1. Superheating methodThe superheating process was originally described in a British patent granted in 1931[23]. Aluminum bearing magnesium alloys benefit from high-temperature treatment in terms of grain refinement. This high-temperature treatment is usually termed as superheating and the process involves heating the melt to a temperaturewell above the liquidus of the alloy often in the range 453K to 573 K for a short time, followed by rapid cooling to, and short holding at, the pouring temperature. Although the grain refinement efficiency of superheating is subjected to many factors, there are some basic characteristics of this technique. Starting with, a significant grain refinement response can only be achieved in Mg-Al alloys with a minimum addition of Mn/Fe content. Then, a specific temperature range above the pouring temperatureis required to maximize the grain refining effect. Finally, rapid cooling from the overheating temperature to the pouring temperature and the short holding time are also crucial requirements to produce fine grains.A model has been proposed on the basis of the recent understanding of the grain refinement of both high purity and commercial purity Mg–Al alloys[23]. It simply involves heating a molten magnesium alloy to a temperature well above the liquidus of the alloy, holding it for a required period, and then cooling rapidly to the required pouring temperature. Figure 1 illustrates three different superheating cycles. Although extensive investigations have been carried out and a number of hypotheses have been proposed since the 1930s, the grain refinement mechanisms remain unclear.Understanding of the controlling mechanism will help foster the development of an effective grain refiner for Mg–Al alloys. This work proposes a new hypothesis for the grain refinement of magnesium alloys by superheating on the basis of the recent developments in grain refinement of magnesium alloys. The model is applied to elucidating the various phenomena observed about superheating. Schematic of typical temperature profiles during a superheating process of Mg–Al alloys are illustrated Fig.1[23].Fig. 1 Schematic of typical temperature profiles during a superheating process ofMg–Al alloys. The superheat temperature T sh is usually in the range of 850–900 °C. The pouring temperature T p is generally around 720 °C. Three different cooling conditions are shown: (1) rapid cooling from T sh to T p with a short holding time before casting; (2) rapid cooling from T sh to T p, but a lengthy holding at T p before casting; and (3) slow cooling from T sh to T p.[23]4.2. Carbon inoculationCarbon inoculation, which developed at the end of World War II, is another major grain-refining process for Mg-Al based alloys. This method is featured with low operating temperature, less fading, short processing time and crucible wear, and therefore favors practical applications [3].As for the grain refinement mechanism of carbon inoculation, the most commonly accepted theory is that Al4C3 particles formed in the Mg-Al melt act as effective nuclei for the Mg grain solidification. It is approved by the fact that the effective addition of carbon inoculant is only confined to Al-containing magnesium alloys. However, no experimental evidence that the Al4C3 particles act as the heterogeneous nuclei of primary α-Mg is observed by micrographs till now. Theapplication range of carbon inoculation method is limited because the grain refining mechanism cannot be understood clearly[24].In recent studies, some researchers proposed that the presence of Mn is necessary to form the heterogeneous nuclei for grain refinement of Mg-Al alloys. Therefore, the role of Mn in grain refinement of Mn containing Mg-Al based alloys should be further investigated to understand the grain refinement mechanism of carbon inoculant treatment. In this work, a novel MgCO3 contained carbon inoculant mixture was developed for grain refinement of AZ91D alloy. The grain refinement process and mechanism of this inoculant on Mg-Al alloy under different processing conditions were investigated experimentally[24].Carbon black is an easily available and inexpensive form of carbon that has nano size morphology. Present work investigates the inoculation potency of these nano particles in Mg–Al alloy melts [25]. It is noted that carbon inoculation grain refinement is only applicable to aluminum-containing magnesium alloys. Accordingly, some researchers put forward that the high-purity carbon powder or the magnesite particles should be added to replace harmful hexachloroethane in the carbon inoculation treatment.4.3. The Elfinal processThe Elfinal process has been invented by the metallurgists of a pioneering German magnesium company based on the hypothesis that iron particles can act as nucleation sites for magnesium grains[26]. It has been reported that Mg-Al-Zn alloys (Al: 4 to 8.5 pct; Zn: 0.5 to 3 pct; no other elements have been mentioned) can be grain refined by additions of 0.4 to 1.0 pct of anhydrous FeCl3 at a temperature range of 1013 K to 1053K. Though the approach has worked satisfactorily in terms of grain refinement but the inventors fail to convince other metallurgists about the mechanism behind it. Different mechanisms have been subsequently proposed. It has been suggested that Fe- containing intermetallic particles or aluminium carbide (Al4C3) particles are possibly the nucleants. According to Emley, hydrolysis of FeCl3 in the magnesium melt gives rise to copious hydrogen chloride (HCl) fumes, which then attack steel crucibles to liberate some carbon into the melt. The other major hypothesis proposed is that Mg grains nucleate on Fe-Mn-Al particles. A detailed examination of this process has been performed to clarify a number of key issues (i) whether Fe is a grain refiner or an inhibitor for Mg-Al alloys (ii) whether iron only grain refines Mg-Al alloys that contain Mn and (iii) the mechanism by which the Elfinal process works[27].For the work stated above, sublimed high-purity magnesium ingots (99.98%) and commercial high-purity aluminium ingots (99.999%) have been used to prepare high-purity Mg-3%Al and Mg-9%Al alloys. Melting has been conducted in an electrical resistance furnace under a protective cover gas of 1.0%SF6 in 49% dry air and 50% CO2.Aluminums titanite crucibles have been used for the reason that they are free of carbon. Anhydrous FeCl3 has been plunged into the melt at 1023 K. Cone sampleshave been taken from the top of the melt using a boron nitride coated cone ladle (Ø 20mm x Ø 30mm x 25 mm), 10 min following addition of FeCl3. No stirring has been applied in each test. The average grain size of each cone sample has been measured from the central region of a longitudinal section of the cone cut through the axis[3].4.4. Control of impurity levelAn interesting observation that has been made about the grain refinement of Mg-Al type alloys is the influence of the source magnesium impurity level. This native refinement in Mg-Al type alloys is said to have occurred when the native grain size is finer than that of commercial purity alloys. It is unclear whether native grain refinement of high purity Mg-Al alloys is conditional upon the C and Al contents. The difficulty of clarifying the role of carbon lies in the difficulty of how to accurately determine a trace level of carbon in magnesium alloy.In a recent work done to understand the mechanism of native grain refinement in Mg-Al alloys, the raw materials used are high-purity aluminium, commercial purity zinc and calcium, and two different sources of magnesium metal, which include sublimed high purity magnesium(99.98%) and commercial purity magnesium(99.7%).It has been found that, Mg-Al alloys with the same basic composition, but made of different sources of magnesium metal, showed an obvious difference in grain size, which is represented in Fig.2. High purity alloys consistently have proved a finer grain size than commercial purity alloys in all cases across the composition range0.5-9%Al. Fig. 3 are microstructural observations corresponding to Mg–9%Al alloys.Fig.2 Effect of source magnesium purity on the grain size of Mg-Al alloys[28].Fig.3. Grain structures in Mg–9%Al alloys made of (a) commercial purity magnesium metal, average grain size (AGS): 200 μm, and (b) high purity magnesium metal, AGS=140 μm.Native grain refinement was observed exclusively in high purity Mg–Al alloys. Mg–Zn and Mg–Ca alloy systems do not show native grain refinement, but rather native grain coarsening in the high purity alloys. The grain size of Mg–9%Al alloys was found to increase with an increase in the proportion of commercial purity magnesium metal used in making these alloys, i.e. impurity level, in the experimental range from 0% to 100% commercial purity magnesium.4.5. Zr addition and Other element additionsZirconium is a potent grain refiner for pure magnesium and is ineffective in magnesium alloys that contain Al, Mn, Si, Fe, Ni, Co, Sn and Sb as zirconium forms stable compounds with these elements[29]. When added to these alloys, where the maximum solubility of zirconium in molten pure magnesium at 927 K is ~ 0.45%, Zr can readily reduce the average grain size to about 50ìm from a few millimeters at normal cooling rates. Moreover, well-controlled grain refinement by Zr can lead to formation of nearly round or nodular grains, which further enhance the structural uniformity of the final alloy. This exceptional grain-refining ability of Zr has led to the development of a number of commercially important magnesium alloys including a few recently developed sand-cast creep resistant magnesium alloys that are aimed at automotive applications such as transmission cases and engine blocks[26]. The most characteristic feature of the microstructure of a magnesium alloy containing more than a few tenths per cent soluble Zr is the Zr-rich cores that exist in most magnesium grains. These Zr-rich cores are usually less than 15 P m in size at normal cooling rates. They are believed to be the products of peritectic solidification. In order to know the mechanism of grain refinement by Zr and capitalize on the grain-refining ability of Zr, it is required to understand the characteristics of these Zr- rich cores[30].At present grain refinement of these alloys is commercially carried out by the addition of a Zr- rich Mg-Zr master alloy, which contains Zr particles ranging from sub-micrometer to 50 Pm in size. It has been found that grain refinement of magnesium alloys by Zr is dictated by both soluble and insoluble Zr contents. However, Zr particles settle very faster in molten magnesium due to the significant difference between the densities of Zr and molten magnesium[31]. As a result, the average grain size increases obviously with increasing residence time of the melt prior to pouring. Moreover, once the Zr particles that are released from a Mg-Zr master alloy added to the melt settle to the bottom of the alloying vessel, little dissolution can be expected of these particles in the absence of stirring. Hence, the particle size distribution in a Mg-Zr master alloy can be understood mainly from a settling point of view rather than from the nucleation point of view. The identification of effective nucleant particles is commonly based on the assumption that after nucleation on any particle added to the melt latent heat release will decrease the likelihood of nucleation on neighbouring particles, which subsequently will be pushed to grain boundaries or into the interdendritic spaces[32]. Therefore, an effective nucleant particle is always expected in the central regions of grains. Compared to the grain refinement of mostother alloys, where it is usually difficult to find a large number of nucleant particles on polished sections, Zr-rich particles that have played a role as nucleation centers in a magnesium alloy can be readily distinguished using a SEM in the BSE image mode, due to the characteristic particle-core structures that form during solidification. Certainly, any information about the size distribution of these particles will help understand the potency and efficiency of Mg-Zr master alloy grain refiner, providing an important basis for improving the design of a grain refiner[33].In magnesium alloys, Zr element has relatively larger GRF value compared with other elements, so it possesses stronger grain refining ability as mentioned previously. Similar to Zr, Ca, Sr and Sb can be effective additions for refining grain size of magnesium alloys[34].4.6. Rapid solidificationIt is well known that rapid solidification processing (RSP) is an important grain refinement method[35]. There are two basic techniques for rapidly solidifying melts: substrate quenching and atomization. Substrate quenching refers to the solidification of the melt against one or two surfaces at a lower temperatures (e.g. room temperature, or near liquid nitrogen)[3]. Substrate quenching includes thermal spray methods,melt-spinning technique, planar flow casting, copper mold casting, twin rolling etc. Atomization is a process of breaking up a molten stream of liquid into small spheres by using gas et. Gas atomization includes high pressure and centrifugal gas atomization et. In substrate quenching, rapid solidification is achieved by increasing the rate of heat extraction and in atomization by increasing the amount of undercooling before nucleation. An average grain size of 0.2- 3 µm can be achieved in the rapid solidification of Mg alloys, and the rapidly solidified Mg-Al-Zn system presented an outstanding ultimate tensile strength of about 500 MPa[5].Besides microstructure refinement, RSP can effectively extend solid solubility in magnesium, for example 1.5 times for Mg-Ag and about 1,000 times for Mg-Ba alloys. The combination of grain refinement and solid solution hardening effect makes RSP a suitable technique for enhancing the mechanical properties and corrosion resistance of Mg alloys[36]. To fabricate structural components, subsequent thermal mechanical processing (e.g. extrusion, forging or rolling and consolidation) is necessary. Depending on the working temperature and processing rate, such hot working significantly impacts the structure of the as-solidified Mg alloys. It should be pointed out that RSP of magnesium alloys poses critical challenges due to the high chemical reactivity of magnesium[37].4.7. Physical grain refining methodsPhysical grain refining methods involve promoting nucleation, dispersion and multiplications of solidified crystals under mechanical force or external physical field without any further chemical additions. Physical grain refinement generally targets creating a favorable condition for nucleation and nuclei survival or breaking thesolidified crystal structures[3].The creation of an ideal condition for nucleation and ensuring high nuclei survival has been employed as a physical grain refinement strategy in the present investigation[38]. Consequently, physical grain refinement increases effective nucleation by tailoring the solidification conditions without necessitating the addition of inoculants. Casting near the liquidus temperature has been known to promote fine equiaxed microstructure. There has been significant controversy in explaining the columnar to equiaxed transition in castings without grain refiner addition. In a comprehensive overview, Hutt and StJohn have discussed the five major available theories and critically assessed the applicability of the proposed mechanisms. It has been concluded by the authors that all proposed mechanisms or a combination of them may be operative depending on the alloy composition, casting conditions or the types of nucleating substrates present[39]. A similar comprehensive analysis of CET and the plausible mechanisms have been discussed by Flood and Hunt. Both of these reviews suggest that in the absence of grain refiner (where constitutional supercooling driven nucleation is important), big bang (also known as free chill crystal or wall mechanism) and dendrite detachment mechanisms are the primary contributors to the creation of equiaxed grains. During low superheat casting the convection associated with the mould filling remains strong as solidification commences. Although it is argued that deformation or melting of the dendrite arms is promoted by the fluid flow, the big bang mechanism becomes progressively important as the melt superheat is reduced[40].5. SummaryThe mechanical behavior of Magnesium alloy has relationship with its microstructure. So it can enhance its deformation ability through refining grain of Magnesium alloy in microstructure, which fine grain size can result in structural uniformity and enhance the mechanical properties, hence improving the service performance of the products. For Magnesium alloys, many grain refinement methods have been developed, but their refining mechanism are still unclear. For example, as the effective grain refinement method, there still is debate in the heterogeneous nuclei for superheating and carbon inoculation of aluminum-containing magnesium alloys. Further investigations are needed for a more comprehensive understanding of the grain refining mechanism, and to develop reliable commercial grain refiners or novel grain refinement processes.References[1] LIU K, MENG J A. Microstructures and mechanical properties of the extrudedMg-4Y-2Gd-xZn-0.4Zr alloys [J]. J Alloy Compd, 2011, 509(7): 3299-3305.[2] ANILCHANDRA A R, BASU R, SAMAJDAR I, et al. Microstructure and compressionbehavior of chip consolidated magnesium [J]. J Mater Res, 2012, 27(4): 709-719.。
AZ91D镁合金半固态压铸工艺研究的开题报告
AZ91D镁合金半固态压铸工艺研究的开题报告
一、研究背景与目的：
AZ91D镁合金作为一种轻量化金属材料，其在航空、汽车、电子等领域具有广阔的应用前景。
半固态压铸是一种新型的金属成形技术，具有高效、低污染、低能耗、高精度等优点，主要适用于高性能铝合金和镁合金的成形加工。
因此，研究AZ91D镁合金半固态压铸工艺对于提高制造工艺的精度和效率，推动轻量化产业的发展具有重要意义。
本文旨在研究AZ91D镁合金半固态压铸工艺，探究半固态压铸工艺对于材料性能和成形质量的影响，以及制造工艺的优化和改进方法。
二、研究内容与方案：
1.半固态压铸工艺参数的确定：通过实验和模拟计算，确定合适的半固态压铸工艺参数，包括半固态成分比例、温度、压力等。
2.半固态压铸成形质量的评价：通过金相分析、机械性能测试等手段，对半固态压铸工艺成形质量进行评价，探究半固态压铸工艺对于材料性能的影响。
3.制造工艺的优化：在半固态压铸工艺的基础上，探究制造工艺的优化和改进方法，提高生产效率和制造精度。
四、研究意义：
1.推动轻量化产业的发展，提高铝合金和镁合金材料的应用水平。
2.优化半固态压铸工艺，提高成形质量和制造效率，减少能耗和污染。
3.为相关行业提供技术支持和指导，促进产业升级和技术创新。
三、预期结果：
1.确定合适的半固态压铸工艺参数。
2.研究半固态压铸工艺对于材料性能的影响，探究工艺优化的方法。
3.提高镁合金半固态压铸的成形精度和效率。
材料工程外文翻译--Mg-9Al-1Zn合金的半固体挤压成型
Semisolid extrusion moldingof Mg-9%Al-1%Zn alloysF. CZERWINSKIDevelopment Engineering, Husky Injection Molding Systems Ltd., Bolton,Ontario, L7E 5S5, CanadaA novel technique in manufacturing net-shape components of magnesium alloys, which combines semisolid processing, extrusion and injection molding, is outlined. For an Mg-9%Al-1%Zn composition, the high-temperature transformations and factors controlling solidification microstructures, are analyzed. C _ 2004 Kluwer Academic Publishers1. IntroductionExtrusion is the plastic deformation process by which a metal is forced to flow by compression through the die orifice of a smaller cross-sectional area than that of the original billet. Since the material is subjected to compressive forces only, the extrusion is an excellent method for breaking down the cast structure of the billet with little or no cracking [1]. Most metals are extruded hot when the billet is preheated to facilitate plastic deformation, but room temperature (cold) extrusion is also exercised. So far, conventional extrusion applications do not utilize preheating materials above the solidus temperature to enter the semisolid range.The advantages of processing metallic alloys in a semisolid state are attributed to the globular solid particles which control their thixotropic properties at high temperatures and reduce the content of dendritic forms after subsequent solidification [2]. It is well established that the benefits associated with semisolid processing, such as low shrinkage porosity, high tolerances and energy savings, are more evident at high solid fractions. Moreover, the ability to cast at higher solid fractions is of interest in improving billet stability and minimizing material loss during handling.Of all semisolid technologies, injection molding provides the largest flexibility in terms of the processed solid contents [3]. This feature is attributed to the fact that injection molding combines the slurry making and component forming operations into one step, and the slurry is accumulated in a direct vicinity of the mold gate. So far, these potentials are not explored and commercial applications are limited to liquid-rich slurries, which, for thin-wall sections, may contain solid volumes as low as 5–10%. As the major obstacle preventing using high solid contents, the premature alloy’s freezing and incomplete filling the mold cavity, is reported [4]. It was,therefore, anticipated that a drastic increase in solid content, especially above 60%, would transform the flow through the machine nozzle, runners, and mold gate into the extrusion, thus activating interaction between solid particles within the slurry which would facilitate the mold filling. The verification of such a hypothesis was the objective of this study.2. Experimental detailsAZ91D magnesium alloy, used in the present study, had a nominal composition of 8.5% Al, 0.75% Zn, 0.3% Mn, 0.01% Si, 0.01% Cu, 0.001% Ni, 0.001% Fe and an Mg-balance. An as-cast ingot was mechanically converted into small chips and processed using a Husky TXM500-M70 prototype system with a clamp force of 500 tons and a 2 m long barrel with a diameter of 70 mm. The component manufactured represented the complex shape with a diameter of 190 mm and a total weight, including sprue and runners, of 582 g [3]. The mold was preheated to 200◦C and the slurry was injected at a screw velocity in the range of 0.7–2.8 m/s. For the gate opening of 221.5 mm2 it converts to the alloy’s velocity at the mold’s gate between 12.2 and 48.6 m/s. In order to examine the role of flow through the gate, the alloy was also injected (purged) into the partly open mold at significantly lower flow velocity at the mold gate. The typical cycle time was approximately 25 s, which corresponds to an average residency time of the alloy within the machine barrel of the order of 100 s. In some cases, the cycle time was deliberately extended up to 4 times. Metallographic samples of the molded alloy were prepared by grinding with progressively finer SiC paper, mechanical polishing with 1 μm diamond paste and colloidal alumina, followed by etching in a 1% solution of nitric acid in ethanol. Stereological analysis was conducted using optical microscopy, equipped with a quantitative image analyzer.3. Results3.1. Structural transformations of the alloy during processingMorphologies of an as-chipped alloy are shown in Fig. 1a. According to size determination by the screen method ASTM E-276-68, the predominant fraction of chips was retained on sieves with openings within the range 0.6–2 mm, and 75% of them did not pass through the 1.4 mm sieve. As a result of interaction with the chipping tool, the alloy experienced a cold work. The chips’ deformation is inhomogeneous with an increased strain in an immediate location of the second phase particles (Fig. 1b).(a)(b)(c)Figure 1The initial state and thermal decomposition of Mg-9%Al-1%Zn feedstock, used during experiments: (a) as-received chips, (b) chip’s crosssection with cold-work features and (c) early stage of chip’s melting, showing disintegration of equiaxed network of recrystallized grains; the chemical segregation contour of former dendritic features, consumed by equiaxed grains is marked as “s”.Figure 2 The schematic diagrams of structural transformations of a magnesium alloy during various stages of the semisolid extrusion molding.As proven by chips melting outside the molding system, dendritic structure disintegrated completely by the solid-state reaction (Fig. 1c). During heating inside the machine barrel, the structure recrystallized by nucleation and the growth of equiaxed grains. The second phase, intermetallic compound Mg17Al12, was distributed mainly along grain boundaries. In addition, grain boundaries were enriched in the solute element Al. After exceeding the solidus temperature, the melting started at grain boundaries leading first to the generation of equiaxed, thenglobular solid particles, surrounded by the liquid metal (Fig. 1c). The summary of major structural transformations during semisolid extrusion molding is shown in Fig.2. Thus, during further conveying of the slurry along the barrel, the globular structures of the unmelted phase experienced breakdown and agglomeration due to the combined effect of external heat and strain. As a result of diffusion, the solid phase was also subjected to coalescence and Ostwald ripening, as discussed in detail previously [5]. After melting of the grain boundary network, the semisolid slurry, with globular solid particles, was essentially ready for the component-forming step,i.e. the injection into a mold cavity.3.2. Solidification microstructuresThe microstructure, typical for ultra-high solids is comprised predominantly of unmelted particles of α-Mg, surrounded by a solidification product of the former liquid phase. As seen in Fig. 3a, the particle’s shape is near globular, but with increasing solid fraction it shows a tendency to be more constrained geometrically, exhibiting shape accommodation. The former liquid cover the grain boundary network with small pools accumulated at triple junctions. There were also randomly distributed larger islands of the liquid; however, they were generally smaller than the solid particle. The portion of the former liquid which was entrapped within solid particles showed high dispersion. Thus, instead of single island morphologies, frequent for medium and lower solid fractions [6], there were numerous randomly distributed precipitates of smaller size and globular shape. It is likely that some of them were formed by solid-state precipitation from the supersaturated solid solution than by solidification of the Al-rich liquid.3.3. Microstructural feedback towards the high-temperature stageThe detailed examination of room-temperature microstructures provided important information regarding phenomena which take place at high temperatures. A strong influence of alloy injection velocity on the solidification morphologies was discovered. An example in Fig. 3b represents the structure formed from the slurry with the same initial solid content as that in Fig. 3a. The major processing difference was approximately a 100% increase in the mold gate opening and component wall thickness. The structure in Fig. 3b solidified in 4 mm-thick sections. Thus, reduction in injection velocity, combined with increased component wall thickness, led to the transition from globular particles surrounded by freshly solidified matrix to equiaxed grains. Some occasional islands of the former liquid were distributed at triple junctions. There is no substantial difference in the volume fraction and dispersion degree of precipitates within grain interiors, as compared to those present withinglobular solid particles (Fig. 3a).The quantitative description of the influence of the injection rate on the molded structure is shown in Fig. 4. The first finding is that for a given initial solid content within the slurry, the higher injection velocity led to an increased liquid content inside the part. In addition, the injection velocity affected the homogeneity of solid particle distribution, particularly across the part thickness. While for the gate velocity of 12.2 m/s the solid fraction across the part thickness of 2 mm remained constant, the velocity of 48.6 m/s caused an increase in the liquid fraction towards the component outer surface. For 12.2–48.6 m/s range of injection velocities, there was an increase in average liquid content from 15 to 25%. It seems that there was not close relationship between the solid content or its distribution and a size of the unmelted phase, which is characterized by the histogram in the insert of Fig. 4. The solid particle size was, however, strongly affected by the alloy’s residency time inside the machine barrel. An increase of average cycle time by four times, resulting in a total residency time of 400 s increased the average particle size from 33 μm to 60 μm.4. Discussion of resultsThe key requirement of semisolid processing is the thixotropic slurry with(a)(b)Figure 3The microstructure of AZ91D alloy obtained from the slurry with the same solid fraction: (a) globular structure, injection velocity of 48.6 m/s, wall thickness of 2 mm and (b) equiaxed structure, injection velocity of 5 m/s, wall thickness of 4 mm.non-dendritic morphologies of the unmelted phase. Our finding [6, 7] that the cold work, imposed on chips during their manufacturing, is a driving force for the generation of globular forms by the mechanism of recrystallization and grain boundary disintegration implies that such a slurry is formed at the very beginning of melting (Fig. 1b and c). Since all benefits are associated with the unmelted solid phase it would appear that the liquid content should be kept as low as possible. However, a certain liquid content is required to ensure processibility, defined as the macroscopically homogeneous and damage free flow. This minimum depends on the alloy type, and for the aluminum alloy A 2014, deformed in an unconstrained compression, it was as low as 20% [8]. In the case of this study, the minimum value was even lower (Fig. 4).It is generally accepted that, at ultra-high solid contents, the alloy represents a deformable, semi-cohesive granular solid, saturated with liquid. When subjected to external strain, the alloy will respond by the disagglomeration of partially bonded grains. It is highly possible that, due to the shear imposed by the screw and relatively short rest time, there is no bond between the solid particles accumulated directlybefore injection. Analyses of as-solidified structures (Figs 3a and 4) and machine operating parameters suggest that phenomena of solid/solid interaction within the slurry with ultra high solid contents are significantly more intense than those described for low and medium solid contents. During semisolid injection molding, the mold filling time is the key factor which controls the entire process [3]. If the material experiences solidification (freezing), it reduces the cross section of the flow channel and increases the effective filling time. It is suspected that the interparticle interference, at the stage of injection, facilitates the mold filling process. This hypothesis is supported not only by the complete mold filling, but also by the fact that the mold filling time was approximately 0.025 s, which is of the same order of magnitude as that measured for low solid contents.Figure 4The inhomogeneity in the distribution of solid particles within the component, extrusion molded in a semisolid state. The inset shows a typical histogram of solid-size distribution at middle-wall thickness.The characteristic feature of microstructures produced by semisolid extrusion molding is the small size of the unmelted phase (Fig. 4, inset). This size being of 34μm, is very similar to the grain size within the recrystallized chip and was believed to be preserved by external shear causing disagglomeration [5, 9]. An increase in the cycle time to 100 s mainly affected the alloy’s residency time in an absence of shear. It is natural that the system reduces its energy by particle growth, althoughmechanisms may differ from those described for low solid contents. At solid fractions above 0.5, coarsening behaviour can better be described by considering the migration of the liquid films separating the grains, than by considering diffusion fields around isolated solid grains, as in the Lifshitz, Slyozov and Wagner analysis [10]. Coarsening cannot be used exclusively to explain the formation of the equiaxed structure in Fig. 3b. Rather, the reduced gate velocity and longer solidification time, caused by thicker alloy section, provide a better answer. According to this mechanism, solidification of the liquid portion of the alloy takes place on pre-existing globular substrates as described earlier for large slurry volumes, which solidified inside the machine barrel [9].The inhomogeneities in the solid phase distribution of as-solidified structures (Fig. 4) result from the flow characteristics during mold filling. The higher solid content in the part than within the runner should be interpreted as resulting from melting during slurry flow through the narrowgate channels. Similar findings were reported for experiments of the back extrusion of Sn-Pb alloy with an effective liquid fraction of less than 0.30, where it was found that the liquid fraction was a function of the extrusion ratio increasing as the wall thickness decreased [11]. Another possibility of the selective flow where the liquid fraction is pushed through the solid skeleton seems to be diminished, from the microstructural analysis of samples taken along the alloy flow path. The relatively high flow velocity of the alloy during mold filling and specific gravity difference between solid and liquid causes phase segregation. For an injection velocity of 49.6 m/s, the liquid content close to the outer surface is about 20% higher than in the center (Fig. 4). Details of the inhomogeneous distribution of the solid phase within the part still requires explanation.5. ConclusionsThe elements of semisolid processing, extrusion and injection molding were successfully combined to manufacture net shape components of Mg-9%Al-1%Zn alloy starting from a slurry with high solid contents of the order of 70%. The resultant solidification microstructures ranged from globular forms surrounded by the former liquid matrix to exclusively equiaxed grains.At high temperatures, the unmelted particles were susceptible to coarsening with increased residency time and the homogeneity in their distribution within the part was influenced by the injection velocity. It is believed that the phenomenon of solid particle interference during slurry flowunder compressive forces facilitated filling the mold cavity.References1. G. E. DIETER, “Mechanical Metallurgy” (McGraw-Hill, New York, 1976).2. M. C. FLEMINGS, Metall. Trans. A 22 (1991) 957.3. F . CZERWINSKI, Adv. Mater. Proc. 160/11 (2002) 31.4. D. M. WALUKAS, R. E. VINING, S . E. LEBEAU, N. TANIGICHI and R. F . DECKER, Advanced Semisolid Processing of Alloys and Composites, edited by Y. Tsutsui, M. Kiuchi andK. Ichikawa, Tsukuba, Japan, 2002, p. 101.5. F . CZERWINSKI, Scripta Mater. 48 (2003) 327.6. F . CZERWINSKI, A. ZIELINSKA-LIPIEC, P . J . PINET and J . OVERBEEKE, Acta Mater. 49 (2001) 1225.7. F . CZERWINSKI, ibid. 50 (2002) 3265.8. E. TZIMAS and A. ZAVALIANGOS, ibid. 47 (1999) 517.9. F . CZERWINSKI, Metall. Mater. Trans. A 33 (2002) 2963.10. E. D. MANSON-WHITTON, I . C. STONE, J . R. JONES,P . S . GRANT and B. CANTOR, Acta Mater. 50 (2002) 2517.11. T. BASNER, R. PEHLKE and A. SACHDEV, Metall. Mater. Trans. A 31 (2000) 57.Received 23 Apriland accepted 12 August 2003材料科学期刊39 (2004)463-468Mg-9Al-1Zn合金的半固体挤压成型凯泽韦斯基开发工程师,赫斯基注塑系统有限公司,博尔顿,安大略,L7E 5 S5,加拿大这种结合半固体加工、挤出和注塑综合作用的新技术，应用在生产制造网状的镁合金部件。
01Effect of semi-solid processing on microstructure and mechanical properties of 5052 aluminum alloy
Effect of semi-solid processing on microstructure and mechanical properties of5052 aluminum alloyLÜ Shu-lin(吕书林), WU Shu-sen (吴树森), ZHU Ze-ming(朱泽明), AN Ping(安萍),MAO You-wu(毛有武)State Key Lab of Materials Processing and Die & Mould Technology,Huazhong University of Science and Technology, Wuhan 430074, ChinaReceived 13 May 2010; accepted 25 June 2010Abstract: The microstructure and mechanical properties of rheocasted 5052 aluminum alloy were investigated. The semi-solid slurry of this alloy was prepared by ultrasonic vibration (USV) process and then shaped by gravity casting (GC) and high press diecasting (HPDC). The experimental results indicate that fine and globular primary α(Al) particles are distributed uniformly in the rheocasting samples. The tensile strength and elongation of the rheo-GC sample are 191 MPa and 7.5%, respectively. Compared with the conventional GC samples, they increase by 22.4% and 82.9% respectively. The tensile strength and elongation of the rheo-HPDC samples reach 225 MPa and 8.6%, respectively, and they are 14.8% and 75.5% higher than those of the conventional HPDC samples, respectively. It is also found that the ductile fracture mode prevails in the rheocasting samples.Key words: ultrasonic vibration;rheocasting; 5052 aluminum alloy; mechanical properties; fracture surface1 IntroductionA lot of researches have been focused on the semi-solid metal (SSM) processing since it was invented in 1970s at Massachusetts Institute of Technology. Components produced by this unique technique have various advantages compared with the conventional liquid casting, such as low porosity, heat treatability and superior mechanical properties[1]. The rheoforming process becomes popular in recent years, because it possesses several advantages over thixoforming process, including increasing mould life, low cost, enhancing casting precision and qualities. Rheocasting involves stirring the melt during solidification to produce a non-dendritic semi-solid slurry, then injecting the slurry directly into a mould or die to give a final product[2]. The ultrasonic vibration (USV) for making semi-solid slurry is a relatively new method in rheocasting process. This novel technology requires less expensive equipment for production and it is easy to be introduced into the melt.The alloys commonly used for SSM processing are certain cast Al alloys, and some wrought Al alloys of 2000 or 7000 series[3-6]. However, little study has been conducted to process 5000 series Al alloys by SSM processing[7-9]. The 5000 alloys have been found a large variety of applications including architectural, household appliances, marine craft, and automotive structures due to its excellent combination of weldability and corrosion resistance[10]. They are generally classified as a non-heat treatable aluminum alloys, and their strength are not high enough to be used as structural components[11]. Therefore, it is necessary to adopt proper processing to obtain higher mechanical properties for these series alloys.The purpose of this work is to investigate the effect of USV rheocasting processing on microstructure and mechanical properties of 5052 Al alloy. The microstructure and tensile properties of rheocasting specimens of this alloy are compared with their liquid casting counterparts.2 ExperimentalThe installation of USV in Ref.[12] was employed for preparing semi-solid slurry of 5052 alloy in this experiment. In this work, the power of the ultrasonic generator was 1.2 kW, and the vibrating frequency was 20 kHz. The rest-work ratio R t of vibration was defined as the ratio of interval resting time t r to ultrasonic time t w in an USV cycle, i.e. R t=t r/t w. t w=1 s and t r=1 s were selected in this work.Foundation item: Project(2007AA03Z557) supported by the National High-Tech Research and Development Program of China; Project (50775086) supported by the National Natural Science Foundation of China.Corresponding author: WU Shu-sen; Tel: +86-27-87556262; E-mail: ssw636@Trans. Nonferrous Met. Soc. China 20(2010) s758-s762LÜ Shu-lin, et al/Trans. Nonferrous Met. Soc. China 20(2010) s758−s762 s759The chemical compositions of the 5052 Al alloys were 2.54% Mg, 0.49% Si, 0.33% Cr, 0.22% Fe, 0.19% Zn, 0.11% Cu, and 0.1%Mn (mass fraction, the same in the following). The raw materials were Al-25.8%Si, Al-5%Cr and Al-10%Mn master alloys, commercial purity Al (99.7%), pure Cu (99.99%), pure Mg (99.9%), and pure Fe (99.9%). The liquidus and the solidus temperature of this alloy were about 649 and 607 ˚C, respectively.The alloy was melted in a resistance furnace at 720-750 ℃, and then degassed for 10 min with argon gas through a graphite lance. The melt was cooled to a pouring temperature of 670-680 ˚C after degassing. The metal cup was preheated to 550 ˚C simultaneously. Subsequently, about 500 g of melt was poured into the metal cup, followed by application of USV. The vibration head was immersed into the melt 15 to 20 mm from the surface. In order to prevent the melt from oxidation, argon gas was introduced as the protective atmosphere during the slurry preparation process. After vibrated for a certain time at the temperature of about 645 ˚C, the liquid melt was converted into semi-solid slurry with certain solid fraction. The semisolid slurry was immediately poured in the iron permanent mould or the shot chamber of the HPDC machine, to produce standard tensile test samples with diameter of 8 mm and 6.4 mm, respectively. The moulds were preheated to about 200 ˚C in this experiment. For comparison, conventional liquid casting specimens, formed at the temperature of 730 ˚C, were also obtainedin the same moulds. The tensile tests were performed using a universal materials testing machine at a crosshead speed of 1.25 mm/min.Specimens for the metallographic examination were cut from the tensile test samples, polished and etched by a solution with 1% HF, 1.5% HCl and 2.5% HNO3. The microstructures and fracture surfaces of the tensile test specimens were examined by an Axiovert 200MAT optical microscope and a Quanta 200 environmental scanning electron microscope (SEM). Micrographs of the samples were analyzed using a quantitative metallographic analysis software[13]. The size of the primary α(Al) crystals was characterized by average particles diameter d and measured by Heyn’s linear intercept method, i.e. d was defined as: d=L T／N, whereL T is the overall length of measured lines used in the software for one metallograph and N is the number of particles passed through the measured lines.3 Results and discussion3.1 MicrostructuresThere are two different solidification stages taking place in USV rheocasting processing. Solidification in the process of making semi-solid slurry is referred to as first solidification, while solidification of the remaining liquid inside the mould is referred to as secondary solidification.Figs.1(a) and (b) show the representative micro- structures of 5052 alloy samples produced by rheo-GC and conventional GC, respectively. It is clear that spherical primary α(Al) particles formed in primary solidification were uniformly distributed throughout the entire cross-section of the rheocasted sample, with average diameter of 143 μm. Primary α′(Al) particles produced by the secondary solidification were rosette-like or spherical, and some were agglomerated together to form a bigger grain, and the average size of these particles was less than 50 μm. In the conventional casting sample, typical dendritic shape of primary α(Al) particle was observed. Microstructural examination indicates that few fine shrinkage pores could be observed in the sample, and the porosity fraction was about 0.64%.Fig.1 Microstructures of 5052 Al alloy samples produced under different processing conditions: (a) Rheo-GC; (b) Conventional GC Microstructures of the rheo-HPDC sample are shown in Figs.2(a) and (b). Unlike microstructure shown in Fig. 1(a), rheo-HPDC sample had a chilled zone with finer α′(Al) particles (with an average diameter of 23 μm) in the edge of the cross-section, and globular α particles (with an average diameter of 79 μm) mainly concentrated in the middle of the cross-section. In the conventional HPDC sample, dendritic morphology was observed, and porosity from entrapped gas could be found. The porosity fraction of the rheo-HPDC and conventional HPDC samples was 0.68% and 1.86%, respectively.LÜ Shu-lin, et al/Trans. Nonferrous Met. Soc. China 20(2010) s758−s762 s760Fig.2 Microstructures of 5052 Al alloy samples produced under different processing conditions: (a) In edge of rheo-HPDC sample; (b) In middle of rheo-HPDC sample; (c) Conventional HPDCIt was obvious that both of the primary α and primary α′particles in the rheo-HPDC sample were finer than those in rheo-GC sample, and this difference of microstructure was caused mainly by the different cooling rates of the slurry. The solidification rate plays an important role in the refinement of the microstructure and significantly affects the mechanical properties. A higher cooling rate in HPDC leads to the formation of a finer microstructureIn the process of semi-solid slurry making by USV, the formation of the non-dendrite microstructure could be mainly attributed to the ultrasonic cavitaion and acoustic streaming induced by USV. On the one hand, when USVis applied to a melt, a large number of tiny cavitation bubbles form. These bubbles expand and collapse instantaneously. The temperature of the cavitation bubble surface decreases during the expansion stage, and nucleation takes place on the cavitation bubble surface. The nuclei are dispersed throughout the melt by acoustic streaming. On the other hand, under the intensive agitation caused by ultrasonic cavitation and acoustic streaming, both the temperature and composition fields inthe melt are extremely uniform. Solute in front of the solidifying interface will be difficult to accumulate, andthe dendritic growth of crystals will be counteracted, thusnon-dendritic structure is formed[14].The remaining liquid in the semi-solid slurry will solidify in the die cavity without USV, which has a uniform temperature and composition throughout the liquid because of the intensive agitation beforehand. Nucleation will occur throughout the entire remaining liquid, and each nucleus would survive and contribute tothe final microstructure, resulting in very fine primary α′particles[15].3.2 Mechanical properties3.2.1 Tensile strengthTable 1 shows the mechanical properties of 5052 alloy under different processing conditions. Compared with the conventional casting counterpart, rheocasting samples had higher mechanical properties; especially, the tensile elongation was enhanced obviously. The tensile strength and elongation of the rheo-GC sample were 191 MPa and 7.5%, respectively, and compared with the conventional GC one, they were increased by 22.4% and 82.9%, respectively. The tensile strength and elongationof the rheo-HPDC sample were 225 MPa and 8.6%, respectively, and they were 14.8% and 75.5% higher than those of the conventional HPDC one, respectively.Table 1 Mechanical properties of 5052 alloy under different processing conditionsProcessing conditionYieldstrength/MPaUltimatestrength/MPaElongation/% Conventional GC 83 156 4.1Rheo-GC 87 191 7.5 Conventional HPDC158 196 4.9Rheo-HPDC 172 225 8.6It has been commonly accepted that the existence of large pores caused by entrapped gas and shrinkage pores reduces the mechanical properties of diecasting samples.The rheocasting sample has better mechanical propertiesLÜ Shu-lin, et al/Trans. Nonferrous Met. Soc. China 20(2010) s758−s762 s761than the conventional casting sample because less shrinkage and associated porosity are expected in the rheocastings as pouring occurs at a temperature below the liquidus. The globular primary solid structure in the mushy semisolid slurry would be more favourable to liquid penetration for feeding[16]. On the other hand, the fine and uniform microstructure throughout the entire sample would effectively reduce stress concentration at the grain boundary under an applied stress. Therefore, the globular shape would improve the tensile strength and ductility of the rheocasting sample.3.2.2 Fracture surfaceFigs.3(a) and (b) show the SEM micrographs of the tensile fracture for the samples produced by rheo-GC and conventional GC, respectively. For both samples, the tensile fracture paths tended to follow the primary α(Al) phase boundaries. But the shearing of primary α(Al) phase is observed in the conventional cast sample. The rheo-GC sample had relatively sharper and longer peaks along the fracture surface. This indicates that the mixed ductile-brittle fracture mode prevailed in the conventional cast sample, and the rheocasting sample had a ductile fracture mode.Fig.3 SEM micrographs of tensile fracture for samples produced under different processing conditions: (a) Rheo-GC; (b) Conventional GCFigs.4(a) and (b) show the SEM micrographs of the tensile fracture for the rheo-HPDC and conventional HPDC samples. In both samples, fractures appeared to initiate and propagate through the grain boundaries, and dimple-rupture morphology prevailed throughout the fracture surface. The dimples on the rheocast sample were smaller than those on the conventional casting sample.Fig.4 SEM micrographs of tensile fracture for samples produced under different processing conditions: (a) Rheo-HPDC; (b) Conventional HPDC4 Conclusions1) Non-dendritic microstructure of 5052 alloy can be obtained by USV rheocasting process, and globular primary α(Al) particles are uniformly distributed throughout the entire rheocasting samples. In addition, the diameters of α(Al) particles in rheo-GC and rheo-HPDC are about 140 and 80 μm, respectively.2) The mechanical properties of the rheocasting samples are considerably higher than the conventional casting ones; especially, the tensile elongation is enhanced obviously.3) In rheocasting samples, factures initiate and propagate through the grain boundaries and dimple-rupture morphology prevails throughout the tensile fracture surface.LÜ Shu-lin, et al/Trans. Nonferrous Met. Soc. China 20(2010) s758−s762 s762References[1]MARTINEZ R A, FLEMINGS M C. Evolution of particlemorphology in semisolid processing [J]. Metall Mater Trans A, 2005,36 (8): 2205-2210.[2]PARK C, KIM S, KWON Y, LEE Y, LEE J. Mechanical andcorrosion properties of rheocast and low-pressure cast A356-T6 alloy[J]. Materials Science and Engineering A, 2005, 391: 86-94.[3]GUO Hong-min, YANG Xiang-jie, ZHANG Meng. Microstructurecharacteristics and mechanical properties of rheoformed wrought aluminum alloy 2024[J]. Transactions of Nonferrous Metals Societyof China, 2008, 18: 555-561.[4]AZPILGAIN Z, HURTADO I, ORTUBAY R, LANDA I, ATXA J.Semisolid forging of 7000 series aluminum alloys [J]. Solid StatePhenomena, 2006, 116/117:758-761.[5]JANG D I, YOON Y O, KIM S K. Thixoextrusion for 7075 Alwrought alloy tube[J]. Solid State Phenomena, 2008, 141/142/143:267-270.[6]WANG Wei-wei, SONG Jian-li, LUO Shou-jing. Preparation oflarge-diameter semi-solid 7075 aluminum alloy billets[J]. Solid StatePhenomena, 2008, 141/142/143: 361-365.[7]DAI S L, DELPLANQUE J P, LAVERNIA E J. Grain growth inreactive spray depositied 5083 alloys [J]. Scripta Materialia, 1999,40(2):145-151.[8]SAKLAKOGLU N, SAKLAKOGLU I E, TANOGLU M, OZTAS O,CUBUKCUOGLU O. Mechanical properties and microstructural evaluation of AA5013 aluminum alloy treated in the semi-solid stateby SIMA process[J]. Journal of Materials Processing Technology2004,148:103-107.[9]PIAO L, MOTEGI T. Solidified structure of a 5052 aluminum alloyobtained by semisolid casting[J].Journal of Japan Institute of LightMetals, 2005, 55(2): 86-90.[10]KAIBYSHEV R, MUSIN F, LESUER D R, NIEH T G. Superplasticbehavior of an Al-Mg alloy at elevated temperatures[J]. MaterialsScience and Engineering A, 2003, 342:169-177.[11]LIN Shuang-ping, NIE Zuo-ren, HUANG Hui, LI Bo-long.Annealing behavior of a modified 5083 aluminum alloy[J].Materialsand Design, 2010, 31:1607-1612.[12]WU Shu-sen, ZHAO Jun-wen, ZHANG Li-ping, AN Ping, MAOYou-wu. Development of non-dendritic microstructure of aluminumalloy in semi-solid state under ultrasonic vibration[J]. Solid State Phenomena, 2008，141/142/143:451-456.[13]WU Shu-sen, XIE Li-zhi, ZHAO Jun-wen, NAKAE H. Formation ofnon-dendritic microstructure of semi-solid aluminum alloy under vibration[J]. Scripta Materialia, 2008, 58: 556-559.[14]WU Shu-sen, WU Xue-ping, XIAO Ze-hui. A model of growthmorphology for semi-solid metals [J]. Acta Materialia, 2004, 52: 3519-3524.[15]YANG Xiang-jie, GUO Hong-min. Microstructure evolution andmechanical properties of rheoformed YL112 aluminiun alloys[J].Solid State Phenomena, 2008, 141/142/143: 163-168.[16]BRABAZON D, BROWNE D J, CARR A J. Mechanical stir castingof aluminium alloys from the mushy state: process, microstructureand mechanical properties[J]. Materials Science and Engineering A,2002, 326: 370-381.(Edited by YANG You-ping)。
南昌大学科技成果——AZ61镁合金半固态坯料的制备方法
南昌大学科技成果——AZ61镁合金半固态坯料的制备方法背景技术
近年来，随着镁合金生产技术的发展，镁合金的应用获得人们的极大关注。
过去镁合金主要应用于冶金行业，而在结构材料方面的应用仅仅局限于航天航空方面少量的要求。
现在，随着半固态金属加工技术广泛应用和推广，以及汽车制造、航空航天、通讯、光学仪器和计算机领域的飞速发展，镁合金作为一种理想的工业材料，被认为是二十一世纪最具有应用和开发前景的商用轻质材料。
技术原理
本项目是由预变形、等温热处理工艺流程组成，其特征是预变形加热温度为300℃-400℃，变形量控制在5%-42%，然后采取阻燃措施，进行等温热处理，加热温度为570℃-610℃，保温时间0min-40min。
本发明所述的制备方法制备出的AZ61镁合金半固态坯料，其球状晶粒组织细小，园整度较好，其晶粒组织的等积圆直径大部分位于90μm-150μm之间，且晶粒组织的圆度主要在1.2-1.8范围，完全满足其后续触变成形的要求，而且工艺简单、安全可靠，无三废污染。
知识产权情况
发明专利1项，ZL200510063947.8。
铝、镁合金半固态浆料的制备与流变成形新工艺
铝、镁合金半固态浆料的制备与流变成形新工艺滕海涛;熊柏青;张永安;李廷举;张小立;谢水生【期刊名称】《中国有色金属学报》【年(卷),期】2012(022)004【摘要】介绍了阻尼冷却管法制备A356铝合金半固态浆料工艺的实验装置及其原理,并进行不同浇注温度的系列实验.结果表明:由于阻尼冷凝管的冷却和搅拌作用,熔体浇注温度越低,在两相温度区间内生成的游离晶核就越多,制备得到半固态铸件的晶粒尺寸就越小,且球化程度越高.在此实验原理及结果分析的基础上,设计真空吸铸-阻尼冷却装置,近液相线温度的AZ91D镁合金液在冷却、剪切的作用下,由液态转变为半固态浆料,然后进入模具完成充型,实现镁合金半固态浆料的制备与铸件流变成形一体化；半固态镁合金熔体具有触变性及更高的黏度,以平稳、层流的充型方式完成充型,能够有效地改善成形件的质量.【总页数】6页(P1019-1024)【作者】滕海涛;熊柏青;张永安;李廷举;张小立;谢水生【作者单位】北京有色金属研究总院有色金属材料制备加工国家重点实验室有研亿金新材料股份有限公司,北京100088;大连理工大学三束材料改性教育部重点实验室,大连116024;北京有色金属研究总院有色金属材料制备加工国家重点实验室有研亿金新材料股份有限公司,北京100088;北京有色金属研究总院有色金属材料制备加工国家重点实验室有研亿金新材料股份有限公司,北京100088;大连理工大学三束材料改性教育部重点实验室,大连116024;大连理工大学三束材料改性教育部重点实验室,大连116024;北京有色金属研究总院有色金属材料制备加工国家重点实验室有研亿金新材料股份有限公司,北京100088【正文语种】中文【中图分类】TG146.2;TG249【相关文献】1.强制对流流变成形制备7075铝合金半固态浆料及其数值模拟 [J], 周冰;康永林;朱国明;郜俊震;祁明凡;张欢欢2.强制对流流变成形制备7075铝合金半固态浆料及其数值模拟 [J], 周冰;康永林;朱国明;郜俊震;祁明凡;张欢欢;3.新工艺制备A356铝合金半固态浆料及其流变压铸组织和性能 [J], 祁明凡;康永林;徐玉昭;李静媛;刘爱森4.新工艺制备A356铝合金半固态浆料及其流变压铸组织和性能 [J], 祁明凡;康永林;徐玉昭;李静媛;刘爱森5.镁合金半固态材料成形与浆料制备工艺研究 [J], 孙建林;康永林;肖邦国;安林因版权原因，仅展示原文概要，查看原文内容请购买。
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Ｓｃｉｅｎｃｅ Ｐｒｅｓｓ
Ａｖａｉｌａｂｌｅ ｏｎｌｉｎｅ ａｔ ｗｗｗ．ｓｃｉｅｎｃｅｄｉｒｅｃｔ．ｃｏｒｎ ●Ｃ●ＥｉｌｌＣＩＥ Ｄ●ＲＥＣＴ ０
Ｔｒａｎｓ．Ｎｏｎｆｅｒｒｏｕｓ Ｍｅｔ．Ｓｏｃ．Ｃｈｉｎａ １ ７（２００７）４６—５０ Ｔｒａｎｓａｃｔｉｏｎｓ ｏｆ
Ｎｏｎｆｅｒｒｏｕｓ Ｍｅｔａｌｓ Ｓｏｃｉｅｔｙ ｏｆ Ｃｈｉｎａ
ｗｗｗ．ＣＳＵ．ｅｄｕ．ｃｎ／ｙｓｘｂ／
Ｐｒｅｐａｒａｔｉｏｎ ｏｆ ｓｅｍｉ—ｓｏｌｉｄ ｂｉｌｌｅｔ ｏｆ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ ａｎｄ ｉｔｓ ｔｈｉｘｏｆｏｒｍｉｎｇ ＪＩＡＮＧ Ｊｕ． （姜巨福），ＬＵＯ Ｓｈｏｕ－ｊｉｎｇ（￣守靖） Ｓｃｈｏｏｌ ｏｆ Ｍａｔｅｒｉａｌｓ Ｓｃｉｅｎｃｅ ａｎｄ Ｅｎｇｉｎｅｅｒｉｎｇ，Ｈａｒｂｉｎ Ｉｎｓｔｉｔｕｔｅ ｏｆ Ｔｅｃｈｎｏｌｏｇｙ，Ｈａｒｂｉｎ １ ５０００ １，Ｃｈｉｎａ Ｒｅｃｅｉｖｅｄ ２７ Ｍａｒｃｈ ２００６；ａｃｃｅｐｔｅｄ １ ０ Ｊｕｎｅ ２００６
Ａｂｓｔｒａｃｔ：Ｐｒｅｐａｒａｔｉｏｎ ｏｆ ｓｅｍｉ．ｓｏｌｉｄ ｂｉｌｌｅｔ ｏｆ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ ａｎｄ ｔｈｉｘｏｆｏｒｍｉｎｇ ｗａｓ ｉｎｖｅｓｔｉｇａｔｅｄ ｂｙ ａｐｐｌｙｉｎｇ ｅｑｕａｌ ｃｈａｎｎｅｌ ａｎｇｕｌａｒ ｅｘｔｒｕｓｉｏｎ ｔｏ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ．Ｔｈｅ ｒｅｓｕｌｔｓ ｓｈｏｗ ｔｈａｔ ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ ｏｆ ＡＺ９ ｌ Ｄ ａｌｌｏｙ ａｔ ｒｏｏｍ ｔｅｍｐｅｒａｔｕｒｅ，ｓｕｃｈ ａｓ ｙｉｅｌｄ ｓｔｒｅｎｇｔｈ（ＹＳ）。ｕｌｔｉｍａｔｅ ｔｅｎｓｉｌｅ ｓｔｒｅｎｇｔｈ（ＵＴＳ）ａｎｄ ｅｌｏｎｇａｔｉｏｎ，ａｒｅ ｅｎｈａｎｃｅｄ ｇｒｅａｔｌｙ ｂｙ ｆｏｕｒ－ｐａｓｓ ｅｑｕａｌ ｃｈａｎｎｅｌ ａｎｇｕｌａｒ ｅｘｔｒｕｓｉｏｎ（ＥＣＡＥ）ａｔ ５７３ Ｋ ａｎｄ ｍｉｃｒｏｓｔｒｕｃｔｕｒｅ ｏｆ ＡＺ９ ｌ Ｄ ａｌｌｏｙ ｉｓ ｒｅｆｉｎｅｄ ｔｏ ｔｈｅ ａｖｅｒａｇｅ ｇｒａｉｎ ｓｉｚｅ ｏｆ ２０ ｇｍ．Ｔｈｒｏｕｇｈ ｕｓｉｎｇ ＥＣＡＥ ａｓ ｓｔｒａｉｎ ｉｎｄｕｃｅｄ ｓｔｅｐ ｉｎ ＳＩＭＡ ａｎｄ ｃｏｍｐｌｅｔｉｎｇ ｍｅｌｔ ａｃｔｉｖａｔｅｄ ｓｔｅｐ ｂｙ ｓｅｍｉ－ｓｏｌｉｄ ｉｓｏｔｈｅｒｍａｌ ｔｒｅａｔｍｅｎｔ，ｓｅｍｉ－ｓｏｌｉｄ ｂｉｌｌｅｔ ｗｉｔｈ ｆｉｎｅ ｓｐｈｅｒｏｉｄａｌ ｇｒａｉｎｓ ｏｆ ２５啪ｃａｎ ｂｅ ｐｒｅｐａｒｅｄ ｓｕｃｃｅｓｓｆｕｌｌｙ．Ｃｏｍｐａｒｅｄ ｗｉｔｈ ｃｏｍｍｏｎ ＳＩＭＡ，ｔｈｉｘｏｆｏｒｍｅｄ ｓａｔｅｌｌｉｔｅ ａｎｇｌｅ ｆｌａｍｅ ｃｏｍｐｏｎｅｎｔｓ ｕｓｉｎｇ ｓｅｍｉ．ｓｏｌｉｄ ｂｉｌｌｅｔ ｐｒｅｐａｒｅｄ ｂｙ ｎｅｗ ＳＩＭＡ ｈａｖｅ ｈｉｇｈｅｒ ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ ａｔ ｒｏｏｍ ｔｅｍｐｅｒａｔｕｒｅ ａｎｄ ｈｉｇｈ
ｔｅｍｐｅｒａｔｕｒｅ ｏｆ ３７３ Ｋ．
Ｋｅｙ ｗｏｒｄｓ：ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ；ｔｈｉｘｏｆｏｒｍｉｎｇ；ｅｑｕａｌ ｃｈａｎｎｅｌ ａｎｇｕｌａｒ ｅｘｔｒｕｓｉｏｎ（ＥＣＡＥ）；ｓｅｍｉ－ｓｏｌｉｄ ｂｉｌｌｅｔ；ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ
ｌ Ｉｎｔｒ０ｄｕｃｔｉ０ｎ Ｅｑｕａｌ ｃｈａｎｎｅｌ ａｎｇｕｌａｒ ｅｘｔｒｕｓｉｏｎ（ＥＣＡＥ）ｉｓ ａ ｐｒｏｍｉｓｉｎｇ ｔｅｃｈｎｉｑｕｅ ｆｏｒ ｏｂｔａｉｎｉｎｇ ｕｌｔｒａ－ｆｉｎｅ ｇｒａｉｎｅｄ ｂｕｌｋ ｍａｔｅｒｉａｌｓ ｗｉｔｈ ｉｎｃｒｅａｓｅｄ ｓｔｒｅｎｇｔｈ ａｎｄ ｄｕｃｔｉｌｉｔｙ ｔｈｒｏｕｇｈ ｓｅｖｅｒｅ ｐｌａｓｔｉｃ ｄｅｆｏｒｍａｔｉｏｎ［１—２］．Ｄｕｒｉｎｇ ＥＣＡＥ， ａ ｂｉｌｌｅｔ ｉｓ ｐｒｅｓｓｅｄ ｔｈｒｏｕｇｈ ａ ｄｉｅ ｔｈａｔ ｃｏｎｓｉｓｔｓ ｏｆ ｔｗｏ ｃｈａｎｎｅｌｓ ｗｉｔｈ ｅｑｕａｌ ｃｒｏｓｓ．ｓｅｃｔｉｏｎ。ｉｎｔｅｒｓｅｃｔｉｎｇ ａｔ ａｎ ａｎｇｌｅ∞．Ｓｉｎｃｅ ｔｈｅ ｃｒｏｓｓ．ｓｅｃｔｉｏｎａｌ ｓｈａｐｅ ｏｆ ｂｉｌｌｅｔ ｒｅｍａｉｎｓ ｎｅａｒｌｙ ｔｈｅ ｓａｍｅ，ＥＣＡＥ ｐｒｏｃｅｓｓｉｎｇ ｃａｎ ｂｅ ｒｅｐｅａｔｅｄ ｆｏｒ ｍａｎｙ ｐａｓｓｅｓ ｔｏ ｃｏｎｔｒｏｌ ｔｈｅ ｍｉｃｒｏｓｔｒｕｃｔｕｒｅ ａｎｄ ｐｒｏｐｅｒｔｉｅｓ ｏｆ ｔｈｅ ｍａｔｅｒｉａｌ［３—６］．Ａｓ ａ ｔｙｐｉｃａｌ Ｍｇ－Ａｌ－Ｚｎ ａｌｌｏｙ，ＡＺ９ ｌ Ｄ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ ｈａｓ ｒｅｃｅｉｖｅｄ ｍｏｒｅ ａｎｄ ｍｏｒｅ ａｔｔｅｎｔｉｏｎ ｄｕｅ ｔｏ ｔｈｅ ｈｉｇｈ ｓｐｅｃｉｆｉｃ ｓｔｒｅｎｇｔｈ ‘ａｎｄ ｌｏｗ ｃｏｓｔ．Ｔｈｉｘｏｆｏｒｍｉｎｇ ｉｓ ａｎ ａｌｔｅｒｎａｔｉｖｅ ｆｏｒ ｐｒｏｄｕｃｉｎｇ ｃｏｍｐｏｎｅｎｔｓ Ｏｆ ＡＺ９１Ｄ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ． Ｔｈｉｘｏｆｏｒｍｉｎｇ ｉｓ ａ ｔｙｐｉｃａｌ ｓｅｍｉ—ｓｏｌｉｄ ｐｒｏｃｅｓｓｉｎｇ，ｗｈｉｃｈ ｉｓ ｏｎｅ ｏｆｔｈｅ ｂｅｓｔ ｍｅｔｈｏｄｓ ｆｏｒ ｆｏｒｍｉｎｇ ＡＺ９ｌＤ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ ｃｏｍｐｏｎｅｎｔｓ ｂｅｃａｕｓｅ ｏｆ ｉｔｓ ｌｏｗ ｒｅｓｉｓｔａｎｃｅ Ｏｆ ｄｅｆｏｒｍａｔｉｏｎ ｃｏｍｐａｒｅｄ ｗｉｔｈ ｓｏｌｉｄ ａｌｌｏｙ ｆｏｒｇｉｎｇ ａｎｄ ｈｉｇｈ ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ ｃｏｍｐａｒｅｄ ｗｉｔｈ ｌｉｑｕｉｄ ａｌｌｏｙ ｄｉｅ ｃａｓｔｉｎｇ［７—８１．ＨＡＧＡ ａｎｄ ＫＡＰＲＡＮＯＳ［９—１ ０】ｓｔｕｄｉｅｄ ｔｈｉｘｏｆｏｒｍｉｎｇ ｏｆ ｌａｍｉｎａｔｅ ｍａｄｅ ｆｒｏｍ ｓｅｍｉ－ｓｏｌｉｄ ｃａｓｔ ｓｔｒｉｐｓ ａｎｄ ｂｉｌｌｅｔｌｅｓｓ ｓｉｍｐｌｅ ｔｈｉｘｏｆｏｒｍｉｎｇ ｐｒｏｃｅｓｓ． ＣＨＡＹ０ＮＧ ｅｔ ａｌ『ｌ１ １ ｒｅｐｏｒｔｅｄ ｔｈｉｘｏｆｏｒｍｉｎｇ ｏｆ ７０７５ ａｌｕｍｉｎｉｕｍ ａｌｌｏｙ．ＨＩＲＴ ｅｔ ａｌ［１ ２１ ｓｔｕｄｉｅｄ ｎｅｔ ｓｈａｐｅ ｃｏｍｐｏｎｅｎｔｓ ｐｒｏｄｕｃｅｄ ｂｙ ｔｈｉｘｏｆｏｒｍｉｎｇ． Ｆｏｒ ｔｈｉｘｏｆｏｒｍｉｎｇ ｏｆ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ，ｔｈｅ ｍｏｓｔ ｉｍｐｏｒｔａｎｔ ｉｓ ｔｏ ｐｒｅｐａｒｅ ｓｅｍｉ－ｓｏｌｉｄ ｂｉｌｌｅｔｓ ｏｆ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ，ｔｈａｔ ｉｓ，ｔｏ ｏｂｔａｉｎ ａｌｌｏｙ ｗｉｔｈ ｎｏｎ－ｄｅｎｄｒｉｔｉｃ，ｏｒ ｓｐｈｅｒｏｉｄａｌ ｍｉｃｒｏｓｔｒｕｃｔｕｒｅ．Ｃｏｍｐａｒｅｄ ｗｉｔｈ ｏｔｈｅｒ ｍｅｔｈｏｄｓ ｆ０ｒ ｐｒｅｐａｒｉｎｇ ｓｅｍｉ．ｓｏｌｉｄ ｂｉｌｌｅｔｓ．ｓｔｒａｉｎ ｉｎｄｕｃｅｄ ｍｅｌｔ ａｃｔｉｖａｔｉｏｎ（ＳＩＭＡ）ｉｓ ａ ｇｏｏｄ ｍｅｔｈｏｄ ｆｏｒ ｐｒｅｐａｒｉｎｇ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ ｓｅｍｉ．ｓｏｌｉｄ ｂｉｌｌｅｔｓ ｗｉｔｈ ｈｉｇｈ ｓｏｌｉｄ ｆｒａｃｔｉｏｎ．Ｉｎ ｇｅｎｅｒａｌ，ｔｈｅｒｅ ａｒｅ ｔｗｏ ｉｍｐｏｒｔａｎｔ ｓｔｅｐｓ ｉｎ ＳＩＭＡ．ｓｔｒａｉｎ ｉｎｄｕｃｅｄ ｓｔｅｐ ａｎｄ ｍｅｌｔ ａｃｔｉｖａｔｅｄ ｓｔｅｐ．Ｉｎ ｓｔｒａｉｎ ｉｎｄｕｃｅｄ ｓｔｅｐ．ｃａｓｔ ｍａｔｅｒｉａｌ ｉｓ ｄｅｆｏｒｍｅｄ ｂｙ ｕｐｓｅｔｔｉｎｇ．Ｉｎ ｍｅｌｔ ａｃｔｉｖａｔｅｄ ｓｔｅｐ，ｔｈｉｓ ｕｐｓｅｔｔｅｄ ｍａｔｅｒｉａｌ ｉｓ ｓｅｍｉ．ｓｏｌｉｄ ｉｓｏｔｈｅｒｍａｌ ｔｒｅａｔｅｄ，ｔｈａｔ ｉｓ，ｕｐｓｅｔｔｅｄ ｍａｔｅｒｉａｌｉｓ ｈｅａｔｅｄ ｔｏ ｓｅｍｉ．ｓｏｌｉｄ ｔｅｍｐｅｒａｔｕｒｅ ａｎｄ ｔｈｅ ｌｏｗ ｍｅｌｔｉｎｇ ｐｏｉｎｔ ｐｈａｓｅｓ ａｒｅ ｐａｒｔｉａｌｌｙ ｍｅｌｔ ｉｎ ｏｒｄｅｒ ｔｏ ｏｂｔａｉｎ ｓｅｍｉ．ｓｏｌｉｄ ｂｉｌｌｅｔ ｗｉｔｈ ｓｐｈｅｒｏｉｄａｌ ｍｉｃｒｏｓｔｒｕｃｔｕｒｅ． Ｈｏｗｅｖｅｒ，ｆｏｒ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ，ｃｏａｒｓｅ ａｎｄ ｎｏｎ－ｇｌｏｂｕｌａｒ ｇｒａｉｎｓ ｏｆｔｅｎ ｅｘｉｓｔ ｉｎ ｔｈｅ ｍｉｃｒｏｓｔｒｕｃｔｕｒｅ ｏｆ ｓｅｍｉ．ｓｏｌｉｄ ｂｉｌｌｅｔｓ ｐｒｅｐａｒｅｄ ｂｙ ｃｏｍｍｏｎ ＳＩＭＡ ｍｅｔｈｏｄ，ｔｈａｔ ｉｓ， ｓｅｍｉ．ｓｏｌｉｄ ｉｓｏｔｈｅｒｍａｌ ｔｒｅａｔｍｅｎｔ ｏｆ ｕｐｓｅｔｔｅｄ ｃａｓｔ ｍａｔｅｒｉａｌｓ．Ｔｈｉｓ ｉｓ ｄｕｅ ｔｏ ｄｉｆｎｃｕｌｔｉｅｓ ｏｆ ｏｂｔａｉｎｉｎｇ ｓｅｖｅｒｅ ｐｌａｓｔｉｃ ｄｅｆｏｒｍａｔｉｏｎ ｗｉｔｈ ｒｅｓｐｅｃｔ ｔｏ ｍａｇｎｅｓｉｕｍ ａｌｌｏｙ ｉｎ ＳＩＭＡ．ｗｈｉｃｈ ｒｅｓｕｌｔｓ ｆｒｏｍ ｔｈｅ ｎａｔｕｒｅ ｏｆ ｃｒｙｓｔａｌ ｓｔｒｕｃｔｕｒｅ． Ｔｈｅｒｅｆｏｒｅ．ａ ｎｅｗ ｓｔｒａｉｎ ｉｎｄｕｃｅｄ ｔｅｃｈｎｉｑｕｅ ｆｏｒ ＳＩＭＡ ｎｅｅｄｓ ｔｏ ｂｅ ａｄｏｐｔｅｄ ｔｏ ｏｂｔａｉｎ ｓｅｖｅｒｅ ｐｌａｓｔｉｃ ｄｅｆｏｒｍａｔｉｏｎ，
Ｆｏｕｎｄａｔｉｏｎ ｉｔｅｍ：ａｒｏｊｅｃｔ（５０６０５０１５）ｓｕｐｐｏｒｔｅｄ ｂｙ ｔｈｅ Ｎａｔｉｏｎａｌ Ｎａｔｕｒａｌ Ｓｃｉｅｎｃｅ Ｆｏｕｎｄａｔｉｏｎ ｏｆＣｈｉｎａ Ｃｏｒｒｅｓｐｏｎｄｉｎｇａｕｔｈｏｒ：ＪＩＡＮＧ Ｊｕ－ｆｕ；Ｔｅｌ：＋８６－４５Ｉ一８６４１５４６４；Ｅ－ｍａｉｌ：ｊｉａｎｇｊｕｆｕ＠ｈｉｔ．ｅｄｕ．ｃｒＩ
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