镁合金---试验翻译

熔剂的反应机理镁合金的熔点不高,热容量较小,在空气中加热时,氧化快,在过热时,易燃烧,在熔融状态下无熔剂保护时,则可猛烈燃烧。

因此镁合金在熔炼过程中必须始终在熔剂和保护气氛下进行。

同时，在转移镁合金及浇注成型过程中，各种工具也需要进行洗涤和保护，这些都要涉及到熔剂。

其实熔剂的基本作用是在熔体表面形成—化学抑制层来防止熔体氧化，并去除熔体中的固态和气态的非金属夹杂物。

要选择比镁氧化亲和性更强的物质来做镁合金熔剂，最好是用碱金属和碱土金属的氯化盐和氟化盐。

镁合金熔剂主要由MgCl2、KCl、CaF2、BaCl2等氯盐及氟盐的混合物组成，它们按一定比例混匀，使熔剂的熔点、密度、黏度及表面性能均较好的满足使用要求。

其中MgCl2是起主要作用的成分，对镁熔体具有良好的覆盖作用及一定精炼能力。

主要表现在以下几个方面：1、MgCl2熔点为981K ，易与其他盐混合形成低熔点盐类混合物，迅速地铺展成一层连续、致密的熔剂层。

2、MgCl2能很好地湿润熔体表面的MgO，并将其包覆后转移到熔剂中去，消除了由MgO 所产生的绝热作用，使Mg在氧化中产生的热量能较快地通过熔剂层散出。

避免镁合金熔体表面温度急剧上升。

3、MgCl2有很大的吸湿性，裸露在空气中很快受潮，与空气中的氧及水气反应生成HCl、Cl2、H2等。

其反应如下：2MgCl2+O2=2MgO+2Cl2 MgCl2+H2O=MgO+2HCl反应生成的Cl2和HCl能迅速的和Mg反应生成MgCl22HCl+Mg= MgCl2+H2↑Mg+Cl2= MgCl2这样HCl、Cl2、H2等保护气氛及MgCl2薄层覆盖均能有效地阻止Mg与O和H2O的作用，防止氧化，抑制燃烧。

4、MgCl2具有化学造渣的作用，形成的产物MgCl2·5 MgO能从熔体中沉淀出来。

MgCl2 + 5 MgO= MgCl2·5 MgO往MgCl2中加入KCl、NaCl能够显著降低熔剂的熔点，密度，黏度，提高熔剂稳定性。

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.。

ITEMS NEEDED FOR AZ91 CASTING TRIALAZ91压铸试验所需要准备工作The following items will need to be prepared in advance of a AZ91 ingot casting trial using the new draft procedure supplied by Applied Magnesium.在根据应用镁公司提供的新的AZ91D合金生产指南草稿进行AZ91D合金锭试验铸造之前，需要提前做好以下的准备事项。

1.One in use and one spare electrically driven casting pump currently located in thestatic cast workshop. Make certain that the following items are checked and/orincluded;在手浇铸车间目前应放置两个电驱动的浇铸泵，一个现在使用，一个作为备用。

确定下列物品已经过检验或者包括在内。

Pump can be hung from crane hook.浇铸泵应可以用吊钩吊起来。

Pump is operational.浇铸泵应可以使用。

Variable speed control to regulate pump speed is included at the alloy workshop.在合金车间，浇铸泵的速度应该有变速控制装置。

Pump to transfer pipe flange is compatible design so that the twoassemblies will fit together.浇铸泵连接移液管的接口装置应该吻合，这两样器具就可以匹配的很好。

Pump is clean and free of debris. MAKE CERTAIN THAT PUMP IS PRE-HEATED IN MOLTEN FLUX JUST PRIOR TO USING INMOLTEN MAGNESIUM浇铸泵应该很干净，无杂质粘连。

Comparison of Microstructure and Mechanical Propertiesof AZ91D Alloy Formed by Rheomoldingand High-Pressure Die Casting()The microstructure and mechanical properties of AZ91D alloy thin-wall parts produced by the 组织 AZ91D镁合金薄壁零件的机械性能产生rheomolding(RM) process were investigated and compared with the same alloy formed by conventional研究传统的high pressure die casting (HPDC). The results indicate that the RM process is able to get such AZ91D 高压压铸parts in which a1-Mg with average size of 27.36 l m are spherical and uniformly distributed in the球形均匀分布matrix, and the matrix is a mixture of numerous fine a2-Mg and intermetallic b-Mg17Al12. High mechanical 矩阵机械properties including ultimate tensile strength (UTS) of 270 MPa, yield strength (YS) of 169 MPa, 极限抗拉强度屈服强度elongation of 7.1%,and Vickers hardness of 102 are obtained in parts formed by RM due to the fine维氏硬度and uniform microstructure and less porosities. Compared with HPDC, the UTS, YS, elongation, and 组织均匀，气孔少压铸hardness of RM AZ91D are increased by 14.4, 9.7, 86.8, and 21.4%, respectively. The solidified grains凝固晶粒in RM AZ91D alloy show a smaller aluminum gradient than that in HPDC. This indicates that the较小铝梯度solidification of the RM AZ91D is closer to equilibrium..1. IntroductionMg-alloys, with a number of desirable properties including light weight, high specific strength理想性能比强度and specific stiffness,excellent damping property and well castability, are thus very 比刚度优良阻尼性能铸造性能attractive for the applications in 3C (computers, communications,and consumer electronics) andautomotive industries (Ref1-4). Over the past decades, with the rapid expansion of Mg-alloyapplications, large-scale thin-wall parts have been developed and implemented by taking full大型薄壁零件实施advantage of high-pressure die casting (HPDC) (Ref 5). However, HPDC partshave high-gas porosity levels, due primarily to the entrapment of air or gas in the melt during 高瓦斯孔隙度空气滞留the high-speed filling of turbulent molten metal into the cavity. The porosities can腔孔隙度severely degrade mechanical properties by acting as local stress concentrators. They also lead 降解浓缩机to problems during heat treatment or welding, where heating causes the expansion of gas in pores,热处理或焊接扩展and result in bubbling and dimensional changes (Ref 6). Also,repositories may have an adverse尺寸变化effect on the corrosion resistance of Mg-alloys (Ref 7).耐腐蚀性In order to solve these problems and meet demands of future applications, alternative castingprocess is developing. Semisolid metal (SSM) processing is a promising manufacturing半固态金属制造route that is capable of producing castings with a high level of quality. SSM processing involves 路线casting a semisolid slurry that exhibits non-turbulent or thixotropic flow behavior (Ref 8-10).铸造半固态浆料展品非湍流触变流动行为Semisolid cast alloys offer several advantages over their HPDC counterparts. For example, the 半固态铸造fraction of pores is lower owing to the laminar mold-filling process that results in less孔隙分数层流充型过程entrapped air (Ref 11, 12). SSM techniques are divided into two categories: thixo (thixomolding 截留空气类别触变(TM) and thixocasting (TC)) and rheo (rheomolding (RM) and rheocasting (RC)) processes.触变铸造流变流变铸造However, TC and RC are difficult to form thin-wall parts for the poor controllability of the melt temperature in chamber.Presently, the only commercially available SSM technology forthin-wall Mg-alloy parts is TM. Though TM has made a great progress and produced the parts withbetter strength and ductility, there still exist many shortcomings, such as poor wear resistance耐磨性差and short service life of the screw and the cylinder liner which are key components of thethixomolder (Ref 13).Moreover, using Mg-alloy particles as raw materials directly results in原材料the increase of production cost (Ref 13, 14).生产成本2. Experiment Procedures2.1 MaterialsCommercial AZ91D alloy was used in this investigation, for which the reported solidus and liquidus temperatures are 468 and 598 C, respectively. The chemical composition of theAZ91D alloy is 9.45% Al, 0.66% Zn, 0.20% Mn, 0.036% Si,0.005% Cu, 0.001% Ni, and Mg balance (by weight).2.2 The RM ProcessThe RM process is an innovative one-step SSM processing technique which, through the use创新一步半固态加工技术of LSP technology, can manufacture near-net shape parts with high integrity directly from liquid制造近净成形件完整性液体alloy without turbulence or gas entrapment.Figure 1 shows the schematic of the NISSEI FMg220-16HM 湍流或气体滞留示意图rheomolder which is composed of melting barrel, blunt gas injection pipe, storage tank, nozzle, injection system, etc.The melting barrel is suitable to accommodate rod-shaped materials with the size of U 609300 熔化筒容纳棒状材料mm, which are melted by heating components. The temperatures of storage tank and融化加热元件储存罐material temperature control barrel are considered as the melt temperature and pouringtemperature, respectively. The semisolid slurry is prepared in injection cylinder, which mainly半固态浆料喷油缸consists of the nozzle, runner, and material measurement room.The injection system is used to 喷嘴，流道和材料测量室注入系统inject the slurry into the mold cavity with high pressure and speed, and the forming parts are 浆phone covers (110960 mm) with the thickness of 0.8 mm.In the RM process, specific parameters were as follows: the melt temperature of 670 C,具体参数pouring temperature of 630 and 610 C, cylinder temperature of 570 C, injection pressure注射压力Of 35 MPa, injection velocity of 1.8 m/s, and mold temperature of 250 C. For the purpose ofcomparison, similar AZ91D phone covers (120955 mm) with a section thickness of 0.8 mm weredie casting on a 400-ton cold chamber HPDC machine. During HPDC, pouring temperature of 630 压铸400吨冷室压铸机。

镁合金再结晶英文全文共四篇示例，供读者参考第一篇示例：Magnesium alloy is a popular material used in various industries due to its lightweight and high strength properties. However, when magnesium alloy undergoes deformation, its microstructure changes and the material may lose some of its original properties. In order to restore the mechanical properties of magnesium alloy, it is important to conduct a process called recrystallization.第二篇示例：Overall, recrystallization is a crucial process in the production of magnesium alloys as it can significantly impact the material properties and performance. By understanding the recrystallization process and implementing appropriate processing techniques, researchers and engineers can develop magnesium alloys with improved mechanical properties and formability for various industrial applications.第三篇示例：The process of recrystallization in magnesium alloys is influenced by a number of factors, including the composition of the alloy, the processing conditions, and the temperature at which it is annealed. By carefully controlling these variables, manufacturers can ensure that their magnesium alloys exhibit the desired properties and performance characteristics.第四篇示例：Magnesium alloys are widely known for their lightweight properties and excellent strength-to-weight ratio, making them a popular choice in various industries such as automotive, aerospace, and electronics. However, due to their hexagonal closed-packed (HCP) crystal structure, magnesium alloys are prone to deformation during processing, making them difficult to form and limiting their applicability.。

重庆理工大学文献翻译二级学院材料科学与工程班级 xxxxxxxxxxx学生姓名 weibinglhr 学号 xxxxxxxxxxx译文要求1、译文内容必须与课题（或专业）内容相关，并需注明详细出处。

2、外文翻译译文不少于2000字；外文参考资料阅读量至少3篇（相当于10万外文字符以上）。

3、译文原文（或复印件）应附在译文后备查。

译文评阅导师评语（应根据学校“译文要求”，对学生外文翻译的准确性、翻译数量以及译文的文字表述情况等作具体的评价）指导教师：年月日镁合金在氟化物溶液中的腐蚀行为及钝化机制LI Jian-zhong(李建中)1, 2, HUANG Jiu-gui(黄久贵)2, TIAN Yan-wen(田彦文)1, LIUChang-sheng(刘常升)11. School of Metallurgy and Materials, Northeastern University, Shenyang110004, China;2. Steel Sheet Cold Rolling Plant, Baosteel Branch, Baosteel Co., Ltd., Shanghai200431, ChinaReceived 24 December 2007; accepted 21 August 2008摘要：采用扫描电子显微镜，x射线光谱分析以及电化学测试方法研究了镁合金在氟化氢溶液中的腐蚀行为及钝化机制。

结果表明镁合金经过HF溶液处理以后，在其表面形成了一层不溶于水并具有导电性能的MgF2涂层。

而且，大多数MgF2涂层都将达到一个稳定的值，即镁合金表面的Mg/F质量比为11.3:1。

当H2SO4与HF的体积比小于1.2时，由于活性强镁合金会在两种混合溶液中产生钝化现象，此时大量的镁合金将不再与溶液反应，当两种溶液的体积比超过1.4时，镁合金将会迅速溶解。

钝化层形成以后，MgF2涂层通过吸附HF2- (或者 H2F3-, H3F4-)离子从而保护镁合金不被氟化物溶液腐蚀。

Bioactivity of Mg-ion-implanted zirconia and titaniumH.Liang a ,Y .Z.Wan a ,*,F.He a ,Y .Huang a ,J.D.Xu b ,J.M.Li b ,Y .L.Wang a ,Z.G.Zhao aaSchool of Materials Science and Engineering,Tianjin University,Tianjin 300072,PR China bInstitute of Semiconductors,Chinese Academy of Sciences,Beijing 100083,PR ChinaReceived 8May 2006;received in revised form 7July 2006;accepted 9July 2006Available online 22August 2006AbstractTitanium and zirconia are bioinert materials lacking bioactivity.In this work,surface modification of the two typical biomaterials is conducted by Mg-ion-implantation using a MEVV A ion source in an attempt to increase their bioactivity.Mg ions were implanted into zirconia and titanium with fluences ranging from 1Â1017to 3Â1017ions/cm 2at 40keV .The Mg-implanted samples,as well as control (unimplanted)samples,were immersed in SBF for 7days and then removed to identify the presence of calcium and phosphate (Ca–P)coatings and to characterize their morphology and structure by SEM,XRD,and FT-IR.SEM observations confirm that globular aggregates are formed on the surfaces of the Mg-implanted zirconia and titanium while no precipitates are observed on the control samples.XRD and FT-IR analyses reveal that the deposits are carbonated hydroxyapatite (HAp).Our experimental results demonstrate that Mg-implantation improves the bioactivity of zirconia and titanium.Further,it is found that the degree of bioactivity is adjustable by the ion dose.Mechanisms are proposed to interpret the improvement of bioactivity as a result of Mg implantation and the difference in bioactivity between zirconia and titanium.#2006Elsevier B.V .All rights reserved.PACS:87.68.+z;85.40.RyKeywords:Bioactivity;Titanium;Zirconia;Ion implantation;Hydroxyapatite;Magnesium1.IntroductionIt seems hard to find an artificial biomedical material which shows integrated advantages such as satisfactory mechanical properties,good biocompatibility,and sufficient bioactivity.For instance,numerous in vivo and in vitro assessments have confirmed that calcium and phosphates (Ca–P),especially HAp,support the attachment,differentiation,and proliferation of relevant cells (such as osteoblasts and mesenchymal cells)[1].Though the excellent biological performance of HAp has been well documented,unsatisfactory mechanical strength,especially its brittleness,limit its applications in engineering of new bone tissue,especially at load-bearing sites.Zirconia (ZrO 2)has been widely used in biomedical applications particularly in the high-load bearing sites such as artificial knee and hip in orthopaedic application and dental post-crown in dental application as ball heads for total hip prostheses due to its good biocompatibility,excellent corrosion resistance,lowfriction,high wear resistance,high strength,and low cost [2,3].Titanium and its alloys (such as Ti–6Al–4V and Ti–6Al–7Nb)are also widely used in biomedical applications as artificial hip,orthopaedic,or dental implants because of their superior mechanical properties,acceptable biocompatibility and light-ness.However,both metallic and ceramic materials are not ideal since bone does not bond directly to these metallic or ceramic materials.Further,the metallic materials may corrode away when implanted.Coating is attractive as advantageous properties of two or more types of materials can be combined by taking advantage of the good mechanical performance of metallic or ceramic substrates which controls the load bearing properties of the implants and the bioactivity of Ca–P which controls the interaction with the surrounding tissue.A variety of approaches have been developed to coat Ca–P (especially HAp)on metallic and ceramic substrates.To date,plasma spray [4–6],electrophoretic deposition [7],sol–gel technologies [8,9],magnetron sputtering [10],and biomimetic growth in simulated body fluid (SBF)[11,12]have been adopted.Among these approaches,the biomimetic method has received more and more attention as it is a low temperature process [13].Prior to biomimetic process,namely immersion in SBF,substrate/locate/apsuscApplied Surface Science 253(2007)3326–3333*Corresponding author.Tel.:+862227405056;fax:+86222705056.E-mail address:yzwantju@ (Y .Z.Wan).0169-4332/$–see front matter #2006Elsevier B.V .All rights reserved.doi:10.1016/j.apsusc.2006.07.038materials are usually subjected to surface modification such as pretreatment in NaOH solution.Very recently,ion implantation has been used to surface treat various substrates such as metals [14],ceramics[15],and even polymers[16]aiming at enhancing capability of inducing Ca–P coating on these substrates.Different kinds of elements such as sodium(Na)and Ca have been ion-implanted and enhanced Ca–P growth has been reported[14,17–20].It is a well-known fact that Mg has similarities to Ca and Na.Mg is also naturally found in bonetissue and is essential to human metabolism[21].More importantly,Kuwahara et al.found that apatite could precipitate on the surface of pure Mg from Hank’s solutions [22–24].It is anticipated that Mg incorporation into bioinert materials like zirconia and titanium is most likely to improve Ca–P nucleation and growth.In this work,Mg ions were implanted into a zirconia based bioceramic and commercial pure titanium using a MEVVA (metal vapor vacuum arc)ion source machine.The aim of this study was to evaluate the possibility of triggering Ca–P growth on a typical bioceramic(zirconia)and metal(titanium)by Mg-ion-implantation and assess the influence of Mg concentration on the bioactivity of the two materials.2.Experimental2.1.MaterialsThe zirconia based bioceramic plates were generously provided by the Department of Inorganic Materials,Tianjin University,Tianjin,China.The tetragonal zirconia was stabilized with3mol%yttria(Y2O3)and further stabilized with15wt.% Al2O3addition aiming at avoiding low-temperature degradation. The zirconia based bioceramic used in this work had an average bending strength of600–700MPa and a fracture toughness of 11Æ1.03MPa m1/2.The bulk density of the plates was>99%of the theoretical density.The titanium material used in this work was commercial pure titanium(Baoji Special Iron and Steel Co. Ltd.,Shanxi,China).The zirconia and titanium samples were polished on one side to the mirrorfinish.The specimens were then ultrasonically washed with acetone and distilled water prior to Mg-ion-implantation.2.2.Ion implantationMg-ion-implantation was performed by using an ion implantation machine with a MEVV A ion source.Mg ions were implanted into one side of zirconia and titanium plates with an acceleration energy of40keV and ion-beam current density of8–10m A/cm2.The implantation dose varied from 1Â1017to3Â1017ions/cm2.The vacuum in the target chamber was kept at1Â10À3Pa.2.3.Soaking in SBFTo examine the bioactivity,the Mg-implanted and unim-planted(control)zirconia and titanium samples were immersed in SBF for a time period of7days.The SBF was prepared by dissolving reagent grade NaCl,KCl,NaHCO3,Na2H-PO4Á3H2O,MgClÁ6H2O,CaCl2,and Na2SO4in distilled water. The ionic concentrations were close to those in human blood plasma as listed in Table1.The SBF solution was buffered at pH7.4with tris-hydroxymethyl-aminomethane((CH2OH)3 CNH3)and hydrochloric acid.The solution temperature was kept at378C.After7days incubation in SBF,samples were taken out,washed with distilled water,and dried.2.4.SEM observationThe morphologies of the SBF-soaked samples were observed by a Philips XL30scanning electron microscope (SEM,the Netherlands)attached with the X-ray energy dispersive spectrometer(EDS).2.5.XRD characterizationThe surface structure of the SBF-soaked samples was characterized by a thinfilm X-ray diffractometer(TF-XRD) (Model D/max2500,Rigaku,Japan).The TF-XRD measure-ment was performed using a Cu K a radiation with a wavelength of0.154nm and operated at the condition of40kV and20mA. The XRD patterns were recorded at a scan rate of48/min and a 2u range of10–808.2.6.FT-IR analysisFT-IR analysis was conducted in a Thermo Nicolet Nexus 560Fourier transform infrared spectrometer.In the preparation of samples,the powders scratched from the surface of the SBF-incubated plates were mixed with KBr powder.The mixtures were then pressed to obtain disc specimens.The FT-IR spectra were recorded in a spectral range of4000–400cmÀ1.3.Results3.1.Surface features of Mg-ion-implanted zirconia and titaniumFigs.1and2compare the surface morphologies of zirconia and titanium samples before and after ion-implantation at a dose of2Â1017ions/cm2.A comparison between Fig.1a and b reveals a slight difference in morphology of the unimplanted and Mg-implanted zirconia samples,the latter showing a rougher surface than the former.However,Fig.2indicates that no detectable difference can be found from the Mg-implanted and unimplanted titanium samples.Roughness measurementH.Liang et al./Applied Surface Science253(2007)3326–33333327Table1Ionic concentrations(mM)of SBF used in this workIonsNa+K+Ca2+Mg2+HCO32ÀClÀHPO42ÀSO42ÀSBF142.0 5.0 2.5 1.5 4.2147.8 1.00.5Humanblood plasma142.0 5.0 2.5 1.5 4.2103.0 1.00.5and a comparison between Figs.1and 2reveal that the Mg-implanted zirconia samples (R a =0.025m m)show rougher surface than the corresponding titanium samples (R a =0.010m m)and that small pores are noted on the surface of the Mg-implanted zirconia whereas the surface of titanium is smooth without discernible pores.These findings reveal that the change of surface morphology as a result of ion-implantation is dependent on the nature of the substrate.This is in agreement with previous results that ion-implantation may roughen or polish the surfaces of materials depending on nature of substrate,ion energy,kind of bombarding ions,ion dose,target temperature,etc.[25–29].EDS analyses demonstrate that Mg exists on the surfaces of the Mg-implanted zirconia and titanium.The measured surface Mg concentration on the two materials is listed in Table 2where all samples studied in this work are designated according to their ion dose.It is found from Table 2that the Mg concentration changes with ion dose and with the type of substrate.The Mg concentration increases consistently with ion dose for zirconia.However,in the case of titanium,it increases initially with ion dose and then levels off after the ion dose is larger than 2.0Â1017ion/cm 2,suggesting the existence ofsaturated dose due to sputter effect [30].Note that Ti 2#shows the highest Mg concentration (5.9wt.%)among the six samples studied and ZrO 23#the highest Mg concentration (3.9wt.%)among the three zirconia samples.3.2.Ca–P precipitation on Mg-implanted zirconium and titaniumFig.3displays representative images of the surfaces of the Mg-implanted zirconia samples at various doses after immersionH.Liang et al./Applied Surface Science 253(2007)3326–33333328Fig.1.SEM photos of unimplanted (a)and Mg-ion-implanted (b)zirconiasamples.Fig.2.SEM photos of unimplanted (a)and Mg-ion-implanted (b)titanium samples.Table 2Dependence of surface Mg concentration and growth rate of Ca–P coating on ion dose for the Mg-implanted zirconia and titanium Sample number Dose (ion/cm 2)Concentration of Mg (wt.%)Average coating thickness (m m)ZrO 21# 1.0Â1017 2.0Æ0.89.9Æ0.5ZrO 22# 2.0Â1017 3.4Æ1.116.5Æ1.6ZrO 23# 3.0Â1017 3.9Æ1.020.1Æ1.4Ti 1# 1.0Â1017 4.3Æ1.28.2Æ1.3Ti 2# 2.0Â1017 5.9Æ0.812.8Æ0.9Ti 3#3.0Â10175.7Æ0.112.5Æ1.5in SBF for 7days.SEM observation reveals that no precipitates exist on the surfaces of the control samples and thus their photos are not given.However,as can be seen from Fig.3,abundant globular aggregates are observed on the surfaces of all Mg-implanted zirconia samples.Moreover,the size of the globular aggregates on ZrO 21#is the smallest among the three zirconia samples.Though there is no obvious difference in compactness among the three zirconia samples,the coating thickness on ZrO 21#is the smallest and ZrO 23#the largest (see Table 2),indicating the influence of surface Mg content.The average thickness of the deposits presented in Table 2was obtained by weighing the as-implanted and the SBF-incubated samples and calculating the volume and thickness of the deposits.The coating thickness is believed to be the volume of the deposits divided by the area on which the deposits grow,assuming the surfaces are completely covered by the deposits.Coating thickness is an indication of the growth rate of the precipitates.Similarly,no precipitates are observed on the control samples after 7days immersion in SBF while globular deposits are noted on the Mg-implanted titanium (Fig.4).There are some differences among the three Mg-implanted titanium samples:Ti 1#shows a less compact deposit layer than Ti 2#and Ti 3#.In addition,Fig.4and Table 2demonstrate that the particle size and coating thickness on Ti 1#are smaller in comparison to Ti 2#and Ti 3#.This suggests that Mg concentration on the surface of titanium samples has an impact on the growth rate and particleH.Liang et al./Applied Surface Science 253(2007)3326–33333329Fig.3.SEM micrographs of deposits grown on Mg-implanted zirconia samples at different ion doses after soaking in SBF for 7days:(a)1Â1017ion/cm 2;(b)2Â1017ion/cm 2;(c)3Â1017ion/cm 2.Fig.4.SEM micrographs of deposits grown on Mg-implanted titanium samples at different ion doses after soaking in SBF for 7days:(a)1Â1017ion/cm 2;(b)2Â1017ion/cm 2;(c)3Â1017ion/cm 2.size of the deposits.This finding is consistent with the trend found with zirconia,suggesting that the particle size and growth rate of deposits are proportional to surface Mg concentration,irrespective of the nature of the substrate.Note that the deposits on Ti 2#and Ti 3#show no discernible difference in morphology and coating thickness due to the fact that their surface Mg content is not significantly different.EDS measurement confirms the existence of Ca and P elements on the surfaces of the Mg-implanted zirconia and titanium samples subjected to 7days immersion in SBF,indicating the presence of Ca–P.3.3.XRD resultsXRD patterns of Ca–P coatings on the surfaces of the SBF-incubated Mg-implanted zirconia and titanium samples are presented in Figs.5and 6,respectively.As shown in Fig.5,three characteristic peaks located at 30.1,50.2,and 59.78can be assigned to zirconia.In addition,other peaks located at 25.9,28.1,31.8,32.6,34.2,and 40.58are observed.These peaks can be attributed to HAp crystals [31],indicating that HAp crystals are formed on the Mg-implanted zirconia after soaking in SBF for 7days.As expected,the XRD patterns of the SBF-incubated Mg-implanted titanium samples,shown in Fig.6,consist of peaks corresponding to HAp and titanium.Note that the characteristic peaks of titanium are still intensive in the three XRD patterns,suggesting the coatings are thin when compared to zirconia,which is consistent with SEM observations.It is obvious,at least in the case of zirconia (see Fig.5),that the relative intensity of the HAp peaks increases with the enhancement in the surface Mg concentration,which agrees well with the trend presented inTable 2.The variation of the relative intensity of the HAp peaks with respect to Mg concentration indicates the thickening kinetics of the HAp coatings.The XRD findings reveal that HAp crystals are grown on the Mg-implanted zirconia and titanium samples after soaking in SBF for 7days.3.4.FT-IR resultsThe deposits grown on the Mg-implanted zirconia and titanium were further characterized by FT-IR.The FT-IR resultsH.Liang et al./Applied Surface Science 253(2007)3326–33333330Fig.5.TF-XRD spectra of the Mg-implanted zirconia samples after 7days imm-ersion in SBF:(a)1Â1017ion/cm 2;(b)2Â1017ion/cm 2;(c)3Â1017ion/cm 2.Fig.6.TF-XRD spectra of the Mg-implanted titanium samples after 7days immersion in SBF:(a)1Â1017ion/cm 2;(b)2Â1017ion/cm 2;(c)3Â1017ion/cm 2.Fig.7.FT-IR spectra of deposits on zirconia implanted with Mg at different does after 7days immersion in SBF:(a)1Â1017ion/cm 2;(b)2Â1017ion/cm 2;(c)3Â1017ion/cm 2.are displayed in Figs.7and8.Fig.7shows representative FT-IR spectra from the deposits formed on the three Mg-implanted zirconia samples after soaking in SBF for7days.In Fig.8,only a typical spectrum of the precipitates formed on Ti2#is presented as no difference can be distinguished among the three Mg-implanted titanium samples.As can be seen from the two figures,the existence of PO43À(571,602,1036cmÀ1)[32,33], P–OH band(866cmÀ1)[34,35],P–O band(1318cmÀ1), intercalated H2O(3450and1650cmÀ1),and CO32À(1420and 1460cmÀ1)can be identified,further confirming that HAp crystals are grown on the surfaces of the Mg-implanted zirconia and titanium after incubation in SBF for7days.Furthermore, the presence of CO32-indicates that PO4sites of the HAp structure are partially substituted by carbonate ions[36],i.e., the HAp formed is carbonate-containing.The FT-IR results shown in Figs.7and8demonstrate that the surface Mg concentration studied in this work has no effect on the structure of the deposits.4.DiscussionSEM observations,combined with XRD and FT-IRfindings, reveal that carbonated HAp crystals are grown on the surfaces of the Mg-ion-implanted zirconia(with surface Mg content of 2.0–3.9wt.%)and titanium(with surface Mg content of4.3–5.9wt.%)whereas no precipitates are observed on the control samples.It is well documented that the bioactivity of a biomaterial is associated with its ability of forming a carbonated HAp layer when it is implanted or in contact with biologicalfluids[37–39].From this perspective,the unim-planted samples do not show any bioactivity while the Mg-implanted zirconia and titanium are bioactive.Moreover, combining SEM observations shown in Figs.3and4with the results presented in Table2,we can learn that the growth rate of the HAp crystals is related to the Mg concentration on the surfaces of zirconia and titanium.The higher the Mg content, the faster the HAp crystals grow.It is inferred that the bioactivity of the Mg-implanted zirconia and titanium is dependent on ion dose,more accurately,on the Mg concentration on their surfaces because the rate of HAp formation determines the bioactivity of the substrates[39].It is thus concluded that Mg-ion-implantation improves the bioac-tivity of zirconia and titanium and that the bioactivity is proportional to Mg concentration on their surfaces.Obviously,the improvement of bioactivity is directly related to the existence of Mg.It has been reported that soluble Si,Ca, P and Na ions released from bioactive glasses lead to the improved bioactivity of the bioactive glasses[40,41].Likewise, the improvement of the bioactivity of zirconia and titanium can be ascribed to the release of Mg ions to adjacent solution when immersed in SBF.The detailed mechanisms are described as follows.It is well recognized that Mg is electrochemically a very active metal.A thin layer of magnesium hydroxide (Mg(OH)2)will form when Mg is exposed to air atmosphere [21,42].Note that Mg(OH)2is not stable and corrodes in SBF where chloride ions are present,forming highly soluble magnesium chloride(MgCl2)and hydrogen gas[21].It is therefore believed that Mg(OH)2will form when the Mg-implanted zirconia and titanium are exposed to air.Further,the release of Mg ions and formation of hydrogen gas will occur concurrently when the Mg-implanted samples are soaked in SBF.The release of Mg ions is thought to increase the concentration of Mg ions in the solution adjacent to the surfaces of specimens,which,in turn,results in elevated calcium concentration[43],namely improved supersaturation level.The enhancement of supersaturation level is responsible for the improvement of Ca–P precipitation on the surfaces of the Mg-ion-implanted samples.This mechanism has also been reported by Pham et al.[18].In addition,as proposed by Krupa et al.[44],formation of hydrogen gas favors Ca–P growth since it increases the pH in the solution,which,in turn,accelerates the apatite nucleation by improving the ionic activity of apatite [45].Clearly,a higher Mg concentration on the substrates means a higher concentration of Mg ions in the adjacent solution,leading to enhanced Ca–P growth,which can elucidate the increase of bioactivity with Mg concentration observed for both zirconia and titanium.It should be mentioned that the Mg-implanted zirconia shows a higher Ca–P growth rate than its titanium counterpart though it has a lower surface Mg concentration.This suggests that the improvement of bioactivity is dependent on the type of the substrate.Although some details remain unknown,the distinct surface roughness and morphologies shown in Figs.1 and2are likely to be responsible for the difference in the Ca–P growth.The surface roughness of the Mg-implanted zirconia is higher than the titanium counterpart.This rougher surface contributes to its higher bioactivity in comparison to the Mg-implanted titanium since a high level of roughness and porosity provides more favorable microenvironments and three-dimen-sional nucleating sites for HAp formation[46].As revealed by FT-IR results,the HAp crystals grown on the Mg-implanted zirconia and titanium are carbonate-containing, which is close to biological apatites that contain4–6% carbonate by weight[47].Therefore,the Mg-ion-implanted zirconia and titanium are expected to show good bone-bonding property and thus good biocompatibility as carbonated HApH.Liang et al./Applied Surface Science253(2007)3326–33333331Fig.8.A typical FT-IR spectrum of deposits on Mg-implanted titanium after soaking in SBF for7days.induces a better tissue response due to its close chemical and crystal resemblance to bone mineral when compared to stoichiometric HAp[48].The results presented in this article suggest that Mg-ion-implantation is an effective method of improving the bioactivity of bioinert materials(metals,ceramics,and even polymers), suggesting that further investigation is worthy.It should be stated that more work is needed tofind direct XRD evidence to prove the presence of Mg(OH)2on the surfaces of Mg-ion-implanted zirconia and titanium and that such issues as optimum surface Mg content,kinetics of release of Mg ions (related to Mg depth profile in the surface layer),in vivo properties,and detailed mechanisms should be addressed. 5.ConclusionsOn the surfaces of Mg-implanted zirconia and titanium samples,spherical deposits are observed while no precipitates are grown on the control samples after7days immersion in SBF.XRD and FT-IR results reveal that these deposits are carbonate-containing hadroxyapatite(HAp).The presence of HAp indicates that the two Mg-implanted materials are bioactive.The current experimentalfindings confirm that Mg-implantation improves the bioactivity of zirconia and titanium.The enhancement of the bioactivity is attributable to the release of Mg ions to the solution during incubation in SBF.It is also found that the degree of bioactivity is adjustable by ion dose during ion-implantation process.In other words,the growth rate of HAp,i.e.,the degree of bioactivity,depends on Mg concentration on the surfaces. Higher Mg content corresponds to quicker HAp growth, namely higher degree of bioactivity.Furthermore,it is interesting to note that the Mg-implanted zirconia shows a higher HAp growth rate than the titanium counterpart though its surface Mg content is lower,suggesting that the bioactivity is also related to the type of the substrate.The higher degree of bioactivity of the Mg-implanted zirconia than titanium is likely to be ascribed to its rougher surface and micropores on its surface layer.AcknowledgmentsThe authors acknowledge thefinancial support given by the Tianjin Municipal Science and Technology Committee through Grant numbers05YFSYSF01800and043111511. 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Hot deformation behavior of Mg–3Gd–1Zn (GZ31) magnesium alloy was studied by hot compression tests over the temperature range of 300–500 C under strain rates of 0.0001–0.1 s-1. This material exhibited typical broad single-peak dynamic recrystallization behavior followed by a gradual drop towards the steady state stress.The constitutive behavior of the tested alloy was studied by the power,exponential,and hyperbolic sine laws.The stress multiplier and the hyperbolic sine exponent were calculated as 0.024 MPa-1 and 3.42,respectively.The deformation activation energy was found to be about 173.2 kJ/mol,which is higher than the lattice self-diffusion activation energy of magnesium (135 kJ/mol).The latter can be ascribed to the presence of gadolinium, which shows the importance of rare earth elements in increasing the deformation resistance at high temperatures.在300–500°温度范围内，0.0001–0.1 s-1的应变速率下进行热压缩试验研究Mg–3Gd–1Zn (GZ31)镁合金的热变形行为。