Effect of tin addition on microstructure and electrochemical properties of rolled AZ61-Sn magnes
Effects of annealing treatment on the microstructure and mechanical properties
Effects of annealing treatment on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIPLijuan Zhou a ,⇑,Hongbo Li b ,Yongting Zheng ba School of Materials Science and Engineering,Shandong University of Technology,ZiBo 255049,PR China bCenter for Composite Materials and Structures,Harbin Institute of Technology,Harbin 150001,PR Chinaa r t i c l e i n f o Article history:Received 14July 2012Received in revised form 10November 2012Accepted 12November 2012Available online 20November 2012Keywords:AlN–SiC–TiB 2ceramic Annealing treatment MicrostructureMechanical propertiesa b s t r a c tAnnealing treatments were carried out at different temperatures (600–1200°C)for various holding times (2–10h)to evaluate the microstructural and mechanical properties changes of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIP.The experimental results show that the annealing treatment is bene-ficial to improve the mechanical properties.In the meantime,the desolution of SiC-rich solid solution and precipitation of fine TiB 2grains occurred.Due to the precipitation strengthening by the SiC-rich solid solution and the grain boundary strengthening by fine TiB 2particles,the improvement of the mechanical properties was obvious with higher annealing temperatures and longer holding times.The highest HRA,bending strength and fracture toughness were 94.1,517.5MPa and 5.79MPa m 1/2,respectively,after the annealing treatment at 1200°C for 10h.Ó2012Elsevier B.V.All rights reserved.1.IntroductionAlN–SiC solid solution has a wide range of engineering appli-cations due to their excellent mechanical properties at high tem-peratures and has been investigated considerably in the past decades years [1,2].Many preparation methods,such as hot pressing [3],reaction synthesis [4],spark plasma synthesis sinter-ing [5]and combustion synthesis [6],were adopted to fabricate AlN–SiC ceramics.TiB 2is another important material for high temperature applications because of its high melting point,mechanical properties and electrical conductivity,and relatively low coefficient of thermal expansion.So AlN–SiC–TiB 2ceramics have wide potential applications in the field of electrical conduc-tivity at high temperatures.It is known that combustion synthesis (self-propagating high-temperature synthesis,SHS)is a novel method for preparation of AlN–SiC–TiB 2ceramics with many advantages,such as high efficiency and energy saving.However,combustion synthesis is a high temperature and velocity reaction process,which often result in large amount of non-equilibrium phases and high residual stress,and this deteriorates the proper-ties of the combustion synthesized products.So annealing treat-ment at elevated temperatures (800–1200°C)has been introduced to reduce the residual thermal stress and improve the mechanical properties [7,8].In this paper,the effects of annealing treatment on properties of the AlN–SiC–TiB 2ceramics prepared by self-propagating high tem-perature synthesis and hot isostatic pressing (SHS–HIP)were stud-ied.The influences of the annealing temperature and holding time on the composition,microstructure and mechanical properties were discussed in detail.2.Experimental procedureThe AlN–SiC–TiB 2ceramics were prepared by SHS–HIP reported in reference [9].The reaction for preparing AlN–SiC–TiB 2ceramics is as follows:Al þSiC þTiB 2þN 2!AlN þSiC þTiB 2ð1ÞThe samples with Al 35wt%–SiC 35wt%–TiB 230wt%were chosen to carry out the annealing treatment tests.The annealing treatments were carried out on the box-type electric furnace (SXK-8-16,Longjiang electrical furnace works,Harbin,China)in vacuum atmosphere.The specimens were enclosed and vacuumed in quartz glass tube.During the annealing treatment test,four different temperatures (600,800,1000and 1200°C)and five different holding times (2,4,6,8and 10h)were adopted to study the effects of annealing temperature and holding time on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramics.The composition of the annealing treated samples was analyzed by X-ray dif-fraction (D/max-rB,Rigaku,Japan).The microstructures were observed by scanning electron microscope (SEM,HITACHI S-4700),energy dispersive X-ray spectrometer,and transmission electron microscopy (TEM,Philips CM12/STEM,Holland).The specimen of 3mm Â4mm Â36mm and 2mm Â4mm Â22mm in dimension for bending strength and fracture toughness tests were sliced respectively.The three-point bending strength and fracture toughness were tested on electronic uni-versal test machine with crosshead speed of 0.5mm/min (Instron5569,USA).The Rockwell hardness test was carried out with a ½00ball and 60kg load applied for 30s on the Rockwell hardness test machine (HR-150A,China).0925-8388/$-see front matter Ó2012Elsevier B.V.All rights reserved./10.1016/j.jallcom.2012.11.080Corresponding author.Tel./fax:+865332781357.E-mail address:zhoulijuan@ (L.Zhou).3.Results and discussion3.1.MicrostructureFig.1shows the XRD patterns of theafter the annealing treatment for10h atAs thefigure indicated,oxidationnot occurred and the main phases wereTiB2after the annealing treatment.temperature and nitrogen pressure duringcontent of hexagonal boron nitride(h-BN)ucts.With the increasing of the annealingtion intensity of h-BN phase decreased,TiN phase appeared at the annealing10h,and slightly increases as the annealingto1200°C.The formation of the smallfavorable to the improvement of thermalfracture toughness because of theirand micro-cracking interactionkind of structural ceramic with excellenttion of h-BN phase during preparation anding treatment will not decrease the properties of the products,onthe contrary,it is helpful to the improvement of mechanical properties.Fig.2shows the XRD patterns of AlN–SiC–TiB2ceramics after annealing treatment at1200°C for different holding times.It can be found that the composition of the annealing treated ceramics was almost unchanged at1200°C for8h.When the holding time increased to10h,a small quantity of TiN was formed in the prod-uct,which was consistent with the XRD results shown in Fig.1. Furthermore,the X-ray diffraction peaks of2H–AlN–SiC phase were becoming narrower and sharper with the increasing of the annealing holding time.Meantime,the XRD peaks of the2H–AlN–SiC phase appeared to shift to larger angles,which indicated 3.2.Strengthening and toughening mechanismFig.3shows the BSE images of AlN–SiC–TiB2samples annealed for10h at600and1200°C.AlN–SiC matrix shows gray color and the dispersed white particles are TiB2phase.Fig.3a shows the sam-ple surface annealed at600°C for10h.White columnar TiB2were dispersed inhomogenously in the matrix,and also the sample sur-face presented many pores with different sizes and shapes.When the annealed temperature was1200°C,the dispersion of TiB2 was relatively homogenous andfiner,and the porosity of the sur-face was relatively decreased,as shown in Fig.3b.It is noted that some light gray phase exists between gray ma-trix and white particles,as shown in Fig.4.Line scanning for the light gray region was conducted to determine the phase composi-tion,as shown in Fig.4.The results show that the light gray phase contained large amount of Si element and tiny content of Al,and no Ti element was detected,which can be indicated that the light gray phase was the SiC-rich phase precipitated during the annealing treatment.After the precipitation of SiC,the AlN-rich and SiC-rich solid solutions were formed correspondingly,which was consis-tent with the phase diagram analysis[10].Phase separation leads to grain size refinement,and then results in the improvement of the yield strength.According to Hall–Petch equation,the relationship between the yield strength and grain size is described as follows[11]:ry¼r0þkdÀ1=2ð2Þwhere r y is the yield strength,r0is a materials constant for the starting stress for dislocation movement,k is the strengthening coefficient,and d is the average grain diameter.As the equation indicated,the grain size refinement is in favor of the improvement of yield strength of the materials.Fig.5shows the TEM image and SAED pattern of AlN–SiC–TiB2 ceramic annealed at1200°C for10h.Light gray matrix was AlN–SiC and the black particles embedded in the matrix were TiB2 phase,as shown in Fig.5a.The SAED pattern of the TiB2particles and the calibration were shown in Fig.5b and c,respectively.The average size of these TiB2particles was much less than those in reactant,and also,these particles were completely embedded in the AlN–SiC matrix,which indicated that thefine TiB2particles were precipitated during the annealing treatment.Since combustion synthesis is a high temperature and velocity reaction process,many dislocations,grain boundaries andFig.1.XRD patterns of AlN–SiC–TiB2ceramics annealed at different temperatures for10h.Fig.2.XRD patterns of AlN–SiC–TiB2ceramics annealed at1200°C for different holding times.500L.metastructures can be formed in the products.In addition,boron and titanium could dissolve in the AlN–SiC matrix or in the grain boundaries because of high reaction temperature.What is more,the standard enthalpy of formation (D H )of TiB 2is relatively low (D H =À293kJ mol À1),and the fine TiB 2particles can be precipi-tated from the supersaturated and non-equilibrium AlN–SiC ma-trix during the high temperature annealing treatment.Precipitating of these fine TiB 2particles decreased the lattice distortion,which also released the inner stress and improved the order degree of the lattice.Meantime,composition gradient formed by the precipitation between the matrix and TiB 2phases enhanced the interface bonding strength.In summary,the improvement of the strength and toughness of the AlN–SiC–TiB 2ceramics can be attributed to the particle dis-persion strengthening effect of the precipitated SiC-rich solid solution.And also the fine TiB 2particles precipitated uniformly at the phase interface enhanced the strength of the grain boundaries.nano-particles in AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)TEM image of AlN–SiC–TiB 2composite;(b)SAED of light gray phase in the AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)SEM image of the surface;(b)line Fig.3.Backscattered electron images of AlN–SiC–TiB 2samples annealed at (a)600°C,10h;(b)1200°C,10Fig. 6.Typical SEM fractograph in annealing treated AlN–SiC–TiB2ceramic composite.Fig.7.Rockwell hardness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.8.Bending strength of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.9.Fracture toughness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.annealing temperature reached800and1200°C.The highest HRA,bending strength and fracture toughness were94.1, 517.5MPa and 5.79MPa m1/2,respectively,after the annealing treatment at1200°C for10h.4.ConclusionsThe effects of annealing treatments on the microstructure and mechanical properties of AlN–SiC–TiB2ceramics prepared by SHS–HIP were studied.The annealing treatment can promote the precipitation SiC-rich solid solution in AlN–SiC matrix andfine TiB2particles on the grain boundaries,which improve the mechan-ical properties obviously.With the annealing temperature increas-ing from600to1200°C and holding times increasing from2to 10h,the strengthening effect was more evident.The highest HRA,bending strength and fracture toughness of AlN–SiC–TiB2 ceramics were94.1,517.5MPa and5.79MPa m1/2,respectively, after the annealing treatment at1200°C for10h.AcknowledgmentsThis work has beenfinancially supported by‘‘A Project of Shandong Province Higher Educational Science and Technology Program(J09LD01)’’and‘‘Shandong Province Natural Science Foundation(ZR2010EM045)’’.References[1]L.J.Zhou,Y.T.Zheng,S.Y.Du,H.B.Li,Mater.Sci.Forum546–549(2007)1505–1508.[2]vrenko,J.Desmaison,A.D.Panasyuk,M.Desmaison-Brut,E.Fenard,J.Eur.Ceram.Soc.25(2005)1781–1787.[3]K.Strecker,M.J.Hoffmann,J.Eur.Ceram.Soc.25(2005)801–807.[4]X.M.Yue,G.J.Zhang,Y.M.Wang,J.Eur.Ceram.Soc.19(1999)293–298.[5]M.Hotta,J.Hojo,J.Eur.Ceram.Soc.30(2010)2117–2122.[6]L.Mei,J.T.Li,Acta Mater.56(2008)3543–3549.[7]A.Varma,A.S.Rogachev,A.S.Mukasyan,S.Hwang,Adv.Chem.Eng.24(1998)79–226.[8]D.Sciti,S.Guicciardi,A.Bellosi,J.Eur.Ceram.Soc.21(2001)621–632.[9]L.J.Zhou,Y.T.Zheng,S.Y.Du,Key Eng.Mater.353–358(2007)1517–1520.[10]A.Zangvil,R.Ruh,J.Am.Ceram.Soc.71(1988)884–890.[11]H.Conrad,J.Narayan,Scripta Mater.42(2000)1025–1030.L.Zhou et al./Journal of Alloys and Compounds552(2013)499–503503。
纳米技术用在化妆品上的作文
纳米技术用在化妆品上的作文英文回答：Nanotechnology has revolutionized various industries, and the beauty and cosmetics industry is no exception. The application of nanotechnology in cosmetics has brought about significant advancements and benefits. One of the key areas where nanotechnology has made a remarkable impact is in the formulation and delivery of skincare products.Firstly, nanotechnology enables the development of more effective and targeted skincare products. Nanoparticles can penetrate the skin more efficiently, allowing active ingredients to reach deeper layers and target specific skin concerns. For instance, nanoparticles can encapsulate antioxidants and deliver them directly to the skin cells, providing enhanced protection against free radicals and oxidative stress.Furthermore, nanotechnology has improved the stabilityand longevity of cosmetic products. By reducing the particle size of ingredients, such as pigments and sunscreens, their dispersion and absorption into the skin are improved. This results in better color payoff and longer-lasting effects. Additionally, nanocapsules can be used to encapsulate volatile ingredients, preventing their degradation and enhancing their shelf life.In addition to skincare products, nanotechnology has also influenced the development of makeup products. Nanoparticles can enhance the texture and performance of makeup formulations. For instance, the use of nanoclays in foundation can provide a smooth and lightweight finish, while nanoscale powders in eyeshadows can improve color intensity and adherence.Moreover, nanotechnology has contributed to the development of innovative delivery systems in cosmetics. Nanoemulsions and nanocapsules can encapsulate active ingredients, allowing for controlled release and improved absorption. This enables the formulation of products with enhanced efficacy and targeted benefits. For example,nanocapsules can deliver anti-aging ingredients to specific areas of the skin, such as fine lines and wrinkles,resulting in more visible and long-lasting effects.中文回答：纳米技术在化妆品领域的应用带来了重大的进展和好处。
Effect of heat treatment on microstructure and
Effect of heat treatment on microstructure andtensile properties of A356 alloysPENG Ji-hua1, TANG Xiao-long1, HE Jian-ting1, XU De-ying21. School of Materials Science and Engineering, South China University of Technology,Guangzhou 510640, China;2. Institute of Nonferrous Metal, Guangzhou Jinbang Nonferrous Co. Ltd., Guangzhou 510340, ChinaReceived 17 June 2010; accepted 15 August 2010Abstract: Two heat treatments of A356 alloys with combined addition of rare earth and strontium were conducted. T6 treatment is a long time treatment (solution at 535 °C for 4 h + aging at 150 °C for 15 h). The other treatment is a short time treatment (solution at 550 °C for 2 h + aging at 170 °C for 2 h). The effects of heat treatment on microstructure and tensile properties of the Al-7%Si-0.3%Mg alloys were investigated by optical microscopy, scanning electronic microscopy and tension test. It is found that a 2 h solution at 550 °C is sufficient to make homogenization and saturation of magnesium and silicon in Į(Al) phase, spheroid of eutectic Si phase. Followed by solution, a 2 h artificial aging at 170 °C is almost enough to produce hardening precipitates. Those samples treated with T6 achieve the maximum tensile strength and fracture elongation. With short time treatment (ST), samples can reach 90% of the maximum yield strength, 95% of the maximum strength, and 80% of the maximum elongation.Key words: Al-Si casting alloys; heat treatment; tensile property; microstructural evolution1 IntroductionThe aging-hardenable cast aluminum alloys, such as A356, are being increasingly used in the automotive industry due to their relatively high specific strength and low cost, providing affordable improvements in fuel efficiency. Eutectic structure of A390 can be refined and its properties can be improved by optimized heat treatment [1]. T6 heat treatment is usually used to improve fracture toughness and yield strength. It is reported that those factors influencing the efficiency of heat treatment of Al-Si hypoeutectic alloys include not only the temperature and holding time [2], but also the as-cast microstructure [3í5] and alloying addition [6í8]. Some T6 treatment test method standards of A356 alloys are made in China, USA, and Japan, and they are well accepted. However, they need more than 4 h for solution at 540 °C, and more than 6 h for aging at 150 °C, thus cause substantial energy consumption and low production efficiency. It is beneficial to study a method to cut short the holding time of heat treatment.The T6 heat treatment of Al-7Si-0.3 Mg alloy includes two steps: solution and artificial aging; the solution step is to achieve Į(Al) saturated with Si and Mg and spheroidized Si in eutectic zone, while the artificial aging is to achieve strengthening phase Mg2Si. Recently, it is shown that the spheroidization time of Siis dependant on solution temperature and the original Si particle size [9í11]. A short solution treatment of 30 minat 540 or 550 °C is sufficient to achieve almost the same mechanical property level as that with a solution treatment time of 6 h [12]. From thermal diffusion calculation and test, it is suggested that the optimum solution soaking time at 540 °C is 2 h [13]. The maximum peak aging time was modeled in terms of aging temperature and activation energy [14í15]. According to this model, the peak yield strength of A356 alloy could be reached within 2í4 h when aging at 170 °C. However, few studies are on the effect of combined treatment with short solution and short aging.In our previous study, it was found that the microstructure of A356 alloy could be optimized by the combination of Ti, B, Sr and RE, and the eutecticFoundation item: Project (2008B80703001) supported by Guangdong Provincial Department of Science and Technology, China; Project (09A45031160) supported by Guangzhou Science and Technology Commission, China; Project (ZC2009015) supported by Zengcheng Science andTechnology Bureau, ChinaCorresponding author: PENG Ji-hua; Tel/Fax: +86-20-87113747; E-mail: jhpeng@DOI: 10.1016/S1003-6326(11)60955-2PENG Ji-hua, et al/Trans. Nonferrous Met. Soc. China 21(2011) 1950í1956 1951melting peak temperature was measured to be 574.4 °Cby differential scanning calorimetry (DSC) [16]. In this study, using this alloy modified together with Sr and RE, the effect of different heat treatments on the microstructure and its mechanical properties were investigated.2 ExperimentalCommercial pure aluminum and silicon were melted in a resistance furnace. The alloy was refinedusing Al5TiB master alloy, modified using Al-10Sr andAl-10RE master alloys. The chemical composition ofthis A356 alloy ingot (Table 1) was checked by readingspectrometer SPECTROLAB. Before casting, the hydrogen content of about 0.25 cm 3 per 100 g in the meltwas measured by ELH-III (made in China). Four bars of50 mm×70 mm×120 mm were machined from the sameingot and heat-treated according to Table 2. Followed thesolution, bars were quenched in hot water of 70 °C.Samples cut from the cast ingot and heat-treated barswere ground, polished and etched using 0.5% HF agent.Optical microscope Leica í430 and scanning electricmicroscope LEO 1530 VP with EDS (Inca 300) wereused to examine the microstructure and fractograph. Toquantify the eutectic Si morphology change of differentheat treatments, an image analyzer Image-Pro Plus 6.0 was used, and each measurement included 800í1200 particles. Table 1 Chemical composition of A356 modified with Ti, Sr and RE (mass fraction, %) Si Cu Fe Mn Mg Ti Zn RE Sr 6.85 <0.01 0.19 <0.01 0.370.23 0.03 0.250.012Table 2 Heat treatments in this study Solution Aging Treatment Temperature/ °C Holding time/h Temperature/°C Holdingtime/hST 550r 5 2 170 2T6 535r 5 4 150 15 Tensile specimens were machined from the heat treated bars. The tensile tests were performed using a screw driven Instron tensile testing machine in air at room temperature. The cross-head speed was 1 mm/min. The strain was measured by using an extensometer attached to the sample and with a measuring length of 50mm. The 0.2% proof stress was used as the yield stressof alloys. Three samples were tested for each heat treatment to calculate the mean value.3 Results and discussion3.1 Microstructural characterization of as-cast alloyThe microstructure of as-cast A356 alloy is shown in Fig. 1(a). It is shown that not only the primary Į(Al) dendrite cell is refined, but also the eutectic silicon is modified well. By means of the image analysis, microstructure parameters of as-cast A356 alloy were analyzed statistically as follows: Į(Al) dendrite cell sizeis 76.1 ȝm, silicon particle size is 2.2 ȝm×1.03 ȝm (length×width), and the ratio aspect of silicon is 2.13. The distributions of RE (mish metal rare earth, more than 65% La among them), Ti, Mg, and Sr in the area shown in Fig. 1(b) are presented in Figs. 1(c)í(f)respectively. It is shown that the eutectic silicon particle is usually covered with Sr, which plays a key role in Siparticle modification; Ti and RE present generallyuniform distribution over the area observed, although alittle segregation of RE is observed and shown by arrowin Fig. 1(d). It is suggested that because the refiner TiAl 3and TiB 2 are covered with RE, the refining efficiency isimproved significantly. In the as-cast alloy, some clustersof Mg probably indicate that coarser Mg 2Si phases exist(arrow in Fig. 1(d))).Ti solute can limit the growth of Į(Al) primarydendrite because of its high growth restriction factor [17].The impediment of formation of poisoning Ti-Si compound around TiAl 3 [18] and promotion of Ti(Al 1íx Si x )3 film covering TiB 2 [19] are very important in Al-Si alloy refining. For Al-Si alloys, the effect of RE on the refining efficiency of Ti and B can be contributed to the following causes [20]: preventing refiner phases from poisoning; retarding TiB 2 phase to amass and sink;promoting the Ti(Al, Si)3 compound growth to cover theTiB 2 phase. In this work, with suitable addition of Reand Sr, the microstructure of A356 alloy was optimized. Especially, eutectic Si is modified fully, which isbeneficial to promote Si to spheroidize further duringsolution treatment. 3.2 Microstructural evolution during heat treatmentThe microstructures of A356 alloys treated withsolution at 550 °C for 2 h and ST treatment are presented in Figs. 2(a) and (b) respectively, while those treatedwith solution at 535 °C for 4 h and T6 treatment are presented in Figs. 2(c) and (d), respectively. From Fig. 1 and Fig. 2, after different heat treatments, the primary Į(Al) has been to some extent and the eutectic silicon has been spheroidized further. Both ST and T6 treatmentsproduce almost the same microstructure. The eutectic Si particle distribution and statistical mean aspect ratio of eutectic Si particle are shown in Fig. 3. After onlysolution at 535 °C for 4 h and 550 °C for 2 h, the meanPENG Ji-hua, et al/Trans. Nonferrous Met. Soc. China 21(2011) 1950í19561952Fig. 1 SEM images (a, b), and EDS mapping from (b) for Ti (c), La (d), Mg (e) and Sr (f) in as-cast alloyFig. 2 Microstructure of A356 alloy with different heat treatments: (a) Solution at 550 °C for 2 h; (b) ST treatment; (c) Solution at 535 °C for 4 h; (d) T6 treatmentPENG Ji-hua, et al/Trans. Nonferrous Met. Soc. China 21(2011) 1950í1956 1953Fig. 3 Statistic analysis of eutectic Si in A356 alloy with different heat treatmentsaspect ratios of Si are 1.57 and 1.54 respectively. After being treated by ST and T6, those aspect ratios of Si do not vary greatly, and they are 1.49 and 1.48, respectively. After solution or solution + aging in this study, the friction of eutectic Si particles with aspect ratio of 1.5 is 50%.The eutectic melting onset temperature of Al-7Si-Mg was reported to be more than 560 °C [16, 19]. 550 °C is below the liquid +solid phase zone. During solution, two steps occur simultaneously, i.e., the formation of Al solution saturated with Si and Mg, and spheroidization of fibrous Si particle. The following model predicts that disintegration and spheroidization of eutectic silicon corals are finished at 540 °C after a few minutes (IJmax ) [9]:2maxs 32ʌ..ln 9kT D U UW JI I§· ¨¸©¹ (1) where I denotes the atomic diameter of silicon; Ȗ symbolizes the interfacial energy of the Al/Si interface; ȡ is the original radius of fibrous Si; D s is the inter-diffusion coefficient of Si in Al; and T is the solution temperature. When the D s variation at different temperatures is taken into account, it is plausible to suggest that IJmax at 550 °C is less than IJmax at 540 °C. From Fig. 2(a), it is actually proved that spheroidization of eutectic Si particle could be finished within 2 h when solution at 550 °C.In a selected area of A356 alloy treated with only solution at 550 °C for 2 h (Fig. 4(a)), the distribution of element Mg is presented in Fig. 4(b). Because there is no cluster of Mg in Fig. 4(b), it means a complete dissolution of Si, Mg into Al dendrite during this solution. From the microstructure of A356 alloy treated with T6 (Fig. 5(a)), the distribution of Mg is shown in Fig. 5(b).Fig. 4 SEM image (a) and EDS mapping (b) of Mg distribution in alloy after only solution at 550 °C for 2 hFig. 5 SEM image (a) and EDS mapping of Mg (b) in alloy after heat treatment with T6For A357 alloy with dendrite size of 240 ȝm, uniform diffusion and saturation of Mg in Al could be finished at 540 °C within 2 h [13]. In this study, the cellPENG Ji-hua, et al/Trans. Nonferrous Met. Soc. China 21(2011) 1950í1956 1954size of primary Į(Al) is less than 100 ȝm. It is reasonable that those solutions treated at 535 °C for 4 h and 550 °C for 2 h, can achieve Į(Al) solid solution saturated with Mg and Si because diffusion route is short, even at a higher solution temperature.During aging, Si and Mg2Si phase precipitation happened in the saturated solid solution of Į(Al) according to the sequence in the Al-Mg-Si alloys with excess Si [21]. The needle shaped Mg2Si precipitation was observed to be about 0.5 ȝm in length and less than 50 nm in width, and the silicon precipitates were mainly distributed in Į(Al) dendrites and few of them could be observed in the eutectic region [22]. Because of the small size, these precipitations could not be observed by SEM in this study. However, it is plausible to suggest that the distribution of Mg in dendrite Al cell zone and eutectic zone is uniform (Fig. 4(b) and 5(b)). According to the study by ROMETSCH and SCHAFFER [15], the time to reach peak yield is 2í4 h and 12í14 h at 170 °C and 150 °C, respectively. From 150 to 190 °C of aging temperature, the peak hardness varies between HB110 and HB120. Hence, it is believed that aging at 170 °C for 2 h produces almost the same precipitation hardening as aging at 150 °C for 15 h.3.3 Tensile properties of A356 alloysThe tensile mechanical properties of A356 alloys are given in Table 3. Due to the microstructure optimization of A356 alloy by means of combination of refining and modification, tensile strength and fracture elongation can reach about 210 MPa and 3.7% respectively. Using T6 treatment in this study, strengthand elongation can be improved significantly. For those samples with T6 treatment, the tensile strength and ductility present the maximum values. 90% of the maximum yield strength, 95% of the maximum ultimate strength, and 80% of the maximum elongation can be reached for samples treated by ST treatment. However,T6 treatment spends about 19 h, while ST treatment takes only about 4 h. Fractographs of samples treated with T6 are presented in Fig. 6. The dimple size is almost similar with different heat treatments, indicating that the size and spacing of eutectic silicon particle vary little with different heat treatments. Shrinkage pore, microcrack inside the silicon particle and crack linkage between eutectic silicon particles were observed on the fracture surfaces.Table 3 Tensile properties of A356 alloys with different heat treatmentsHear treatmentıb/MPa ı0.2/MPa į/% As-cast 210 í 3.7 ST 247 178 5.6T6 255 185 7.0Fig. 6 Fractographs of samples with different heat treatments: (a), (b) T6; (c), (d) STPENG Ji-hua, et al/Trans. Nonferrous Met. Soc. China 21(2011) 1950í1956 1955It is well known that shrinkage pores have a great effect on the tensile strength and ductility of A356 alloys. In-situ SEM fracture of A356 alloy indicates the fracture sequence as follows [4]: micro-crack initiation inside silicon particle; formation of slipping band in the Al dendrite; linkage between the macro-crack and micro-crack, and the growth of crack. During tensile strain, inhomogeneous deformation in the microstructure induces internal stresses in the eutectic silicon and Fe-bearing intermetallic particles. Although the full modification of eutectic Si particle was reached in this study, those samples treated with T6 treatment do not perform as well as expected. The main reason is probably due to the higher gas content (0.25 cm3 per 100 g Al). Our next step is to develop a new means to purify the Al-SI alloys to further improve their mechanical properties.4 Conclusions1) The solution at 535 °C for 4 h and the solution at 550 °C for 2 h can reach full spheroidization of Si particle, over saturation of Si and Mg in Į(Al). The heat treatments of T6 and ST produce almost the same microstructure of A356 alloy.2) After both T6 and ST treatments, the aspect ratio of eutectic Si particle will be reduced from 2.13 to less than 1.6, and the friction of eutectic Si particles with aspect ratio of 1.5 is 50%.3) The T6 treatment can make the maximum strength and fracture elongation for A356 alloy. After ST treatment, 90% of the maximum yield strength, 95% of the maximum ultimate strength, and 80% of the maximum elongation can be achieved.References[1]WAN Li, LUO Ji-rong, LAN Guo-dong, LIANG Qiong-hua.Mechanical properties and microstructures of squeezed and casthypereutectic A390 alloy [J]. Journal of Huazhong University ofScience and Technology: Natural Science Edition, 2008, 36(8):92í95. (in Chinese)[2]RAINCON E, LOPEZ H F, CINEROS H. Temperature effects on thetensile properties of cast and heat treated aluminum alloy A319 [J].Mater Sci Eng A, 2009, 519(1í2): 128í140.[3]MANDAL A, CHAKRABORTY M, MURTY B S. Ageingbehaviour of A356 alloy reinforced with in-situ formed TiB2particles [J]. Mater Sci Eng A, 2008, 489(1í2): 220í226.[4]LEE K, KWON Y N, LEE S. Effects of eutectic silicon particles ontensile properties and fracture toughness of A356 aluminum alloysfabricated by low-pressure-casting, casting-forging, and squeeze-casting processes [J]. J Alloys Compounds, 2008, 461(1í2):532í541. [5]VENCL A, BOBIC I, MISKOVIC Z. Effect of thixocasting and heattreatment on the tribological properties of hypoeutectic Al-Si alloy[J]. Wear, 2008, 264 (7í8): 616í623.[6]BIROL Y. Response to artificial ageing of dendritic and globularAl-7Si-Mg alloys [J]. J. Alloys Compounds, 2009, 484(1): 164í167. [7]TOKAJI K. Notch fatigue behaviour in a Sb-modifiedpermanent-mold cast A356-T6 aluminium alloy [J]. Mater Sci Eng A,2005, 396(1í2): 333í340.[8]KLIAUGA A M, VIEIRA E A, FERRANTE M. The influence ofimpurity level and tin addition on the ageing heat treatment of the356 class alloy [J]. Mater Sci Eng A, 2008, 480(1í2): 5í16.[9]OGRIS E, WAHLEN A, LUCHINGER H, UGGOWITZER P J.Onthe silicon spheroidization in Al-Si alloys [J]. J Light Metals, 2002,2(4): 263í269.[10]SJOLANDER E, SEIFEDDINE S. Optimisation of solutiontreatment of cast Al-Si-Cu alloys [J]. Mater Design, 2010, 31(s1):s44ís49.[11]LIU Bin-yi, XUE Ya-jun. Morphology transformation of eutectic Siin Al-Si alloy during solid solution treatment [J]. Special Casting &Nonferrous Alloys, 2006, 26 (12): 802í805. (in Chinese)[12]ZHANG D L, ZHENG L H, STJOHN D H. Effect of a short solutiontreatment time on microstructure and mechanical properties ofmodified Al-7wt.%Si-0.3wt.%Mg alloy [J]. J Light Metals, 2002,2(1): 27í36.[13]YU Z, ZHANG H , SUN B, SHAO G. Optimization of soaking timefor T6 treatment of aluminium alloy [J]. Heat Treatment, 2009, 24(5):17í20. (in Chinese)[14]ESTEY C M, COCKCROFT S L, MAIJER D M, HERMESMANNC. Constitutive behavior of A356 during the quenching operation [J].Mater Sci Eng A, 2004, 383(2): 245í251.[15]ROMETSCH P A, SCHAFFER G B. An age hardening model forAl-7Si-Mg casting alloys [J]. Mater Sci Eng A, 2002, 325(1í2):424í434.[16]TANG Xiao-long, PENG Ji-hua, HUANG Fang-liang, XU De-ying,DU Ri-sheng. Effect of mishmetal RE on microstructures of A356alloy [J]. The Chinese Journal of Nonferrous Metals, 2010, 20(11):2112í2117. (in Chinese)[17]EASTON M A, STJHON D H. A Model of grain refinementincorporation alloy constitution and potency of heterogeneous nucleant particles [J]. Acta Mater, 2001, 49(10): 1867í1878.[18]QIU D, TAYLOR J A, ZHANG M X, KELLY P M. A mechanismfor the poisoning effect of silicon on the grain refinement of Al-Sialloys [J]. Acta Mater, 2007, 55(4): 1447í1456.[19]JUNG H, MANGELINK-NOEL N, BERGMAN C, BILLIA B.Determination of the average nucleation undercooling of primaryAl-phase on refining particles from Al-5.0wt% Ti-1.0wt% B inAl-based alloys using DSC [J]. J Alloys Compounds, 2009, 477(1í2):622í627.[20]LAN Ye-feng, GUO Peng, ZHANG Ji-jun. The effect of rare earthon the refining property of the Al-Ti-B-RE intermediate alloy [J].Foundry Technology, 2005, 26(9): 774í778. (in Chinese)[21]EDWARDS G A, STILLER K, DUNLOP G L, COUPER M J. Theprecipitation sequence in Al-Mg-Si alloys [J]. Acta Mater, 1998,46(11): 3893í3904.[22]RAN G, ZHOU J E, WANG Q G. Precipitates and tensile fracturemechanism in a sand cast A356 aluminum alloy [J]. J Mater ProcessTechnol, 2008, 207(1): 46í52.PENG Ji-hua, et al/Trans. Nonferrous Met. Soc. China 21(2011) 1950í19561956⛁ ⧚ A356䪱 䞥㒘㒛㒧 㛑ⱘ1ē 1ē 1ē ԃ 21. ⧚ ⾥ Ϣ 䰶ˈ 510640˗2. 䞥䙺 㡆 䞥 䰤 㡆䞥 ⷨお ˈ 510340㽕˖⫼ϸ⾡ϡ ⱘ⛁ ⧚ ⿔ 䬊㓐 㒚 䋼ⱘA356 䞥䖯㸠 ⧚ˈϔ⾡ 䭓 䯈 ⧚ T6(535 °C ⒊4 h+150 °C 15 h)ˈ ϔ⾡ ⷁ 䯈ⱘ⛁ ⧚ ST(550 °C ⒊2 h+170 °C 2 h)DŽ䞛⫼ 䬰ǃ ⬉䬰 ⏽ Ԍ 偠ㄝ ↉ ⛁ ⧚ A356 䞥 㾖㒘㒛 Ԍ 㛑ⱘ DŽ㒧 㸼 ˖ 550 °Cϟ ⒊2 h ҹ㦋 MgǃSi䖛佅 Ϩ ⱘĮ(Al) ⒊ԧˈ Փ ⸙Ⳍ⧗ ˗ 㒣170 °CҎ 2 h ˈ ҹ䖒 Ӵ㒳T6 ⧚ⱘ DŽ Ԍ 偠㒧 㸼 ˈA356䪱 䞥㒣Ӵ㒳T6 ⧚ њ 催ⱘ Ԍ 㺖Ԍ䭓⥛˗䗮䖛STⷁ ⛁ ⧚ ˈ Ԍ ǃ Ԍ䭓⥛ ҹ䖒 T6 ⧚ ⱘ90%ˈ95% 80%DŽ䬂䆡˖Al-Si 䞥˗⛁ ⧚˗ Ԍ 㛑˗ 㾖㒘㒛ⓨ(Edited by LI Xiang-qun)。
低真空条件对混凝土中介质迁移性能的影响
第 54 卷第 9 期2023 年 9 月中南大学学报(自然科学版)Journal of Central South University (Science and Technology)V ol.54 No.9Sep. 2023低真空条件对混凝土中介质迁移性能的影响刘梦晶1，谢友均1，上官明辉1，龙广成1，高策2(1. 中南大学 土木工程学院，湖南 长沙，410000；2. 中铁工程设计咨询集团有限公司，北京，100055)摘要：为探索低真空条件对混凝土中介质迁移性能的影响规律，保障低真空管道等严酷环境下混凝土结构服役性能，首先，通过模拟低真空环境条件，研究了经低真空条件处理的混凝土气体渗透性、毛细吸水性和氯离子渗透性能；其次，分析了粉煤灰、硅灰、消泡剂和吸水树脂(SAP)等外加剂对低真空条件下混凝土中气体、离子、水介质迁移性能的影响规律；最后，基于灰色关联理论，探讨了各因素与混凝土中介质迁移性能的关联度。
研究结果表明：低真空条件对混凝土气体渗透性和毛细吸水性产生显著不利影响，复合掺入粉煤灰、硅灰及消泡剂可较好地降低混凝土气体渗透性和毛细吸水性，并且提高混凝土抗氯离子渗透性能，但复合掺入SAP 对混凝土抗渗性改善作用较低；低真空条件处理会显著增加混凝土的水分蒸发量，扩大孔结构，从而导致其抗介质迁移性能比标准养护的差。
复合掺入粉煤灰、硅灰及消泡剂可制得适用于低真空条件的高抗介质迁移性能的混凝土。
关键词：混凝土；低真空条件；介质迁移；气体渗透性；毛细吸水性；氯离子渗透性中图分类号：TQ178 文献标志码：A 文章编号：1672-7207（2023）09-3535-09Effect of low vacuum condition on medium migrationperformance in concreteLIU Mengjing 1, XIE Youjun 1, SHANGGUAN Minghui 1, LONG Guangcheng 1, GAO Ce 2(1. School of Civil Engineering, Central South University, Changsha 410000, China;2. China Railway Engineering Design and Consulting Group Co. Ltd., Beijing 100055, China)Abstract: In order to explore the influence of low vacuum condition on the migration performance of media in concrete and ensure the service performance of concrete structures in harsh environments, such as low vacuum收稿日期： 2022 −09 −12； 修回日期： 2022 −11 −07基金项目(Foundation item)：湖南省重点研发计划项目(2020SK2005)；中国中铁股份有限公司科技研究开发计划项目(2021−重大−09)；中南大学中央高校基本科研业务费专项资金资助项目(2022ZZTS0621) (Project(2020SK2005) supported by the Hunan Provincial Key R & D Plan Program; Project(2021−major −09) supported by the Technology Research and Development Plan of China Railway Corporation Limited; Project(2022ZZTS0621) supported by the Fundamental Research Funds for the Central Universities of Central South University)通信作者：龙广成，博士，教授，从事土木工程材料研究；E-mail ：**********************.cnDOI: 10.11817/j.issn.1672-7207.2023.09.014引用格式： 刘梦晶, 谢友均, 上官明辉, 等. 低真空条件对混凝土中介质迁移性能的影响[J]. 中南大学学报(自然科学版), 2023, 54(9): 3535−3543.Citation: LIU Mengjing, XIE Youjun, SHANGGUAN Minghui, et al. Effect of low vacuum condition on medium migration performance in concrete[J]. Journal of Central South University(Science and Technology), 2023, 54(9): 3535−3543.第 54 卷中南大学学报(自然科学版)pipelines, firstly, by simulating low vacuum environmental condition, the gas permeability, capillary water absorption, and chloride ion permeability of concrete treated under low vacuum condition were studied. Secondly, the influence rules and corresponding mechanisms of admixtures such as fly ash, silica fume, anti-foaming agent, and water absorbing resins on the migration performance of gas, ions, and water in concrete under low vacuum condition were analyzed. Finally, based on the grey correlation theory, the correlation between each factor and the media migration performance in concrete was discussed. The results show that low vacuum condition has significant adverse effects on concrete gas permeability and capillary water absorption. The composite addition offly ash, silica fume and anti-foaming agent can effectively reduce the gas permeability and capillary water absorption of concrete and improve the chloride ion permeability of concrete. However, the composite addition of SAP has a relatively small effect on improving the permeability of concrete. Low vacuum condition significantly increases the water evaporation of concrete, expands the pore structure and leads to poorer resistance to medium migration compared to standard curing. Concrete with high resistance to medium migration suitable for low vacuum condition can be prepared by combining fly ash, silica fume and anti-foaming agent.Key words: concrete; low vacuum condition; media transportation; gas permeability; capillary water absorption; chloride permeability随着现代经济社会和科学技术发展，混凝土结构不断向高原、高海拔乃至深空等环境延伸，近年来，兴起了低真空管道运输系统建设[1]。
纳米技术的利用英语作文
纳米技术的利用英语作文英文回答：In the realm of scientific advancement, nanotechnology has emerged as a game-changer, offering unprecedented opportunities to manipulate matter at the atomic and molecular scale. This extraordinary field holds immense promise for revolutionizing diverse industries and sectors, spanning from healthcare and energy to manufacturing and environmental protection.Healthcare:Nanotechnology presents groundbreaking possibilities in healthcare by enabling targeted drug delivery, early disease detection, and regenerative medicine. Nanoparticles can be engineered to encapsulate and deliver therapeutic agents directly to diseased cells, enhancing efficacy and reducing side effects. Nanosensors can detect minute concentrations of disease markers, facilitating earlydiagnosis and personalized treatment plans. Moreover, nanomaterials can serve as scaffolds for tissue engineering, offering hope for organ repair and regeneration.Energy:Nanotechnology holds the key to unlocking clean and sustainable energy sources. Nano-structured materials can enhance the efficiency of solar cells and fuel cells, enabling more efficient conversion of renewable energy. Nanomaterials also find applications in energy storage systems, improving the capacity and lifespan of batteries. By paving the way for cleaner energy generation and storage, nanotechnology contributes to a more sustainable future.Manufacturing:In the manufacturing sector, nanotechnology introduces novel materials with tailored properties. Nanoscalecoatings can enhance the durability and performance of surfaces, while nano-engineered composites offer improved strength and lightness. Nanotechnology also enables thefabrication of miniaturized and multifunctional devices, leading to advancements in electronics, robotics, and sensing technologies.Environmental Protection:Nanotechnology has the potential to revolutionize environmental protection. Nanoparticles can be employed as catalysts for the degradation of pollutants, reducing air and water contamination. Nanofilters can effectively remove contaminants from water sources, ensuring clean and safe drinking water. Moreover, nanomaterials can enhance the efficiency of solar energy conversion, promoting sustainability and mitigating greenhouse gas emissions.Challenges and Future Directions:Despite its immense potential, the deployment of nanotechnology faces ethical and safety considerations. The potential risks associated with the environmental and human health implications of nanomaterials require careful evaluation. Ongoing research focuses on developingsustainable and biocompatible nanomaterials, as well as establishing robust regulatory frameworks for their responsible use.As we move forward, nanotechnology is poised to shape the future across multiple disciplines. By unlocking new frontiers of scientific exploration and technological innovation, it has the potential to address global challenges, improve human health, and forge a sustainable path for generations to come.中文回答：纳米技术在各个领域中的应用。
Effect of cerium addition on microstructure and texture of aluminum foil for electrolytic capaci
JOURNAL OF RARE EARTHS,Vol.28,No.1,Feb.2010,p.134F j y I M I S &T y j ()I M T y R &D j W NG y (f 6@63)DOI 6S ()666Effect of cerium addition on microstr ucture and texture of aluminum foil for electr olytic capacitor sWANG Haiy an (王海燕)1,2,L I Wenx ue (李文学)1,REN Hu iping(任慧平)1,HUANG L iy in g (黄丽颖)1,WAN G Xia ngy ang(王向阳)1(1.School of Material and Metall urgy Engineering,Inner Mongolia University of Science and Technology,Baotou 014010,China;2School of Materials S cience and En-gineering,University of Science and Technology B eijing,Beijing 100083,China)Received 23October 2008;revised 27August 2009Abstract:Anode foil of aluminum electrolytic capacitor,which requires large surface area for high capacitance,were prepared by rolling,an-nealing and electrochemical etching.Effects of cerium addition on the capacitance of aluminum electrolytic capacitors were investigated.Mi-crostructure of the aluminum foil surface was observed by optical microscopy (OM)and scanning electron microscopy (SEM).Electron back scattered diffraction (EBSD)was also employed to reveal texture evolvement of cold-rolled aluminum foil after annealed.The results showed that addition of cerium contributed to higher <100>texture and higher etch pit density as well.Moreover,the distribution of etch tunnels was more uniformcompared with that without cerium addition.Perfect specificsurface area was obtained with 0.0074%addition of cerium.Keywords:electrolytic capacitor;aluminum foil;rare earthsAluminum electrolytic capacitors are widely used in all types of electronic equipments.High voltage electrolytic ca-pacitors consist of an electrolyte put on an anode made from a 40–100μm thin high-purity aluminum foil with an oxide layer serving as dielectric [1].The capacitance of dielectrics depends on the dielectric constant,the thickness,and the surface area.Among those factors,especially for aluminum electrolytic capacitors,the enlargement of the surface area is most effective from the practical viewpoint because of the trend toward capacitor size miniaturization [2,3].The effective surface can be greatly increased through an etching process that forms narrow channels along the crystallographic <001>directions into the foils.Thus,with regard to the cube-tex-ture in the aluminum foils for electrolytic capacitors,produc-tion has to be optimized so as to provide maximum number of cube-nuclei as well as good growth prospect of the cube-orientation [1,4].A strongly textured material consisting of grains with <001>direction parallel to the sheet normal is favorable.In practical application this is achieved by pro-ducing sheets with a strong cube-texture {001}<100>.To increase the surface area,investigation of pit formation seems very important.Many studies have discussed the pits using high-purity aluminum foils during DC etching in HCl solutions,including the shape,morphology of pitting corro-sion,the tunnel pit growth and the effect of impurities on etching behavior [5].In addition,since the thickness of these foils is only 100–110μm,the tunnel length must be con-trolled to less than 50μm,in order to form a non-piercing layer after etching on double surfaces of Al foil [6].Thus,it is necessary to study the principle of tunnel growth to controltunnel length of pits and maintain a non-piercing layer.It has been recognized that the solution composition in-duces certain mechanical,physical and chemical changes,which can lead to activation of chemical and electrochemical reactions,thus result in accelerated material removal and corrosion [7].In general,hydrochloric acid is used for the etching solution,but the size and distribution of etch pits are not uniform with low density as a result of high corrosive-ness.However,effect of solution composition,especially for rare earth element on pitting corrosion,has not been exten-sively or systematically studied,it can be expected that addi-tion of rare earth during electrochemical etching of alumi-num may induce optimum etch structure,and produces high surface of aluminum foil.Therefore,to increase the specific internal area by reducing the rate of dissolution concentrated on local area and producing dense etch pits,studies of rare earth addition was carried out in this work.1ExperimentalHigh-purity (>99.9972%)aluminum foils with different cerium additions of 0.006wt.%,0.0074wt.%,and 0.012wt.%were used.The impurities in the aluminum are as follows (wt.%):Fe 0.00071,Si 0.00070,Cu 0.00125,Zn 0.00014,and Ti 0.00005.The thickness of aluminum ingot was 18mm,after homogenized treatment,reheating,hot rolling and reversibly cold rolling with a thickness reduction of as much as 98%to a final thickness of 0.11mm.Finally,all speci-mens were isothermally annealed for various times in salt bath furnace.For DC electrolytic etching test,specimensound at ion it em:P ro e ct supported b the nne r ongolia mporta nt cie nc e e chnolog Pro ec t 20071911a nd nner ongolia e chnolog ese arch evelopment Pro ec tCorre sponding a uthor :A H ai an E-ma il:w ind :10.101/1002-07210900-1WANG Haiyan et al.,Effect of cerium addition on microstructure and texture of aluminum foil for electrolytic capacitors 135were cut into 10mm ×55mm,and exposed with an area of 1cm 2to the etching solution.The rest of the specimen was masked with polyester tape.All the specimens were pre-treated in HCl+H 2SO 4(3:1)solution and then rinsed with deionized water before electrolytic etching test.Surface morphology of the recrystallized foils and etch tunnels were observed by means of optical microscopy and scanning electron microscopy (SEM).The evolution of cube-texture during isothermal annealing of the cold rolled foils was tracked by crystallographic texture using electron back scattered diffraction (EBSD)micro texture analysis [8].2Results and discussionFig.1(a –d)show the shape and distribution of the etch pit formed in the solutions with cerium additions of 0wt.%,0.006wt.%,0.0074wt.%,and 0.012wt.%respectively.It is found that there are noticeable differences in the morpholo-gies of etch pits,depending on the additive.The intensity of the cube-orientation {001}<100>changed with the increas-ing of cerium content.As shown in Fig.1(a),when cerium is not applied,about 20%of the tunnels observed at 15μm become alive in the process of tunnel growth,while more tunnels are observed with cerium addition and the survival rate increases in Fig.1(b),(c)and (d).Compared with those with additives,the density of etch pits is low in pure alumi-num foil.The length,widths and depths of individual pits weremeasured for each specimen.As shown in Fig.1,under the conditions without additive and with 0.006wt.%Ce,0.0074wt.%Ce,and 0.012wt.%Ce,the density of etch pit is 0.9×105,3.2×105,4.3×105and 6.1×105pits/cm 2,and the av-erage size is 14.0,12.2,9.0and 8.6μm,respectively.The number of etch pits generated in the etching solutions with rare earth additives is about twice as much as that formed without additive as shown in Fig.1(a).This system is typical of the Al-RE (rare earth)series.Ce-rium exists as stable intermetallic compounds in aluminum,and with a negative enthalpy of mixing in the liquid phase [9].There are some kinds of micro impurity,such as Fe,Si,Mn in aluminum,Ce will form compound of CeFe 2Al 10,CeSi 2A13,CeMn 4Al 8,Ce 2Mn 7Al 10,(CeMg)Al 2,etc.[10]These com-pounds can purify the matrix and eliminate the obstruction of impurity,especially the effect of Fe on nucleation and growth of cube orientation grains during annealing process,and thus increase the cube orientation of aluminum foil,which agrees with the etch pits measurement shown in Fig.1.Fig.2shows the texture evolution of samples with differ-ent rare earth contents cold-rolled and annealed at 530°C for 600s.In Fig.2(a),some minor intensities of cube-orienta-tion {001}<100>are observed.By contrast,with cerium addition,the texture has changed in Fig.2(b).The ODF (ΔΦ=5°)is characterized by very sharp cube-orientation dominating over the rolled texture,for the aluminum foil with 0.0074wt.%pare to the Fig.2(b),the 0.019wt.%Ce addition showed weaker cube-orientationtextureFig.1Recrystallized microstructureof aluminumfoil without Ce (a),with 0.006wt.%Ce (b),with 0.0074wt.%Ce(c)and with 0.012wt.%Ce(d)F T x f f 53f 6(),%()%()ig.2e tureevolution o the samples a terannealed at 0o C or 00s without Ce a with 0.0074wt.Ce b and with 0.012wt.Ce c136JOURNAL OF RARE EARTHS,Vol.28,No.1,Feb.2010(Fig.2(c)).To sum up,annealed samples with cerium addi-tion have higher intensity of cube-orientation {001}<100>and the cube texture reaches the maximum with 0.0074wt.%Ce.With further addition of Ce,cube texture density weak-ens.In order to develop greater understanding of texture evolvement,{111}pole figures,determined from EBSP,for the above specimens with different rare earth contents cold-rolled and annealed at 530°C,are shown in Fig.3,where separate measurements were made on small and large grains.The sample of cold rolled and annealed high purity alu-minum had a random texture with few grains exhibiting cube,S and Cu orientations (see Fig.3(a)),while addition of 0.0074wt.%cerium resulted in sharp cube texture,i.e.,more than 90%{100}<001>orientations with little other texture component are observed in Fig.3(b).With more Ce addi-tions,Fig.3(c)indicates that the orientations are similar to those of the grains shown in Fig.3(b),i.e.,cube texture with minor other texture components.However,the cube texture is less sharp,the contour line tend to sparse,with other ori-entations starting to be presented.It is known that addition of rare earth can refine crys-tal [11,12],and induce boundary corrosion,since higher bound-ary density reduces segregation of impurity atoms and dis-tribution inhomogeneity of current density.Therefore,it is possible to increase the specific electric capacity through de-creasing the grain size under special corrosion conditions [4].With increasing rare earth content,higher defect density and corresponding local concentration of corrosion current den-sity around grain boundaries induce preferred corrosion of rge corrosion holes or channels will appear frequently along the grain boundaries,which lead to the in-homogeneity of corroded structure.Taking the microstructure and texture evolution into ac-count,we conclude that control of cerium addition produces different recrystallization textures during annealing of cold-rolled aluminum,from which the favorite corrosion orienta-tion can be found.When aluminum is electrochemically etched,cubic pits are initially formed and turned into tunnels as the pits grow up.The SEM images of the aluminum etched pit shown in Fig.4corroborate the microstructural evolvement.Fig.4(a)shows that the tunnel length is approximately 20μm.Fig.4(b)indicates that many dark and small tunnels distribute along the tunnel depth in the etched aluminum foils with 0.0074wt.%Ce addition.It is easy to verify that cerium additive induces higher density of small etch pits and contributes their further growth into tunnels by enhancing the mass transport inside one-dimensional tunnel structure.Uniform tunnel length distribution increases the surface area which favors capacitance.It should be noted,however,that large etch pit forms in Fig.4(c)which is rooted in the connections of many small pits ，thus the tunnel wall along its depth will pare to the smaller pits,the surface of tunnel will decrease,as a result of blocking of small pits or tunnels by oxides with the increasing of cerium.The specific capaci-tance obtained by using alternating current impedance [13]shows that the capacitance at first increases,and then de-creases with increasing rare earth addition,it means that the content of rare earth should be controlled to certain extent,and the highest specific surface area was obtained from the specimen with 0.0074wt.%Ce.The effect of cerium on the corrosion of aluminum has been investigated by incorporating cerium into aluminum substrates and observing the subsequent anodic oxidation behavior of the alloys [14].The corrosion resistance of alumi-Fig.3{111}pole figures for sample after annealed at 530°C for600s without Ce (a),with 0.0074wt.%Ce (b)and with 0.012wt.%Ce(c)F S M f f (),%()%()ig.4E images o the aluminum oil without Ce a with 0.0074wt.Ce b and with 0.012wt.Ce cWANG Haiyan et al.,Effect of cerium addition on microstructure and texture of aluminum foil for electrolytic capacitors137num in aqueous media can be attributed to rapid formation of surface oxide film[15].Hydrochloric acid is used as etching solution,and Cl–plays an important role in spot corrosion and development,which has great effect on the oxide scale dissolution of the aluminum foil surface.Cl–is adsorbed on the special location of oxide surface,dissolves the oxide scales and formed Al(OH)2Cl,Al(OH)Cl2,AlCl2+,AlCl3, Al(OH)Cl+,finally produces etch pits on the surface of Al. Cerium and aluminium are oxidized at the alloy/film inter-face and incorporated into the anode film.In general,the oxide layer spread over the macroscopic surface,improves active hole forming[16],thus increases the specific internal area.However,when the content of cerium keeps increasing, the distribution inhomogeneity induces the substantial in-crease of chemistry activity,and results in large number of etching pits,which generates film breakdown over alumi-num surface.During the growth of tunnel,metal dissolution occurs from its surface,while the tunnel sidewalls are cov-ered with oxide film which restrain corrosion,which results in inferior capacitance[17].As mentioned above,we can estimate that the introduction of cerium reduces the thickness of oxide film on aluminum metal and induces the nucleation of etch pits.Therefore,ad-dition of rare earth has certain potential to be used for the application of aluminum electrolytic capacitors,and the in-volved researches are still underway to be carried out.3ConclusionsCerium was introduced to control the etch pits and tunnel length distribution on aluminum foil for electrolytic capaci-tors.The dissolution of cerium during electrochemical etch-ing might induce the optimum etch structure,produce high surface area and finally increase the capacitance,since addi-tion of cerium prohibited initial oxide film formation during etching but favors the formation of corrosion pits.In this work,the most perfect surface area was obtained when the addition of Ce reached0.0074wt.%.References:[1]Engler O,Huh M Y.Evolution of the cube texture in high pu-rity aluminum capacitor foils by continuous recrystallization and subsequent grain growth.Mater.Sci.Eng.A,1999,271(1): 371.[2]Scherer J,Magnussen O M,Ebel T,Behm R J.Initial stages ofthe anodic etching of aluminum foils studied by atomic force microscopy.Corros.Sci.,1999,41(1):35.[3]Oh H J,Lee J H,Ahn H J,Jeong Y S,Park N J,Kim S S,ChiC S.Etching characteristics of high-purity aluminum in hy-drochloric acid solutions.Mater.Sci.Eng.A,2007,449-451(25):348.[4]Mao Weimin,Chen Lei,Sa Liman,Yu Yongning,Li Yunfeng.Influence of grain boundaries on corrosion structure of low voltage aluminum foil.T he Chinese Journal of Nonferrous Metals(in Chin.),2004,14(1):1.[5]Osawa N,Fukuoka K.Pit nucleation behavior of aluminiumfoil for electrolytic capacitors during early stage of DC etching.Corros.Sci.,2000,42(3):585.[6]Wang Mei,He Yedong.Growth and passivation of aluminumetch tunnels at on-off controlling direct current in6wt.%HCl solution.Rare Metals,2008,27(2):205.[7]Liu Guili,Zhang Guoying,Li Rongde.Electronic theoreticalstudy of the interaction between rare earth elements and impu-rities at grain boundaries in steel.J.Rare Earths,2003,21(3): 372.[8]Baker I,Li J.An EBSP study of isothermally-annealed cold-rolled nickel.Microscopy Research and T echnique,2004,63: 289.[9]Kang Y B,Peltona A D,Chartranda P,Fuerst C D.Criticalevaluation and thermodynamic optimization of the Al-Ce,Al-Y, Al-Sc and Mg-Sc binary put.Coupling Phase Diagrams Thermochem.,2008,32:413.[10]Deng Yunlai,Zhang Xinmin,Liu Chunming,Chen Zhiyong,Zhou Zhuoping.The effect of trace RE on the evolution of tex-tures in high voltage aluminum capacitor foils.Functional Material,2002,33(1):60.[11]Li Chunlong,Wang Yunsheng,Chen Jianjun,Liu Chengjun,Jiang Maofa.Effects of rare earth on structure and mechanical properties of clean BNbRE steel.J.Rare Earths,2005,23(4): 470.[12]Lin Qin,Guo Feng,Zhu Xingyuan.Behaviors oflanthanum and ce-rium on grain boundaries in carbon manganese clean steel.J.Rare Earths,2007,25(4):485.[13]Wang Xiangyang.Study of the Electrolytic Etching of Ce Alumi-num Foil for Gigh Voltage Anode of Electrolytic Capacitor.Inner mongolia:InnerMongoliaUniv.Sci.&Tech.,2008.5.[14]Tomcsany L,V arga K,Bartik Ietal.Study of the interaction ofCl–with A passive film on Al and initiation of pitting corrosion.Electrochim.A cta,1989,34(5):855.[15]Mishra A K,Balasubramaniam R.Corrosion inhibition of alu-minium by rare earth chlorides.Mater.Chem.Phy s.,2007, 103(2-3):385.[16]Crossland A C,Thompson G E,Skeldon P,Wood G C,SmithC J E,Habazaki H,Shimizu K.Anodic oxidation of Al-Ce Al-loys and inhibitive behavior of Cerium species.Corros.Sci., 1998,40(6):871.[17]Kang J W,Shin Y H,Tak Y S.Growth of etch pits formedduring sonoelectrochemical etching of aluminum.Electro-chimica A cta,2005,51(5):1012.。
Effect of heat treatment on microstructure and
Effect of heat treatment on microstructure and mechanical properties of extruded ZM61 magnesium alloyF.G.Qi1,2,D.F.Zhang*1,2,Z.T.Zhu1,2,X.X.Xu1,2and G.L.Shi3The effects of T5and T6heat treatments on the microstructure and mechanical properties of extruded Mg–6Zn–1M n alloy were investigated in the present study.The results showed that T5 and T6treatments could markedly improve the strengths of extruded ZM61alloy,and the precipitate strengthening effect of double aging was better than that of single aging.The precipitates formed in the T6treatment were much finer and more dispersive than the ones in the T5treatment,resulting in stronger precipitation strengthening effect.However,due to additional grain size refinement strengthening effect,the T5treatment could improve ductility without sacrificing strength over the T6treatment.Scanning electron microscopy observation and tensile test indicated that different cooling methods after extrusion,such as air cooling and water quenching,had no obvious influence on the microstructure and strengths of extruded and subsequent heat treated Mg–6Zn–1Mn alloys.In addition,the Mn element mainly existed as fine Mn phase particles,which were well dispersed in the matrix.Dispersed Mn particles could be found in rod-like b’1precipitates,but not in the disc shaped b’2precipitates.Keywords:Mg–Zn–Mn,Heat treatment,Precipitate,Microstructure,Mechanical propertiesIntroductionAs the lightest structure metal materials with high specific strength and stiffness,good damping capacity, excellent machinability and good castability,magnesium alloy is the most attractive material in aerospace, transportation and mobile electronics.1–4However,due to limited strength,poor formability and high cost of expensive composition elements used,the application of magnesium alloy is still limited.5–8Therefore,it is pressing to develop some new wrought magnesium alloys with high strength and low cost.Mg–Zn alloys are the most widely used wrought magnesium alloy.9,10 Owing to the problems with hot shortness and coarse and uneven grain size in binary Mg–Zn alloy,commer-cially Mg–Zn alloys are always grain refined by the addition of Zr.11In addition,RE and Cu have been added to improve both casting characteristics and mechanical properties at elevated temperature.12–15 However,the additions of these elements also increase the alloy’s cost.Mg–6Zn–1Mn(ZM61)alloy is a new promising alloy, which is developed to meet the above requirements. Zhang et al.16,17reported that the mechanical properties of ZM61alloy with solution and aging treatment can achieve the level of the commercial wrought ZK60A alloy.16Recently,Park et al.18,19investigated the effect of the Al addition on the microstructure and tensile properties of ZM61alloy and found that ZM61–1Al alloy exhibited excellent tensile properties as a result of refined precipitates by Al addition.More recently,the microstructure and mechanical properties of the Mg–x Zn–1Mn alloy have been reported.20According to the reports,the Mg–6Zn–1Mn(ZM61)alloy had the best comprehensive mechanical properties.It is well known that strengthening via grain size control is particularly effective in magnesium alloys because of the higher Hall–Petch coefficient.21,22For the most part of Mg–Zn series alloys,solution treatment after plastic deformation can significantly coarsen the grain so that the mechanical properties of T6aged alloys are worse than T5aged alloys.23,24It was reported that the hardness and strength of extruded ZK60alloy under T5condition were higher than those under T6condition because the second phase precipitated during the aging process wasfiner and more dispersive under T5 condition than that under T6condition.23The micro-structure and mechanical properties of forged ZK60-Y alloy under various heat treatments have been reported.24The results showed that T5treated alloy had superior tensile strength and plasticity compared with T4and T6treated alloys.1College of Materials Science and Engineering,Chongqing University, Chongqing400045,China2National Engineering Research Center for Magnesium Alloys,Chongqing University,Chongqing400044,China3State Key Laboratory for Fabrication and Processing of Non-Ferrous Metals,General Research Institute for Non-Ferrous Metals,Beijing 100088,China*Corresponding author,email zhangdingfei@1426ß2012Institute of Materials,Minerals and MiningPublished by Maney on behalf of the InstituteReceived15May2012;accepted27July2012DOI10.1179/1743284712Y.0000000095Materials Science and Technology2012VOL28NO12Although some researches on the microstructure of ZM61alloy have been carried out,no systematical study was focused on heat treatment of extruded ZM61alloy.In the present study,the effect of T5and T6heat treatment on the microstructure and mechanical proper-ties of extruded ZM61alloy were investigated.This study also aims to investigate the relationship between precipitations and mechanical properties and to opti-mise the heat treatment parameters.ExperimentalThe nominal composition (in wt-%)of the alloy used in the present study is Mg–6Zn–1Mn.The experimental alloy was prepared from commercial high purity Mg (.99?9%),Zn (.99?95%)and Mg–4?1%Mn master alloy by melting in an electrical resistance furnace under a SO 2z CO 2protective gas and then casting them into a steel mould.The actual composition of alloy was analysed by XRF-800CCDE X-ray fluorescence spec-trometer,and the result is Mg–5?9300Zn–1?0200Mn–0?0094Al–0?0049Fe–0?0058Si–0?0015Cu–0?0005Ni (wt-%).Experimental detail is schematically presented in Fig.1.First,cast ingots were homogenised at 330u C for 24h with air cooling.Before the ingots were extruded,the ingots and extrusion die were heated to 420u C for 90min.To study the effect of the preheating treatment on the microstructure,small samples for microstructure observation were also heat treated with the same heating regime and then quenched in water to retain the high temperature microstructures.Then,the homogenised ingots were hot extruded into bars 16mm in diameter at 420u C.The extrusion ratio was 25:1,and the ram speed was set at 3m min 21during extrusion.To investigate the effect of cooling methods after extrusion on the microstructure and mechanical proper-ties of extruded and subsequent heat treated alloys,different cooling methods of air cooling and water quenching were used.Following this,the samples were given T5or T6heat treatment.In the case of T5treatment,the extruded bars were merely single aged (180u C for 16h)and double aged (90u C for 24h followed by 180u C for 16h)respectively.In the case of T6treatment,the extruded bars were solution treated at 420u C for 2h followed by water quenching and thenimmediately single aged (180u C for 16h)and double aged (90u C for 24h followed by 180u C for 16h)respectively.Cylindrical tensile samples,50mm in gauge length and 5mm in gauge diameter,were machined from the extruded and aged bars along the extrusion direction.Tensile tests were conducted on a Sans CMT-5105electronic universal testing machine at room tempera-ture with a displacement rate of 3mm min 21.Each test condition was repeated at least three times for repeat-ability and accuracy.Microstructure was observed by an optical microscope (NEOPHOT30),a scanning electron microscope (SEM)(TESCAN VEGAII)equipped with an Oxford INCA Energy 350energy dispersive X-ray (EDS)spectrometer.Precipitates were examined using a transmission electrical microscope (Zeiss LIBRA 200FE)operating at 200kV.Phase constitutions were determined by a Rigaku D/max 2500PC X-ray diffractometer with the use of Cu K a radiation and a scanning rate of 4u min 21.Results and discussionMicrostructure of as cast and as homogenised alloysFigure 2shows the microstructures of the as cast and as homogenised ZM61alloys.As shown in Fig.2a and c ,the as cast microstructure of the experiment alloy consists of a -Mg matrix and eutectic compounds.The eutectic com-pounds are Mg 7Zn 3phase by X-ray diffraction (XRD)analysis as shown in Fig.3a .Mn exists as pure a -Mn.The average grain size of as cast alloy is y 160m m.After homogenisation at 330u C for 24h,some of the eutectic compounds in the grain boundary dissolve into the matrix as shown in Fig.2b and d .Figure 3b shows the XRD pattern of the as homogenised ZM61alloy.It is clearly seen that the peaks of the Mg 7Zn 3phase become weaker,and some peaks of the MgZn 2phase are detected,indicating that MgZn 2is precipitated during the Zn diffusion.Microstructure of extruded and solution treated alloysThe preheating microstructures of ZM61alloy at 420u C for 90min and quenching in water is shown in Fig.4.1Extrusion and heat treatment scheduleQi et al.Effect of heat treatment on extruded ZM61magnesium alloyMaterials Science and Technology 2012VOL28NO121427After homogenisation at 330u C for 2h (Fig.3),some of the Mg–Zn eutectic compounds in the grain boundary cannot dissolve completely into the matrix.These undissolved compounds,however,are found to dissolve into the matrix during the preheating of the ingots before extrusion,indicating a low thermal stability of these Mg–Zn compounds.Microstructural changes after hot extrusion with air cooling and water quenching are shown in Fig.5a and b .Owing to the deformation and the occurrence of dynamic recrystallisation during the hot extrusion process,equiaxed grain microstructure is formed,and the average grain size is y 9m m.The effect of preheating treatment at 420u C on the microstructure has already been studied.As is stated above,almost all the eutectic compounds are solutionised into the matrix aftera ,b optical micrographs;c ,d SEM images2Microstructures of a ,c as cast and b ,d as homogenised ZM61alloys3X-ray diffraction patterns of a as cast and b as homo-genised ZM61alloys 4Preheating microstructure of ZM61alloy at 420u C for90min and quenching in waterQi et al.Effect of heat treatment on extruded ZM61magnesium alloy1428Materials Science and Technology 2012VOL28NO12homogenisation and preheating treatment at 420u C.Therefore,little second phase particles are retained,and the complete dynamic recrystallisation happens during extrusion at 420u C,resulting in equiaxed grain.In addition,it is found that there is no difference on the microstructure of extruded ZM61alloys with different cooling methods including air cooling and water quenching.Figure 5c presents an SEM image of ZM61alloy after solution treatment at 420u C for 2h.The average grain size of the solution treated is y 25m m.The dynamic recrystallised grains of the investigated alloy grew up sharply,and all the broken particles dissolved into the matrix,resulting in a high Zn solid solution concentration.The phase evolution was further determined by XRD analysis.Figure 6shows XRD patterns of the extruded and solution treated samples.It is obvious that the diffraction patterns of extruded specimens mainly contain a -Mg matrix,Mn and MgZn 2phase.However,the weak diffraction patterns of the MgZn 2precipitates in the extruded alloy significantly broadened.According to the Scherrer formula,25peak broadening qualitatively illustrates a decrease in grain5a ,b images (SEM)of ZM61alloy after extrusion with a air cooling and b water cooling and c SEM and d TEM imagesof ZM61alloy after solution treated at 420u C for 2h6X-ray diffraction patterns of ZM61alloy after extrusionwith a air cooling and b water cooling and c X-ray dif-fraction patterns of ZM61alloy after solution treated at 420u C for 2hQi et al.Effect of heat treatment on extruded ZM61magnesium alloyMaterials Science and Technology 2012VOL28NO121429size in the corresponding phase,implying that some nanosized MgZn 2precipitates form during the cooling after extrusion.After solution treatment at 420u C for 2h,the diffraction patterns show that the MgZn 2phase disappears,which suggests that a uniform solid–solution structure is produced,as shown in Fig.5c .In addition,the detailed microstructure inside the a -Mg after solution treatment is shown in Fig.5d .From the TEM image,only one spherical phase can be observed.No other phases are detected after solution treatment.Based on the XRD result and previous studies,19,20we can preliminarily conclude that the spherical phase is pure Mn particle.Microstructure of aged alloysFigure 7shows the SEM images of ZM61alloy in the T5(single aging)and T6(single aging)state.Since the SEM images of single aged alloys are very similar to those of double aged alloys,one of them is displayed here.By comparing Figs.5a and 7a ,the alloy in the T5(single aging)state shows the similar microstructure to the extruded alloy.The average grain size of the T5aged alloy is y 11m m.As shown in Figs.5c and 7c ,there is little difference on microstructure between T4treated and T6aged alloys under the SEM observation.In fact,many nanosized Mg–Zn precipitates that are formed during the aging treatment are observed in Fig.8c .Figure 7a and b shows the SEM microstructures of ZM61alloy in the T5(single aging after extrusion with air cooling and water quenching respectively)state.It can be found that there is no obvious change on microstructures under SEM between the two.It is well known that magnesium metal and its alloys have high thermal diffusivity,high thermal conductivity and high efficiency of heat release.26The diameter of extruded bars is only 16mm,so the extruded alloys with air cooling and water quenching have same macrostruc-tures.In addition,the average grain size of the T5treated alloy is much finer than that of the T6treated alloy due to high temperature solution treatment in the latter.Figure 8shows TEM images of ZM61alloy in the T5and T6treatment states.It is observed that two kinds of precipitates formed during aging treatments.Based on previous studies,27–30we can conclude that the two precipitates are rod-like b ’1and disc shaped b ’2phases respectively.The interface between b ’1and the matrix is coherent,while semicoherent between b ’2and the matrix.b ’1phases,which formed as rods with their long axis parallel to the [0001]a direction of the a -Mg matrix,can act as a more enormous impediment to the motion of dislocations than b ’2formed as plates on (0001)a ,as reported in previous studies.27–30In all samples,the precipitates after double aging (Fig.8b and d )are much finer and more dispersed than those after single aging (Fig.8a and c ).It is because the nanosized G.P.zones,which formed during the preaging at 90u C for 24h,could provide more effective nuclei for b ’1phase during the second aging.On the other hand,b ’1and b ’2precipitates in T5treated alloys are relatively less than those in T6states.This is because the Zn solid solubility in T6states is slight higher than that in T5states,and a few broken particles formed after extrusion are grown and retained after T5treatment.In addition,it is observed that many spherical phases are well dispersed in the matrix,which are found in thea ,b T5(single aging after extrusion with a air cooling and b water quenching);c T6(single aging)7Images (SEM)of ZM61alloy at different single agingtreatments conditionsQi et al.Effect of heat treatment on extruded ZM61magnesium alloy1430Materials Science and Technology 2012VOL28NO12rod-like b ’1precipitates but not found in disc shaped b ’2precipitates.As mentioned above,the spherical phase for the solution treated alloy is initially speculated to Mn.Figure 9shows a high angle annular dark field scanning TEM image of ZM61alloy in the T6(double aging)state and the typical EDS result of a spherical phase.It can be seen that the spherical phase is pure a -Mn particle,which can further illustrate the existence form of Mn element.As mentioned above,there are only Mn particles observed in the solution treated alloy,suggesting that Mn particles have a higher thermal stability than MgZn bined with Figs.5d and 8,it is observed that some rod-like b ’1precipitates nucleate on the pre-existing Mn particles.Mechanical properties of ZM61alloyThe mechanical properties of ZM61alloy in different conditions are shown in Fig.10.Figure 10a shows the effect of heat treatment conditions on mechanical properties of extruded ZM61alloy.On one hand,it is noted that double aging could result in a significantincrease in tensile and yield strength as compared with single aging.This is because large amount of G.P.zones,which could act as nuclei for b ’1precipitates,formed during the pre-aging at 90u C for 24h.Therefore,it results in finer and more dispersed b ’1and b ’2precipitates in the second step aging at 180u C for 16h.On the other hand,it is interesting to note that there is no obvious difference in the strengths between T5and T6treated alloys,and the elongations of T5treated alloys are higher than T6treated alloys,as shown in Fig.10a .The strengths of the aged alloys are determined by the combined contributions of grain size refinement strengthening and precipitation strengthening.As the precipitate size is finer and the volume fraction and distribution is larger in the T6treated sample,the precipitation strengthening effect is stronger.However,the average grain size of T5treated alloys is much finer than that of T6treated alloy due to high temperature solution treatment of T6treatment.Therefore,T5treatment can improve ductility without sacrificinga T5(single aging);b T5(double aging);c T6(single aging);d T6(double aging)8Images (TEM)of ZM61alloy at different aging treatments conditionsQi et al.Effect of heat treatment on extruded ZM61magnesium alloyMaterials Science and Technology 2012VOL28NO121431strength over T6treatment due to additional grain size refinement strengthening effect,although precipitation strengthening effect is marginally lower compared to T6treated samples.Figure 10b shows the effect of cooling methods after extrusion on the mechanical properties of extruded and subsequent heat treated alloys.There is no difference in strengths of extruded and subsequent heat treated ZM61alloys with different cooling methods including air cooling and water quenching.As mentioned above,magnesium alloys have high thermal diffusivity,high thermal conductivity and high efficiency of heat release,so the extruded and heat treated ZM61alloys with air cooling and water quenching have same microstructures observed under a scanning electron microscope.There-fore,there are no difference in strengths of extruded and subsequent heat treated ZM61alloys with air cooling and water quenching.ConclusionsThe effects of T5and T6heat treatments on the microstructure and mechanical properties of extruded Mg–6Zn–1Mn alloy have been investigated.The main conclusions can be summarised as follows.1.T5and T6treatments can markedly improve the strengths of extruded ZM61alloy,and T5treatment can improve ductility without sacrificing strength over T6treatment.The precipitates formed in T6treatment are finer and more dense than in T5treatment,resulting in stronger precipitation strengthening effect.However,the grain size of T5treated alloy is much finer compared to T6treated alloy.2.Scanning electron microscopy observation and tensile test reveal that different cooling methods after extrusion,such as air cooling and water quenching,have no obvious influence on microstructure and strengths of extruded and subsequent heat treated ZM61alloys.3.Mn element mainly exists as fine Mn phase particle,which are well dispersed in the matrix.Some rod-like b ’1precipitates nucleate on the Mn dispersoid particles.AcknowledgementsThis work was sponsored by the National Great Theoretic Research Project (grant no.2007CB613700),the National Science and Technology Support Project (grant no.2011BAE22B01-3),the National Natural Science Foundation of China (no.50725413),the International Cooperation Project (grant nos.2010DFR50010and 2008DFR50040),the Chongqing Science and Tech-nology Project (grant no.2010CSTC-HDLS)and the Fundamental Research Funds for the Central Universities (grant no.CDJXS10132202).ReferencesD.Eliezer,E.Aghion andF.H.Froes:‘Magnesium science,technology and applications’,Adv.Perform.Mater.,1998,5,(3),201–212.X.Cao,M.Jahazi,J.P.Immarigeon and W.Wallace:‘A review of laser welding techniques for magnesium alloys’,J.Mater.Process.Technol.,2006,171,(2),188–204.9a high angle annular dark field scanning TEM image ofZM61alloy after T6(double aging)and b corresponding EDS results of point A 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纳米银—金复合颗粒对葡萄糖生物传感器响应灵敏度的增强效应
维普资讯
鄹 窜 既 目 Ｅ固
２０ ０ ２年 第 ２卷 第 ３期
０６ ４ １ ４ ．６
２ ３０７ ０００
任湘菱
唐芳琼 ［ 中国科 学院理化技术研 究所 （ 中国科 学院感 光化学研 究所部分 ） 原 ，
北 京 １０ ０ 】 ０ １１
ｔｒ ｒ ｎｔｕｅ ｔ ｈ ｔｇａ ｈｃＣ ｅ ｓｙｏ ＣＡ ）Ｂ ｉｎ ０ ｏｍｅ ｓ ｔｔｏ、 ｏｏ ｒｐ ｉ ｈｍｉｒ ｆ Ｓ． ｃ ｉｇ１０ Ｉ Ｉ ｉ Ｐ ｔ ｊ １ ） ０
纳米银． 复合颗 粒对 葡萄糖 生物传 感器 响应灵 敏度 的增 强 金
效应
Ｅｎ ａ ｃ ｍ ｅ ｔ ｌ ｃ ｓ ｏ ｅ ｓ ｒＳｅ ｓｔ ｉ ｙ Ａｄ ｉ ｏ ｆＡｇ Ａ ｕ ｈ ｎ ｅ ｎ Ｇ ｕ ｏ ｅ Ｂｉ ｓ ｎ ｏ ｎ ｉ ｖ ｔ ｂ ｄ ｔ ｎ ｏ — ｏｆ ｉ ｙ ｉ
ｍａ ｒｘ ｗｈｉｈ ｉ ｏ ｏ ｅ ｏ Ａｇ Ａｕ ａ ｏ ａ ｔｃｃ ａ ｄ ｏｌｖｉ ｙ ｔ ｉ ｃ ｓ ｃ ｍｐ ｓ ｄ ｆ — ｎ ｎ ｐ ｒｉｌ ｓ ｎ Ｐ ｙ ｎ ｌ
ｂ ｔ ｒｌ（ ｕｙ ａ ＰＶＢ） ｂ ａ ｏ— ｅ ｍｅ ｈ ｄ ｙ ｓ ｌｇ ｌ ｔｏ Ｔｈ ｅ ｐ ｒ ｎｓ ｈ ｗ ｔａ ｔ ｅ ｅ ｘ ｅｉ ｍｅ ｔ ｓ ｏ ｈｔ ｈ
【Ａｂｔａ ｔ ｓｒｃ 】 Ｇｌｃｓ ｉｓｎｏｓｈ ｖ ｂｅ ｆｒ ｄ ｗｉｌ ｌｃｓ ｕｏｅ ｂｏｅ ｓｒ ａ ｅ ｅ ｎ ｏ ｍｅ ｆ ｕ ｏｅ ｌ ｇ
ｏｘｄ ｓ （ ｉａｅ ＧＯＤ ）ｉ ｏ ｉｉｅ ｎ ｃ ｍｐ ｓｔ ｍ ｍｏ ｉｉａｉｎ ｍｅ ｒ ｎ ｍｍ ｂ ｌ ｄ ｉ ｏ ｏ ｉ ｉ ｚ ｅ ｂｌｚ ｔ ｍｂ ａ ｅ ｏ
Ｒ Ｎ ａ ｇ ｉｇ， ＡＮＧ ｎ ｑｏ ｇ （ ｅ ｈ ｉａ ｎ ｔ ｕｅｏ’ ｈ ｓｃ ｎ ｅ ｓｒ ｆ ｉｅｅＡｃ ｄ ｍｙｏ ｉｎ ｅ （ ｈ Ｅ Ｘｉｎ ｌ ｎ Ｔ Ｆａ ｇ ｉｎ Ｔ ｃ ｎｃ ｌ ｓｉ ｔ ｔ ｙ ｉｓａ ｄＣｈ ｍｉｔｙｏ ｎ ｓ ａ ｅ ｆ ｅ ｃ ｓ Ｔ ｅ Ｉ ｔ Ｐ Ｃｈ Ｓｃ