Microstructure and Mechanical Behaviuor of Rotary Friction Welded Titanium Alloys

Microstructure and mechanical properties of wroughtMg-4.1Li-2.5Al-1.7Zn-1Sn alloyRuizhi Wu1,2, a, Dayong Li1,b, Xuhe Liu2,c, and Milin Zhang2,d1College of Materials Science & Engineering, Harbin University of Science & Technology, Harbin,P.R. China 1500802Key Laboratory of Superlight Materials & Surface Technology (Harbin Engineering University),Ministry of Education, Harbin, P.R. China 150001a ruizhiwu2006@,b dyli@,c liuxuhe@,d zhangmilin@Keywords: Mg-Li alloy, deformation, microstructure, mechanical properties.Abstract.An Mg-Li-Al-Zn-Sn alloy was prepared by vacuum melting. The actual content of the elements in the alloy was determined using inductively coupled plasma-atomic emission mass spectrometry (ICP-AEMS). The density of the alloy was detected using Archimedes’ method. Extrusion and rolling deformation were carried out on this alloy. Its microstructures and mechanical properties were then studied with an optical microscope (OM), scanning electronic microscope (SEM), X-ray diffractometer (XRD), energy dispersive spectrometer (EDS), and tensile tester. The extruded alloy was composed of α-Mg and Mg2Sn phases and had good strength and elongation properties as well as a good comprehensive performance. After further rolling deformation, an Al-Li phase appeared due to atomic diffusion during the hot rolling process. Strain-hardening and the strengthening effect of the Al-Li phase further improved the strength of the alloy but decreased its elongation capacity.IntroductionSince the Mg-Li alloy was discovered in 1910, it has attracted a lot of attention from researchers because of its low density, high specific strength, stiffness, good processing performance, and dimensional stability. With these properties, it has shown great potential for use in applications in the aerospace, automotive, electronics, and defense industries [1-3].The Mg-Li alloy is the lightest metal material. Its binary phase diagram shows that, when the lithium content is less than 5.7%, the alloy displays an α single-phase that resembles a close-packed hexagonal structure. When the lithium content is more than 10.3%, the alloy shows a β single-phase, which is a body-centered cubic structure. At lithium contents between 5.7-10.3%, the alloy shows a two-phase organization [4]. Studies have shown that the addition of lithium causes the length of the c-axis in the hexagonal close-packed (HCP) structure to decrease, thus bringing about a decrease in the axial ratio c/a and rendering the alloy more able to undergo dislocation slip. This factor improves the deformability of the alloy [5].Al and Zn are two other elements that are commonly used in alloys. Appropriate amounts of Al added to alloys not only increases their strength and hardness but also improves their ductility and corrosion resistance. The addition of Zn can improve the deformation capacity of alloys [6, 7].Xiang Qi et.al[8] studied the influence of Sn on the microstructure and mechanical properties of a Mg-Li-Al-Zn alloy. Their results indicated that the addition of Sn refined the alloy, thereby improving its strength due to the formation of an Mg2Sn strengthening phase. When the Sn content was 1%, the grain size of the alloy was at the minimum size.Extrusion and rolling are the main deformation methods used for Mg-Li alloys. Deformation not only eliminates some casting defects but also causes dynamic recrystallization under certainconditions [9, 10]. This allows for the formation of refined alloys with improved comprehensiveIn this paper, a Mg-Li-Al-Zn-Sn alloy was prepared. After being subjected to two deformation modes, the microstructure and mechanical properties of the alloy obtained were determined.1. Materials and MethodsPure Mg, Li, Al, Zn, and Sn were used as raw materials. The alloy was created by vacuum induction melting using Ar gas as the protection gas in low-carbon steel molds. Homogenization treatments were performed at 300 o C for 24 h. The alloy ingot was extruded from Φ56 mm to 14 mm at 350 o C and denoted as an as-extruded alloy. As a last step, the extruded alloy was reheated and rolled at 260 o C to a final thickness of 3 mm.The chemical composition of the alloy was tested by inductively coupled plasma-atomic emission mass spectrometry (ICP-AEMS). The density of the alloy was determined by Archimedes’ method. The specimens for microscopic examinations were prepared using standard metallographic sample preparation methods. In brief, the specimens were etched with 1 vol% of nitric acid alcohol for 5-10 s.A LEICA DMIRM and JSM-6480 scanning electron microscopes (SEM) were used to observe the surface and fracture morphology of the alloy. A TTR III X-ray diffractometer (XRD) was utilized to identify the different phases in the alloy. An Energy Dispersive Spectrdmeter (EDS) was used to analysis micro-area composition.The tensile specimens were prepared according to the ASTM E8M-04 standard procedure. Tensile tests were carried out on an Instron4505 electronic universal testing machine with a speed of 1.5 mm/min. Five samples for each test were subjected to analysis along the extrusion and rolling directions.2. Results and Discussion2.1 The composition and density of the alloyThe analysis determined that the alloy composition was made of Mg-4.1Li-2.5Al-1.7Zn-1Sn. The density of the alloy was found to be 1.57 g/cm 32.2 The microstructure of the alloyThe microstructure of the extruded Mg-Li-Al-Zn-Sn alloy is shown in Figures 1a and 1b. The alloy was composed of a single α-phase, although some black matter appeared to be distributed in the matrix material. The grain size was small and its shape was equiaxial. These observations are typical of a material that has undergone dynamic recrystallization. Thus, it can be said that dynamic recrystallization occurred during the extrusion process.The microstructure of the alloy after rolling is shown in Figures 1c and 1d. Except for the black material, there also existed some eutectic compounds in the crystals, which may impact the performance of the alloy. The grain size for as-rolled samples was bigger than that for extrusion alloys.Figure 1. The microstructure of the alloy: (a) As-extruded alloy, (b) Magnified as-extruded alloy, (c) As-rolled alloy, and (d) Magnified as-rolled alloy.3.3 Phase analysisbelonged to Mg 22Sn in the 2Sn exists as a phase in the alloy.There were also sections of the rolled alloy that2θ/(°)I nte nsi t y /a .u .Figure 2. The XRD patterns of the alloy.30µmElements Wt./% At./%Mg 33.68 68.77Al 1.07 2.43Zn 1.47 1.36Sn 63.78 27.47A3.4 Mechanical properties of the alloyThe stress-strain curves of the extruded and rolled alloy are shown in Figure 4. Furthermore, Table 1 lists the mechanical properties of the two deformation state of the alloy. It also lists the corresponding performance parameters of commercially available LA141 Mg-Li alloy.Compared with the LA141 alloy, the as-extruded Mg-4.1Li-2.5Al-1.7Zn-1Sn had a higher tensile and yield strength with a considerable elongation capacity (>20%). Its specific strength and modulus are significantly higher than those of the LA141 alloy. These differences may be due to the α-phase (i.e., the α-phase alloy has higher strength than the β-phase alloy). It is also possible that Mg 2Sn, which was extensively distributed throughout the matrix, hinders dislocation slips when the alloy is deformed, thus playing a role in second phase strengthening.After rolling, the strength of the alloy was further increased, and its tensile strength reached 290.26 MPa. The elongation capacity of the alloy decreased but was still above 10%. The increase in strength may be explained in part by several factors, including work-hardening, additional deformation processes, and an increase in internal dislocation density. The latter causes flow stress to increase and improves the strength of alloys. The existence of an Al-Li phase after rolling could also be another reason for the increase in alloy strength. These two strengthening mechanisms, however, contribute to a decline in alloy plasticity. The alloy grain size increased, leading to a decline in its plasticity too.0510152025050100150200250300A s-extruded A s-rolledT e nsi le stre s s/MP a Strain/%Figure 4. Stress-strain curves of the alloy.Table 1. Mechanical properties of the alloyCondition As-extruded As-rolled LA141Tensile strength, MPa 267.51 290.26 144.69Specific tensile strength, cm ×105 170.39 184.87 105.03Yield strength, MPa 161.27 192.19 124.14Specific Yield strength, cm ×105 102.72 122.41 90.12Elastic modulus, GPa — 57.3 42.1Specific Elastic modulus,×106 — 36.49 31.12Elongation, % 21 11 24Density, g/cm 3 1.57 1.57 1.353.5 Fracture microstructure of the alloyThe fracture microstructure of the alloy is shown in Figure 5. The fracture microstructure of the as-extruded alloy was composed of a large number of small dimples. In some individual dimples,mechanism of the as-extruded alloy was ductile in nature. The fracture microstructure of the as-rolled alloy consisted of cleavage planes and a small number of dimples, which indicate that the fracture mechanism of the as rolled alloy can be ascribed to quasi-cleavage fractures.Figure 5. Fracture microstructure of the alloy. (A) As-extruded alloy and (B) As-rolled alloy.Summary1) An ultra-light Mg-Li-Al-Zn-Sn alloy was prepared by vacuum melting, then it was extruded and rolled. The density of the resulting alloy is 1.57 g/cm 3.2) The Mg(4.1)-Li(2.5)-Al(1.7)-Zn-Sn alloy ingot was subjected to two kinds of deformation processes: extrusion and rolling. The as-extruded alloy was found to be composed of the α-Mg and Mg 2Sn phases. After further rolling deformation, however, the alloy was found to consist of the α-Mg, Mg 2Sn, and Al-Li phases.3) Rolling deformations could further improve the strength of the alloy, but this resulted in a decrease of the elongation capacity.AcknowledgmentsThis work was supported by the National Natural Science Foundation of China (No. 51001034), China Postdoctoral Science Foundation(No. 20100481016) and Heilongjiang Postdoctoral Science Foundation.References[1]. R.Z. Wu, M.L. Zhang: Rev. Adv. Mater. Sci. Vol. 24 (2010), p.14[2]. H.Y Wu, Z.W. Gao and J.Y. Lin: J. Alloys Compd. Vol. 474 (2009), p.158[3]. Z.K. Qu, X.H. Liu, R.Z. Wu and M.L. Zhang: Mater. Sci. Eng. A Vol. 527 (2010), p.3284.[4]. L.Y.Wei, G.L.Dunlop and H.Westengen: Mater. Sci. Technol. Vol. 12 (1996), p.741[5]. C.H.Chiu, H.Y.Wu and J.Y.Wang: J. Alloys Compd. Vol. 460 (2008), p.246[6]. R.Z Wu, M.L Zhang: Mater. Sci. Eng. A Vol. 520 (2009), p.36[7]. T.C.Chang, J.Y.Wang and C.L.Chu: Mater. Lett. Vol. 60 (2006), p.3272[8]. Q. Xiang, R.Z.Wu and M.L. Zhang: J. Alloys Compd. Vol. 477 (2009), p.832[9]. T.C. Chang, J.Y. Wang and C.L. Chu: Mater. Lett. Vol. 60 (2006), p.3272[10]. R. Ninomiya and K. Niyake: J. Jpn. Inst. Met. Vol. 10 (2001), p.509[11]. R.Z Wu, Y.S Deng and M.L Zhang: J. Mater. Sci. Vol. 44 (2009), p.4132[12]. D.K. Xu, L. Liu and Y.B. Xu: Scripta Mater. Vol. 57 (2007), p.285(a)Material and Manufacturing Technology IIdoi:10.4028//AMR.341-342Microstructure and Mechanical Properties of Wrought Mg-4.1Li-2.5Al-1.7Zn-1Sn Alloydoi:10.4028//AMR.341-342.31。

第27卷第1期粉末冶金材料科学与工程2022年2月V ol.27 No.1 Materials Science and Engineering of Powder Metallurgy Feb. 2022DOI:10.19976/ki.43-1448/TF.2021083快速制备C f/SiC复合材料的组织与力学性能王洋，李国栋，于士杰，姜毅(中南大学粉末冶金国家重点实验室，长沙410083)摘要：以SiC粉末、醇溶性酚醛树脂粉末以及炭纤维毡、炭纤维无纬布为原料，采用料浆刷涂−针刺−温压固化−高温碳化工艺，在料浆中酚醛树脂的体积分数分别为10%和15%、温压固化压力分别为8 MPa和20 MPa条件下制备C f/SiC多孔预制体，然后通过化学气相渗透法沉积SiC，快速制备C f/SiC陶瓷基复合材料。

观察和分析复合材料的形貌和组织结构，测定材料的密度、孔隙率、抗弯强度和断裂韧性等性能。

结果表明：料浆中的酚醛树脂体积分数较低时，C f/SiC复合材料的性能较好，并且随固化压力增加而提高。

在酚醛树脂体积分数为10%、温压固化压力为20 MPa条件下得到开孔隙率为13.1%的高致密C f/SiC复合材料，该材料的基体较致密，且纤维束和基体之间基本没有孔隙；当材料受到外加载荷时，通过纤维拔出、纤维脱粘和裂纹偏转来提高复合材料的强度和韧性，断裂方式为假塑性断裂，抗弯强度和断裂韧性都较高，分别为570 MPa和18.6 MPa∙m1/2。

关键词：C f/SiC复合材料；料浆针刺；化学气相渗透；抗弯强度；断裂韧性中图分类号：TB332文献标志码：A 文章编号：1673-0224(2022)01-92-10Microstructure and mechanical properties ofrapid prepared C f/SiC compositesAll Rights Reserved.WANG Yang, LI Guodong, YU Shijie, JIANG Yi(State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China)Abstract: C f/SiC ceramic matrix composites were prepared byslurry brushing-needle punching-warm pressure curing-high temperature carbonization proces using SiC powders, alcohol-soluble phenolic resin powders, carbon fiber felt andcarbon fibers weft free cloth as raw materials. The C f/SiC porous preform was prepared by slurry brush coating, puncture,temperature pressure curing and high temperature carbonization. The volume fraction of phenolic resin in the slurry wasrespectively 10% and 15%, and temperature and pressure curing pressures was 8 MPa and 20 MPa, respectively. Andthen SiC was deposited by chemical vapor infiltration method to quickly prepare C f/SiC ceramic matrix composites. Themorphology and microstructure of the materials were observed and analyzed. The properties such as density, porosity,bending strength and fracture toughness of the materials were measured. The results show that the properties of C f/SiCcomposites are better when the volume fraction of phenolic resin in the slurry is low, and the properties increase with theincrease of curing pressure. When the volume fraction of phenolic resin was 10% and the curing pressure was 20 MPa, ahigh density C f/SiC composite with open porosity of 13.1% is obtained. The matrix of C f/SiC composite was relativelycompact, and there was almost no pores between the fiber bundles and the matrix. When the composite was subjected tothe applied load, the strength and toughness of the composite were improved by fiber pulling out, fiber debonding andcrack deflection. The fracture mode of the composite was pseudoplastic fracture, with the highest bending strength andfracture toughness of 570 MPa and 18.6 MPa·m1/2, respectively.Keywords: C f/SiC composites; slurry needling; chemical vapor infiltration; bending strength; fracture toughness基金项目：轻质高强结构材料国家重点实验室基金资助项目收稿日期：2021−09−16；修订日期：2021−10−29通信作者：李国栋，教授，博士。

Interfacial microstructure and mechanical properties of aluminium –zinc-coated steel joints made by a modifiedmetal inert gas welding –brazing processH.T.Zhang a,⁎,J.C.Feng a ,P.He a ,H.Hackl baState Key Laboratory of Advanced Welding Production Technology,Harbin Institute of Technology,Harbin 150001,Heilongjiang Province,PR ChinabFronius.Internation GMBH,A4600Wels-Thalheim,AustriaReceived 10May 2006;accepted 4July 2006AbstractThe microstructure and properties of aluminium –zinc coated steel lap joints made by a modified metal inert gas CMT welding –brazing process was investigated.It was found that the nature and the thickness of the high-hardness intermetallic compound layer which formed at the interface between the steel and the weld metal during the welding process varied with the heat inputs.From the results of tensile tests,the welding process is shown to be capable of providing sound aluminium –zinc coated steel joints.©2006Elsevier Inc.All rights reserved.Keywords:Welding –brazing;Heat input;Intermetallic compound1.IntroductionIn order to reduce pollution and save energy,it is attractive to make car bodies lighter by introducing some aluminium parts as substitutes for the previous steel structures [1,2].Therefore,joining aluminium to steel has become a major problem,requiring resolution.Direct solid-state joining can be used to make these dissimilar metal joints by controlling the thickness of the interme-tallic compound layer that develops within a few micrometers of the joint interface [3–9].However,the shape and size of such solid-state joints are extremely restricted.Thus,the joining of aluminium to steel byfusion welding methods has been widely studied.As is well known,the joining of aluminium to steel by fusion welding is difficult because of the formation of brittle interface phases which can deteriorate the mechanical properties of the joints.However,Kreimeyer and Sepold [10]have shown that if the layer is less than 10μm thick,the joint will be mechanically sound.In addition,the authors also deem that the existence of a zinc coating increases the wettability of the Al to the steel substrate.As another approach,Achar et al.[11]reported that the thickness of the intermetallic compound layer formed during TIG arc welding of Al to steel is decreased by the use of an Al alloy filler metal containing Si.Murakami et al.[12]and Mathieu et al.[13]both point out that the temperature probably determines the thickness of the intermetallic compound layer of the joint and recom-mended the use of lower heat input to obtain a sound joint.Materials Characterization 58(2007)588–592⁎Corresponding author.Tel.:+8645186412974;fax:+8645186418146.E-mail address:hitzht@ (H.T.Zhang).1044-5803/$-see front matter ©2006Elsevier Inc.All rights reserved.doi:10.1016/j.matchar.2006.07.008The cold metal transfer process,identified here as CMT,is a modified metal inert gas welding process which invented by the Fronius Company.The principal innovation of this method is that the motions of the welding wire have been integrated into the welding process and into the overall control of the process.Every time the short circuit occurs,the digital process-control both interrupts the power supply and controls the re-traction of the wire.The wire retraction motion assists droplet detachment during the short circuit,thus greatlydecreasing the heat input during welding.In this study,we selected the CMT process to join aluminium to zinc-coated steel using a lap geometry. The main purpose of this effort was to reveal the rela-tionship between heat input and the microstructure of the joint.Hardness testing was also used to characterize the phases formed during the welding process.In ad-dition,the quality of the joints was assessed by tensile testing.2.ExperimentalDeep drawn sheets of hot-dip galvanized steel and sheets of pure Al1060with thickness of1mm were used in the welding experiments.An Al sheet was lapped over a Zn-coated steel sheet on the special clamping fixture, and the ending of the weld wire was aimed at the edge of the aluminium sheet,as shown in Fig.1.The MIG welding–brazing was carried out using the CMTwelding source with an expert system and1.2-mm-diameter Al–Si filler metal wire.Argon was used as the shielding gas at a flow rate of15L/min.The surface of the samples was cleaned by acetone before welding.Two sets of welding parameters of different heat inputs were selected,as shown in Table1.The heat input,J,is calculated using the equation:J=(60×UI)/v,where U is the mean welding voltage,I is the mean welding current and v is the welding speed.Typical transverse sections of the samples were observed using optical microscopy(OM)and scanning electron microscopy(SEM).The composition of the intermetallic compound layer at the interface between the steel and the weld metal was determined by energy dispersive X-ray spectroscopy(EDX).Hardness values were obtained using a microindentation hardness tester with a load of10g,and a load time of10s.In addition, the samples were cut in10mm widths,and transverse tensile tests(perpendicular to the welding direction) were used to measure the joint tensilestrength.Fig.1.Schematic plan of the welding process.Table1The welding parametersSamplenumberMeanweldingcurrent(A)Meanweldingvoltage(V)Wire feedrate(m/min)Weldingspeed(mm/min)Weldheatinput(J/cm)Sample A6611.8 3.9762613.2Sample B11013.3 5.4913961.5Fig.2.Front(upper)and back(lower)appearances of typical jointswith different heat inputs:(a)Sample A;(b)Sample B.589H.T.Zhang et al./Materials Characterization58(2007)588–5923.Results and discussion 3.1.Macro-and microstructuresThe appearance of the weld seams for different heat inputs are shown in Fig.2.For all welding cases,a smooth weld seam was made.The molten metal wetted the steel better when using lower heat input,i.e.,compare Sample A at lower heat input to Sample B.This may be related to the different degree of evapo-ration of the zinc coating at different heat inputs.While improving the heat input,the greater evaporation of zinc reduces the wettability of the molten metal on the steel.Fig.3shows a typical cross-section of the joints.Higher heat input (Sample B)resulted in a decrease in the contact angle between the steel and the weld metal.Meanwhile,a special zone with lighter colour at the toe of the weldments can be found (designated by white arrows in Fig.3).Optical micrographs shows that a visible intermetallic compound layer has formed be-tween the steel and weld metal during the welding process,Fig.4.The thickness of the intermetallic com-pound layer changes not only with the location within a given joint but also with the varying heat input between different joints.The thickness of the intermetallic compound layer in the center is greater than at the edge of the seam within one joint.For Sample A,the maximum thickness of the compound layer is about 10μm but is 40–50μm for Sample B.The microstructure of the intermetallic compound is shown in greater detail in the SEM micrographs inFig.5.At lower heat input (Sample A),the inter-metallic compound presents a serrated shape oriented toward the weld metal.When the heat input was increased (Sample B),the compound layer became much thicker and grew into the weld metal with tongue-like penetrations.Anisotropic diffusion is a possible explanation for this irregularity.The intermetallic compounds that form under these conditions generally have an orthorhombic structure (see below).Because of the high vacancy concentration along the c -axis of the orthorhombic structure,Al atoms can diffuse rapidly in this direction and cause rapid growth of the inter-metallic compound.EDX analysis was used to determine the phases of the intermetallic compound layer.The results show that the intermetallic compound layer of the joint made by lower heat input consists entirely of Fe 2Al 5.But when the heat input is increased,the intermetallic compound layer consists of two different phases,the FeAl 2phase near the steel surface and a FeAl 3phase which penetrates toward the weld metal.Thus it is clearthatFig.4.Optical microstructures of interface between steel and weld metal:(a)Sample A;(b)SampleB.Fig.3.Cross-section image at limit of penetration in the joint,showing change in contact angle with increased heat input.Arrows point to an intermetallic compound at the tip of the weld metal:(a)Sample A;(b)Sample B.590H.T.Zhang et al./Materials Characterization 58(2007)588–592the intermetallic compound layer that forms is closely related to the heat input during the welding process.With regard to the special zone designated by white arrows in Fig.3,dendritic-appearing structures can be distinguished on a high-magnification SEM micrograph (Fig.6).EDX analysis results show that such dendrite-shaped crystals of an Al-richα-solid solution containing residual zinc routinely formed at this location.3.2.Hardness measurementsHardness testing results also confirm the presence of a hard intermetallic compound layer.The hardness of the interface layer is much higher than that of the base metal and the weld metal and is found to vary for the corresponding intermetallic compound phases.For the high heat input weld(Sample B)the hardness is much higher,Fig.7.Fig.8.The location where the fracture occurred during tensile testing (designated by white arrows):(a)Sample A;(b)SampleB.Fig.7.Microindentation hardness test results of the joints made using different heatinputs.Fig.6.Dendrite crystal structure at the toe of the weldment(SampleB).Fig.5.SEM micrograph of interface between steel and weld metal:(a)Sample A;(b)Sample B.591H.T.Zhang et al./Materials Characterization58(2007)588–5923.3.Tensile test resultsThe tensile tests were performed to provide a qualitative measure of the joint strength and behavior. These results show that the bond strength is excellent, with the fractures occurring in the HAZ of the Al even when the thickness of the intermetallic compound layer was greater than40μm,Fig.8.From a general view-point,the thickness of the intermetallic compound layer should be controlled to less than10μm in order to obtain a sound joint.This implies that the joint made with higher heat input should have a lower intrinsic strength than the other because of the thicker brittle intermetallic compound layer.However,the intrinsic strength of the joints cannot be determined when the fracture occurs in the HAZ of the pure Al.Nevertheless, according to the thickness of the compound layer,we can presume that the intrinsic strength of the joints should be decreased when increasing the welding heat input.4.ConclusionsBased on the experimental results and discussions, conclusions are drawn as follows1)Dissimilar metal joining of Al to zinc-coated steelsheet without cracking is possible by means of a modified metal inert gas(CMT)welding–brazing process in a lap joint.2)Fe–Al intermetallic compound phases were formedat the interface between the steel and the weld metal.The thickness and the composition of the interme-tallic compound layer varied with weld heat input.3)Despite the formation of the intermetallic compoundphases,the interface between steel and weld metal is not the weakest location of the joints.Tensile tests of the joints caused fractured in the Al HAZ,even when the intermetallic compound layer thickness exceeded 40μm.AcknowledgementsThe authors wish to acknowledge the financial support provided by the National Natural Science Foundation under Grant No.50325517for this work. References[1]Schubert E,Klassen M,Zerner I,Walz C,Sepold G.Light weightstructures produced by laser beam joining for future applications in automobile and aerospace industry.J Mater Process Technol 2001;115:2.[2]Schubert E,Zernet I,Sepold ser beam joining of materialcombinations for automotive applications.Proc SPIE 1997;3097:212.[3]Oikawa H,Ohmiya S,Yoshimura T.Resistance spot welding ofsteel and aluminium sheet using insert metal sheet.Sci Technol Weld Join1999;2:80.[4]Czechowski M.Stress corrosion cracking of explosion weldedsteel–aluminum joints.Mater Corros2004;6:464.[5]Fukumoto S,Tsubakino H.Friction welding process of5052aluminium alloy to304stainless steel.Mater Sci Technol 1999;9:1080.[6]Ochi H,Ogawa K,Suga Y,Iwamoto T,Yamamoto Y.Frictionwelding of aluminum alloy and steel using insert metals.Keikinzoku Yosetsu1994;11:1.[7]Shinoda T,Miyahara K,Ogawa M,Endo S.Friction welding ofaluminium and plain low carbon steel.Weld Int(UK) 2001;6:438.[8]Uzun H,Donne CD.Friction stir welding of dissimilar Al6013-T4to X5CrNi18-10stainless steel.Mater Des2005;1:41. [9]Adler L,Billy M,Quentin G.Evaluation of friction-weldedaluminum-steel bonds using dispersive guided modes of a layered substrate.J Appl Phys2001;12:6072.[10]Kreimeyer M,Sepold ser steel joined aluminium-hybridstructures.Proceedings of ICALEO'02,Jacksonville,USA;2002.[11]Achar DRG,Ruge J,Sundaresan S.Joining aluminum to steel,with particular reference to welding(III).Aluminum1980;4:291.[12]Murakami T,Nakata K.Dissimilar metal joining of aluminum tosteel by MIG arc brazing using flux cored wire.ISIJ Int 2003;10:1596.[13]Mathieu A,Mattei S,Deschamps A.Temperature control in laserbrazing of a steel/aluminium assembly using thermographic measurements.NDT&E Int2006;39:272.592H.T.Zhang et al./Materials Characterization58(2007)588–592。

“龙消防员”可远距离灭火
科学家找到提升仿生珍珠母断裂韧性新机制日本研究人员发明了一台会喷水灭火的新型机
器人——“龙消防员”。

“龙消防员”的消防水带可
改变形状并朝向火焰，由后方的轮式推车中的控制
单元控制。

这辆推车通过供水管连接到一辆装有1.4
万升蓄水箱的消防车上。

喷嘴以6.6升每秒的速度喷水，压强高达1兆帕斯卡。

软管的顶端包含一个传统热成像摄像头，有助于找到发生火灾的位置。

不过，该机器人设计仍有局限。

如对“龙消防员”的被动减振机制仍未落实，这导致其准备飞行的时间过长。

在户外应用中，火灾产生的热量会导致包裹水管和电缆的波纹管产生变形。

（来源：《科技日报》）
中国科学技术大学俞书宏院士团队研究人员将
两种模型整合至仿生结构陶瓷中：一种是常见的仿
珍珠母微米尺度的“砖–泥”结构模型，另一种是广
泛应用于生物和人工材料的纳米尺度梯度结构模型。

研究人员通过构筑氧化石墨烯与有机物的混合框架，
利用框架诱导矿化生长的方法制备具有氧化石墨烯梯度的仿生珍珠母。

结果发现，与没有梯度的仿生珍珠母相比，具有氧化石墨烯梯度的材料表现出更高的内部和外部断裂韧性。

纳米压痕测试表明，具有梯度的仿生珍珠母的基元片能更有效地进行滑动和裂纹偏转，从而实现更高的外部增韧。

研究人员表示，他们目前的工作证明了多尺度结构设计增韧的可行性，为实现超强结构材料的制备提供了一种策略和思路。

（来源：《科技日报》）11。

第27卷第1期粉末冶金材料科学与工程2022年2月V ol.27 No.1 Materials Science and Engineering of Powder Metallurgy Feb. 2022DOI:10.19976/ki.43-1448/TF.2021070基于高通量的原位制备网状结构TiC增强TC4复合材料的组织与性能杜康鸿，柳中强，张建涛，温利平，肖志瑜(华南理工大学国家金属近净成形工程技术研究中心，广州510640)摘要：以不同粒度的TC4合金粉末为基体材料，以VC作为碳源，采用高通量热压烧结工艺，原位制备具有不同网状结构尺寸和不同TiC体积分数(分别为2%、4%和6%)的TiC/TC4钛基复合材料，研究TiC含量和TC4粉末粒度对复合材料组织与性能的影响。

结果表明，TiC/TC4复合材料中的TiC增强颗粒呈网状分布。

与TC4合金相比，TiC/TC4复合材料的组织明显细化。

随TiC含量增加，TiC网状结构的厚度增大，材料的抗拉强度与伸长率先升高后下降，TiC含量为2%的复合材料综合性能最优。

随TC4粉末粒度减小，TiC/TC4复合材料中的基体组织逐渐细化，基体的连通性提高，材料抗拉强度与伸长率同时提高。

采用粒度为40～80 μm的TC4合金粉末为原料制备的2%TiC/TC4复合材料，网状结构尺寸小，综合性能最优，屈服强度、抗拉强度和伸长率分别为946 MPa、1058 MPa和18.1%，较TC4合金分别提高29.6%、31.6%和118.1%。

关键词：钛基复合材料；原位碳化钛；高通量制备；网状结构；显微组织；力学性能中图分类号：TB331文献标志码：A 文章编号：1673-0224(2022)01-56-10Microstructure and properties of high-throughput in situ networkAll Rights Reserved.structure TiC reinforced TC4 composite materialsDU Kanghong, LIU Zhongqiang, ZHANG Jiantao, WEN Liping, XIAO Zhiyu(National Engineering Research of Net-Shape Forming for Metallic Material,South China University of Technology, Guangzhou 510640, China)Abstract: Using TC4 alloy powders with different particle sizes as the matrix material, using VC as the carbon source,using high throughput hot press sintering process, TiC/TC4 composites were prepared with different network structuresizes and TiC volume fractions (2%, 4%, 6%). The effect of TiC content and TC4 powder particle sizes on themicrostructure and properties of composite materials were studied. The results show that the TiC reinforcement particlesin TiC/TC4 titanium matrix composites are distributed in a network. Compared with TC4 alloy, the microstructure ofTiC/TC4 composite materials are significantly refined. As the TiC content increases, the thickness of the TiC networklayer increases, and the tensile strength and elongation of the material first increase and then decrease. The material with2%TiC has the best overall performance. As the particle size of TC4 decreases, the microstructure of the TiC/TC4composite is gradually refined, the connectivity of the matrix increases, and the tensile strength and elongation of thematerial increase at the same time. 2%TiC/TC4 composite material prepared by TC4 alloy powders with a particle size of40−80 μm has a small network structure and the best overall performance. The yield strength, tensile strength andelongation reach 946 MPa, 1058 MPa and 18.1%, respectively, which are 29.6%, 31.6%, and 118.1% higher than TC4alloy.Keywords:titanium matrix composites; in-situ titanium carbide; high-throughput preparation; network structure;microstructure; mechanical properties基金项目：广东省重大科技攻关项目(2019B010942001)；国家自然科学基金资助项目(51627805)收稿日期：2021−08−17；修订日期：2021−10−26通信作者：肖志瑜，教授，博士。

三维亥姆霍兹线圈驱动的仿生鳐鱼微机器人目录1. 内容描述 (2)1.1 研究背景 (2)1.2 研究意义 (3)1.3 文献综述 (4)1.4 研究方法与技术路线 (5)1.5 论文结构安排 (6)2. 三维亥姆霍兹线圈原理 (8)2.1 亥姆霍兹线圈的物理模型 (9)2.2 三维亥姆霍兹线圈的磁场分析 (9)2.3 磁场与流体动力学相互作用 (10)3. 仿生鳐鱼微机器人设计 (11)3.1 鳐鱼运动机制研究 (13)3.2 仿生鳐鱼微机器人的总体设计 (15)3.3 驱动系统的设计与选择 (16)3.4 仿生机构与控制策略 (17)4. 三维亥姆霍兹线圈驱动系统的设计与实现 (18)4.1 线圈系统的设计 (20)4.2 电源系统的设计 (22)4.3 线圈与微机器人的集成 (23)4.4 系统调试与性能评估 (24)5. 仿生鳐鱼微机器人的运动测试 (25)5.1 水下测试环境设置 (27)5.2 运动特性的测试与分析 (28)5.3 传感反馈与自主导航能力的测试 (29)5.4 应用场景模拟与性能评估 (30)6. 结论与展望 (31)6.1 研究结论 (32)6.2 存在的问题与不足 (33)6.3 未来研究方向与展望 (34)1. 内容描述本文介绍了一种基于三维亥姆霍兹线圈驱动的仿生鳐鱼微机器人。

该微机器人采用仿生设计，结构简化且生物相容性强，模仿了鳐鱼柔性尾鳍的运动模式。

通过利用三维亥姆霍兹线圈技术实现无接触驱动，微机器人能够在水中产生流畅的自适应运动，并具备灵活的转向和控制能力。

本文将详细阐述微机器人的设计理念、结构特点、工作原理以及运动性能评估结果。

探索了该仿生微机器人在水下微环境探测、生物医学应用等方面的潜在应用前景。

1.1 研究背景随着微纳米技术的飞速发展，微型机器人（micbotics）已经成为科学研究的热门领域，特别是在生物医学、工业检测和环境保护等方面展现出巨大的应用前景。

塑料作为一种被广泛应用的有机合成聚合物材料，在为我们生活提供便利的同时，也带来了后续的环境问题。

据估算，到2050年，将有大约12000万t 塑料垃圾被埋入垃圾填埋场或自然环境中[1]，塑料垃圾进入到环境后会逐渐破碎变成微塑料（<5mm ），导致其在土壤和水体中的丰度逐年递增[2-3]。

2015年的第二届联合国环境大会已将微塑料污染列为环境与生态领域的第二大科学问题[3]。

目前，人们对水体中微塑料的认识已较为系统[4-6]，土壤微塑料逐步成为新的研究热点[7-8]。

2012年Rillig [9]首次提出微塑料会影响土壤理化性质，这引起了人们对土壤微塑料的关注，后续研究表明微塑料可被植物吸收并积累[10-11]，最终通过食物链进入人体。

厘清微塑料对植物生长的影响及其机制，有助于系统掌握其在土壤-植物体陈欣，郭薇，李济之，等.土壤微塑料影响植物生长的因素与机制研究进展[J].农业环境科学学报,2024,43（3）：488-495.CHEN X,GUO W,LI J Z,et al.Research progress on the influencing factors and mechanisms of soil microplastics on plant growth[J].Journal of Agro-Environment Science ,2024,43（3）：488-495.土壤微塑料影响植物生长的因素与机制研究进展陈欣1，郭薇1，2，李济之1，2，迟光宇1*（1.中国科学院沈阳应用生态研究所，污染生态与环境工程重点实验室，沈阳110016；2.中国科学院大学，北京100049）Research progress on the influencing factors and mechanisms of soil microplastics on plant growthCHEN Xin 1,GUO Wei 1,2,LI Jizhi 1,2,CHI Guangyu 1*（1.Key Laboratory of Pollution Ecology and Environmental Engineering,Institute of Applied Ecology,Chinese Academy of Sciences,Shenyang 110016,China ；2.University of Chinese Academy of Sciences,Beijing 100049,China ）Abstract ：Microplastics in soil can affect plant growth in a variety of ways,accumulate in plants,and eventually enter the human body via the food chain.Clarifying the mechanisms and main factors whereby microplastics influence plant growth can contribute to a systematic understanding of their environmental behavior in soil-plant systems.Both the occurrence state and physicochemical characteristics of microplastics can influence their effects on plants.In this paper,from the perspectives of particle size,shape,concentration,and type,plastic additives,and aging degree of microplastics,we review the main factors and mechanisms underlying the effects of soil microplastics on plant growth.The key direction of future research is proposed,which will provide a reference for further clarifying the impact of microplastics on soil ecosystems.Keywords ：soil;microplastics;plant;influencing factor;mechanism of action收稿日期：2023-04-13录用日期：2023-06-19作者简介：陈欣（1968—），男，辽宁沈阳人，博士，研究员，研究方向为农业生态。

第27卷第1期粉末冶金材料科学与工程2022年2月V ol.27 No.1 Materials Science and Engineering of Powder Metallurgy Feb. 2022DOI:10.19976/ki.43-1448/TF.2021088放电等离子体烧结Al-4.5Cu合金的组织与性能穆迪琨祺1，曹磊1，张震1，梁加淼1，张德良2，王俊1(1. 上海交通大学上海市先进高温材料及其精密成型重点实验室，上海200240；2. 东北大学材料科学与工程学院，沈阳110819)摘要：利用放电等离子烧结制备Al-4.5Cu(质量分数，%)合金，并对其进行固溶、淬火和时效处理。

通过X 射线衍射、扫描电镜和透射电镜进行结构表征以及拉伸力学性能测试，研究颗粒界面结构和界面析出行为及其对力学性能影响。

结果表明：放电等离子烧结Al-4.5Cu合金颗粒界面由Al2O3纳米颗粒、粗大CuAl2相和纳米微孔组成。

热处理后，Al-4.5Cu合金颗粒界面附近析出尺寸为150～600 nm的CuAl2相，同时形成宽度为40～60 nm 的无析出区。

屈服强度和抗拉强度分别从95 MPa 和229 MPa增加至280 MPa和378 MPa，断后伸长率从11.8%下降为6.0%。

强度增加主要归因于热处理过程中析出相的弥散分布，以及材料的致密化；塑性下降主要是由于拉伸变形过程中，无析出区率先发生塑性变形，导致位错从无析出区向颗粒界面附近的CuAl2相堆积，造成应力集中，促使裂纹沿颗粒界面扩展，材料伸长率下降。

关键词：粉末冶金；颗粒界面；析出；微观组织；力学性能中图分类号：TG156文献标志码：A 文章编号：1673-0224(2022)01-24-10The microstructure and mechanical properties ofAl-4.5Cu alloyAll Rights Reserved.fabricated by spark plasma sinteringMU Dikunqi1, CAO Lei1, ZHANG Zhen1, LIANG Jiamiao1, ZHANG Deliang2, WANG Jun1(1. Shanghai Key Laboratory of Advanced High-Temperature Materials and Forming,Shanghai Jiao Tong University, Shanghai 200240, China;2. School of Materials Science and Engineering, Northeastern University, Shenyang 110819, China)Abstract:Al-4.5Cu (mass fraction, %) alloy was prepared by spark plasma sintering (SPS) followed by solution,quenching and aging. The X-ray diffraction, scanning electron microscopy, transmission electronmicroscopy and tensiletests were carried out. The effect of interparticle boundary (IPB) and precipitation behavior on mechanical properties ofthe Al-4.5Cu alloy were investigated in detail. The results show that the IPB consists of Al2O3 nanoparticles, CuAl2 phaseand residual nanopores. After T6 aging, coarse CuAl2 phases with a diameter of 150−600 nm precipitate at the IPB, andthe precipitation free zone (PFZ) with a width of 40−60 nm is formed. An improvement of yield strength and ultimatetensile strength from 95 MPa and 229 MPa to 280 MPa and 378 MPa is achieved respectively after T6 aging, while theelongation to fracture decreases from 11.8% to 6%. The increase in strength is mainly due to the well dispersion ofprecipitates and the densification of the material during T6 aging. The decrease in plasticity may result from the earlierplastic deformation in the PFZ during tensile deformation, leading to the accumulation of dislocations from PFZ to CuAl2phase nearby the IPB, as a result, stress concentrationis formed, which consequently promotes cracksexp and along theIPB and decreases the ductility of the material.Keywords: powder metallurgy; interparticle boundary; precipitation; microstructure; mechanical properties基金项目：国家自然科学基金资助项目(51971143)收稿日期：2021−10−25；修订日期：2021−12−01通信作者：梁加淼，工程师，博士。

B33 T型通道挤压变形ZK60镁合金的组织与力学性能孔晶侯文婷彭勇辉康志新李元元（华南理工大学机械与汽车工程学院 国家金属材料近净成形工程技术研究中心，广东省广州市 510640）摘要：采用一种新型剧塑性变形工艺—T型通道挤压 (TCP) 对ZK60镁合金在673 K温度下以A和Bc两种路径进行1~4道次挤压变形，通过光学显微镜观察了变形镁合金的显微组织。

结果表明，经两种路径TCP变形后晶粒尺寸均明显细化，其中1道次变形后变形过程不均匀，变形量最大部位为试样中间部位的最底部，组织特征为大晶粒和细小晶粒的混合体，大晶粒呈拉长的流线状；随着道次的增加，由于变形过程发生动态再结晶，晶粒不断细化，经4道次变形后试样底部的组织细小均匀，A路径的平均晶粒尺寸由原始铸态的88.5 μm可最小细化至2.4 μm，Bc路径的平均晶粒尺寸则细化至4.6 μm。

对TCP变形镁合金的不同部位以应变速率4×10-3s-1的条件进行室温拉伸，结果表明变形后强度与塑性都得到提高，在相同道次TCP变形后A路径的屈服强度都优于Bc路径，但抗拉强度和塑性却弱于后者，其中以A路径4道次TCP变形后抗拉强度、屈服强度和伸长率分别为305.1 MPa、223.4 MPa和16.4 %，Bc路径分别为312.3 MPa、194.6 MPa和24.8 %；此外，试样最底部的抗拉强度和屈服强度均高于顶部，以Bc路径经2道次变形后底部与顶部的抗拉强度与屈服强度分别相差31.8和39.2 MPa；随着道次增加，试样顶部与底部的变形趋于均匀，在4道次变形后抗拉强度和屈服强度分别只相差3.1和4.6 MPa。

关键词：镁合金；T型通道挤压；剧塑性变形；显微组织；力学性能Microstructure and Mechanical Properties of ZK60 MagnesiumAlloy Processed by T-shape Channel PressingKONG Jing, HOU Wenting, PENG Yonghui,KANG Zhixin, LI Yuanyuan(National Engineering Research Center of Near-Net-Shape Forming for Metallic Materials,School of Mechanical & Automotive Engineering, South China University of Technology,Guangzhou 510640, China)Abstract: ZK60 magnesium alloy was deformed by a new process of severe plastic deformation (SPD) —T-shape channel pressing (TCP) from 1 pass to 4 passes at 673 K using route A and route B c.Microstructure was observed through optical microscope. The experimental results show that TCPed grain size is greatly refined after two routes. Deformation process is heterogeneous after 1 pass only, the biggest deformation is located at the bottom of sample, and microstructure character is combination of large and small grains which large grains are elongated with the shape of streamline. As the passes increased, the grain was refined gradually due to the dynamic recrystallization during deformation. The grains of the bottom are基金项目：广州市科技支撑计划资助项目（2009Z2-D811）作者简介：孔晶，女，在读硕士，师承康志新教授，从事高性能镁合金研究；E-mail：******************。

International Journal of Minerals, Metallurgy and Materials Volume 25, Number 9, September 2018, Page 1080https:///10.1007/s12613-018-1659-7Corresponding author: Hai-tao Zhang E-mail:haitao_zhang@© University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018The evolution of microstructure and mechanical properties duringhigh-speed direct-chill casting in different Al–Mg2Si in situ compositesDong-tao Wang, Hai-tao Zhang, Lei Li, Hai-lin Wu, Ke Qin, and Jian-zhong CuiKey Laboratory of Electromagnetic Processing of Materials, Ministry of Education, Northeastern University, Shenyang 110819, China(Received: 1 November 2017; revised: 9 May 2018; accepted: 10 May 2018)Abstract: The effect of high-speed direct-chill (DC) casting on the microstructure and mechanical properties of Al–Mg2Si in situ composites and AA6061 alloy was investigated. The microstructural evolution of the Al–Mg2Si composites and AA6061 alloy was examined by optical microscopy, field-emission scanning electron microscopy (FE-SEM) and transmission electron microscopy (TEM). The results revealed that an increase of the casting speed substantially refined the primary Mg2Si particles (from 28 to 12 μm), the spacing of eutectic Mg2Si (from 3 to 0.5 μm), and the grains of AA6061 alloy (from 102 to 22 μm). The morphology of the eutectic Mg2Si transformed from lamellar to rod-like and fibrous with increasing casting speed. The tensile tests showed that the yield strength, tensile strength, and elongation improved at higher casting speeds because of refinement of the Mg2Si phase and the grains in the Al–Mg2Si composites and the AA6061 alloy. High-speed DC casting is demonstrated to be an effective method to improve the mechanical properties of Al–Mg2Si composites and AA6061 alloy billets.Keywords: Al–Mg2Si in situ composite; casting speed; grain size; primary Mg2Si; mechanical property1. IntroductionAl–Mg2Si in situ composites have potential applications in the aviation and automotive industries [1–6]. These com-posites offer numerous attractive advantages, including a high elasticity modulus, low density, good wear resistance, and a low thermal expansion coefficient, because of the bene-ficial properties of the Mg2Si intermetallic compound [7–13]. In as-cast Al–Mg2Si composites, the primary Mg2Si particles are coarse and exhibit an irregular morphology [14–19], which is detrimental to the composites’ mechanical proper-ties, thereby limiting their range of applications. Therefore, improving the microstructure and mechanical properties remains a critical issue in the further development of Al–Mg2Si composites. Moreover, Al–Mg–Si alloys take a large proportion of the total aluminum alloys, which are used as structural materials. As a conventional Al–Mg–Si alloy, AA6061 alloy exhibits favorable forming properties, high strength, high corrosion resistance, and good welding performance [20–22]. Therefore, Al–Mg–Si alloy has been widely used in industrial and construction applications. However, the coarse intermetallic phase and low solid solu-bility limit the mechanical properties of AA6061 billet be-cause of the low cooling rate in conventional DC casting.As an important method for producing aluminum alloy billets, DC casting can be used to produce aluminum alloy billets in large quantities. DC casting is important for indus-trial applications of aluminum alloys. However, the conven-tional DC casting process has many disadvantages, includ-ing a low cooling rate, coarse precipitated phase, low solid solubility, and slow melt flow; these shortcomings result in a coarse microstructure and poor mechanical properties of the cast aluminum alloy billets.In the DC casting process, the casting speed can affect the cooling rate, geometry of the liquid sump, and the melt flow. Increasing the casting speed can refine the grain and intermetallic phase, increase the solid solubility, decrease the thickness of the segregation layer on the billet surface, and improve the surface quality of the billet [23–24]. These are very advantageous for the industrial production and ap-D.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting (1081)plication of aluminum alloys. Generally, researchers have focused on the effect of casting speed (100–200 mm·min–1) on the microstructures of aluminum alloys [25–28]. Howev-er, the low casting speed is not sufficient to substantially in-crease the cooling rate, resulting in a limited refinement ef-fect on the grain and intermetallic phase in aluminum alloys.In the present work, we developed a high-speed (300 mm·min–1) DC casting process to improve the microstruc-ture and mechanical properties of Al–Mg2Si in situ compo-sites and AA6061 alloy. Moreover, in order to satisfy the strong cooling demand in a high-speed DC casting process, we improved the design of the cooling water system. The aim of this paper is to investigate the effects of high-speed DC casting on the microstructure and mechanical properties of Al–Mg2Si composites and AA6061 alloy.2. ExperimentalFig. 1 shows a schematic of the high-speed DC casting experiment. First, the starting block was positioned in the copper mold. The melt was poured into the copper mold, which formed a solid shell immediately in the primary cooling region. The starting block was then steadily with-drawn from the copper mold. Finally, the billet surface was directly cooled by water jets (secondary cooling) to achieve the DC casting process. In the conventional DC casting process, cooling water is jetted onto the billet surface by only single-row nozzles, which is insufficient to provide strong cooling during the high-speed DC casting process and results in the melt breaking out. We designed an im-proved secondary cooling system with the nozzles of three rows (Positions 1-3 in Fig. 1). This multiple cooling water system improved both the availability of cooling water and heat transfer and provided sufficient cooling during the high-speed DC casting process.Fig. 1. Schematic of the high-speed direct chill casting process.The chemical compositions of different alloys are listed in Table 1. Commercial pure Al, pure Mg, and Al–20wt%Si master alloy were used to prepare the Al–16%Mg2Si and Al–8%Mg2Si composites and the AA6061 alloy in a 100-kW resistance furnace. When the pure Al and Al-20wt%Si were melted, pure Mg was added to the melt at 650–670°C. The melt was subsequently heated to 800°C and maintained at this temperature for 20 min. The melt was degassed with hexachloroethane dry tablets (0.5wt% of the molten alloy) at 780°C for 6 min, following slag removal. After being held for 15 min, the melt was stirred manually for 3 min to en-sure thorough mixing. Finally, the melt was poured into a copper mold to start the DC casting process. The casting speed was varied in the range 50–300 mm·min–1. Billets with a diameter of φ106 mm were successfully prepared via the high-speed DC casting process.Table 1. Chemical compositions of the alloys wt% Alloy Mg SiFeCuCr Zn MnAl–16%Mg2Si10.11 5.890.07 0.03 0.020.030.06 Al–8%Mg2Si 5.05 2.950.07 0.03 0.020.030.06AA6061 1.020.610.520.310.250.250.15Specimens were cut from the same positions of the soli-dified billets for microstructural observation. The micro-structures were observed on a Leica optical microscope after being etched with 1vol% HF solution. The morphologies of the primary and eutectic Mg2Si crystals were examined on a Zeiss ULTRA PLUS field-emission scanning electron mi-croscope after the samples were deeply etched with 20vol% NaOH solution for 20 s at room temperature. The ImagePro Plus software was used to analyze the size and the area frac-tion of primary Mg2Si particles, the area fraction and the spacing of eutectic Mg2Si, and the grain size and area frac-tion of the intermetallic phase. The eutectic Mg2Si phases were observed on a FEI TECNAI G220 transmission elec-tron microscope operated at 200 kV.The tensile test samples were machined according to standard ASTM B557M. The tensile test was conducted on a SHIMADZU AG-X100 kN tension machine at a cross-head speed of 1 mm·min–1. Each tensile result was the average value of eight tension specimens. The fracture surfaces of the tensile test specimens were observed by field-emission scanning electron microscopy (FE-SEM).3. Results and discussion3.1. MicrostructuresFigs. 2(a)–2(d) shows the microstructural evolution of1082 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018the Al–16%Mg 2Si composite at different casting speeds. The increased casting speed effectively refines the primary Mg 2Si; the morphology of the primary Mg 2Si transforms from an irregular to a regular polygon shape. As evident in Figs. 2(e)–2(h), the increase of casting speed substantial-ly refined the eutectic Mg 2Si in the Al–8%Mg 2Si compo-site. In the Al–Mg 2Si composites, the eutectic Mg 2Si ex-hibits a lamellar morphology and a large eutectic spacing when cast at low speed (Figs. 2(i) and 2(k)). The increase of the casting speed refines the eutectic Mg 2Si and sub-stantially decreases the eutectic spacing (Figs. 2(j) and2(l)).Fig. 2. Microstructures of Al–Mg–Si alloys at different DC casting speeds: Al–16%Mg 2Si cast at (a) 50 mm·min –1, (b) 100 mm·min –1, (c) 200 mm·min –1, and (d) 300 mm·min –1; eutectic morphology of Al–16%Mg 2Si cast at (i) 50 mm·min –1 and (j) 300 mm·min –1; Al–8%Mg 2Si cast at (e) 50 mm·min –1, (f) 100 mm·min –1, (g) 200 mm·min –1, and (h) 300 mm·min –1; eutectic morphology of Al–8%Mg 2Si cast at (k) 50 mm·min –1 and (l) 300 mm·min –1.D.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting …1083Figs. 3(a)–3(h) shows the microstructural evolution of the AA6061 alloy cast at different speeds. The Mg 2Si phase was refined and transformed from bulk and strip-like shapes to dot-like shapes with increasing casting speed; the reticular α-AlFeSi phase was also substantially refined at high casting speeds. Figs. 3(b) and 3(g) shows the evolution of the grain structures of the AA6061 alloy specimens. The grain size ofthe AA6061 alloy decreases with the increase of casting speed.Fig. 3. Microstructures of AA6061 alloy at different casting speeds: (a) optical image, (b) grains, and (c) SEM image of AA6061 al-loy cast at 50 mm·min –1; optical images of the AA6061 alloy cast at (d) 100 mm·min –1 and (e) 200 mm·min –1; (f) optical image, (g) grains, and (h) SEM image of AA6061 alloy at 300 mm·min –1.Figs. 4(a)–4(c) shows the area fraction and the size of the primary Mg 2Si at the different casting speeds. With increas-ing casting speed, the area fraction of the primary Mg 2Si decreases from 6.9% (50 mm·min –1) to 4% (300 mm·min –1)1084 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018and the size of the primary Mg 2Si decreases from 28 (50 mm·min –1) to 11 μm (300 mm·min –1). Moreover, the spacing of the eutectic Mg 2Si decreases from 2.7 (50 mm·min –1) to 0.4 μm (300 mm·min –1) in Fig. 4(d). These results indicate that the high-speed DC casting adequately refines the mi-crostructure of the Al–16%Mg 2Si composite. Figs. 4(a) and 4(d) shows the area fraction of the eutectic Mg 2Si and the spacing of eutectic Mg 2Si in the Al–8%Mg 2Si samples cast at different speeds. Notably, the eutectic spacing decreasesfrom 3.5 (100 mm·min –1) to 0.4 μm (300 mm·min –1). The area fraction of the eutectic Mg 2Si increases from 7.3% to 13.2%. As shown in Figs. 4(a) and 4(b), the grain size of the AA6061 alloy decreased from 105 (50 mm·min –1) to 26 μm (300 mm·min –1) and the area fraction of the intermetallic phases decreased from 5.2% (50 mm·min –1) to 2.8% (300 mm·min –1) with increasing casting speed. These results in-dicate that the high casting speed improves the solid solubil-ity of the different alloy elements in the α-Al solid solution.Fig. 4. Area fraction, grain size and eutectic spacing of the different Al–Mg–Si alloys: (a) area fraction; (b) grain size of the AA6061 alloy; (c) the primary Mg 2Si size; (d) the eutectic spacing as functions of the casting speed.Fig. 5 shows TEM images of the eutectic Mg 2Si of Al–Mg 2Si composites cast at high DC casting speed. Com-bined with the morphology observations of the eutectic Mg 2Si in Figs. 2(j) and 5(a), these results show that the eu-tectic Mg 2Si in the Al–16%Mg 2Si composite exhibits a rod-like and fibrous shape, with a diameter of 200 nm. By contrast, the eutectic Mg 2Si in the Al–8%Mg 2Si composite exhibits a long and lath-shaped morphology, with a width from 100 to 200 nm and length from 2 to 4 μm, as shown in Fig. 5(b). The eutectic Mg 2Si in the different Al–Mg 2Si composites exhibits different morphologies and sizes at the highest casting speed.The growth rate of Mg 2Si crystals differs substantially in different directions because of the faceted growth style. The growth of Mg 2Si crystals is slower than that of α-Al. Theα-Al can surround the eutectic Mg 2Si in the growth process. Therefore, the continual growth of Mg 2Si crystals depends on re-nucleation. The increased cooling rate promotes nucleation and growth of the eutectic and thus results in the observed fine eutectic structure with rod-like and fibrous shapes. 3.2. Mechanical propertiesFig. 6(a) shows the yield strength of the different alloys cast at different speeds. The yield strengths of the different alloys all improve with increasing casting speed. Notably, the Al–8%Mg 2Si composite exhibits higher yield strength than the Al–16%Mg 2Si composite and the AA6061 alloy. High casting speeds effectively increase the cooling rateD.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting (1085)during the DC casting process, resulting in a finer Mg 2Si phase in the Al–Mg 2Si composites and finer grains in the AA6061 alloy. The refinement of grains and the Mg 2Si phase causes an increase in the number of grain boundaries, which can act as obstacles to dislocation motion [29–30]. The increase in yield strength due to grain boundaries, ΔσGB ,is described by the Hall–Petch equation [30]:12GB y k d σ-∆= (1) where d is the average grain diameter and k y is the Hall–Petch slope. Decreasing grain size and refining the Mg 2Si phase improve the yield strength of the different al-loys via the augmentation of grain boundaries.Fig. 5. TEM images of the eutectic Mg 2Si of the Al–Mg 2Si in situ composites cast at high DC casting speed (300 mm·min –1): (a) Al–16%Mg 2Si; (b) Al–8%Mg 2Si.Figs. 6(b) and 6(c) shows the tensile strength and the elongation at the different casting speeds. Both the tensile strength and the elongation improve with increasing casting speed in the different alloys. The fine and regular Mg 2Si par-ticles strengthen the load-bearing capacity and suppress crack propagation along the particles, which results in enhance-ment of the strength and the ductility of the Al–16%Mg 2Si composite. Moreover, the refinement of grains and the re-duction of intermetallic phases results in increases of thetensile strength and the elongation in the AA6061 alloy.The Al–16%Mg 2Si composite cast at the highest casting speed (300 mm·min –1) exhibits lower elongation (4.33%) than the AA6061 alloy (18.85%) because of the brittle Mg 2Si particles. The Al–8%Mg 2Si composite exhibits theFig. 6. Tensile properties of different Al–Mg–Si alloys: (a) yield strength; (b) tensile strength; (c) elongation.1086 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018highest yield and tensile strength; it shows the lowest elon-gation (0.82%) because of its high area fraction and long lath shape of eutectic Mg 2Si (Fig. 5(b)). Fig. 7 shows the engineering stress–strain curves of the different alloys. The Al–8%Mg 2Si composite does not exhibit the evident plastic deformation stage during the tensile process, consistent with its low elongation. The engineering stress–strain curves of the Al–16%Mg 2Si composite show some extent of plastic deformation; thus, its elongation increases in samples cast at high speeds. The AA6061 alloy undergoes the obvious plas-tic deformation stage and greater elongation with increasing casting speed. During the high-speed DC casting process,the high cooling rate is the decisive factor for improving themicrostructure and mechanical performance.Dislocation theory can explain the improvement of ten-sile strength [29–30]. The process of plastic deformation can yield a large number of moving dislocations. In the Al–Mg 2Si composites, the grain boundary of the Mg 2Si phase hinders the dislocation glide and leads to dislocation accumulation. Substantial dislocation accumulation will produce the driving force for dislocation glide. With in-creasing Mg 2Si crystal size, an increase in dislocation ac-cumulation implies a higher driving force of dislocation glide. The accumulated dislocations are easier to glide, re-sulting in a diminished strengthening effect. Therefore, finer Mg 2Si phases result in less dislocation accumulation and a weaker driving force of dislocation glide; it blocks disloca-tion glide during the deformation process and thereby im-proves the tensile strength. Moreover, solution strengthening is also a factor responsible for tensile-strength improvement in the AA6061 alloy because of the increased solid solubili-ty of the different alloying elements. 3.3. FractographyFigs. 8(a)–8(f) shows the fracture surfaces of the different alloys cast at different speeds. Fig. 8(a) shows a typical fracture surface of the Al–16%Mg 2Si composite cast at low speed. The fracture surface includes the complete Mg 2Si particles, which implies that the fracture occurs at the inter-face between the particles and the matrix. In conventional DC casting, coarse and irregular Mg 2Si particles result in an increase in both the stress concentration and crack initiation. The cracks often initiate at susceptible and weak points (i.e., coarse and irregular Mg 2Si particles) along the interface between the matrix and the particle. The particles are diffi-cult to bear higher local active stress in comparison with their intrinsic yield strength. Therefore, the coarse Mg 2Si particles break the continuity of the Al matrix and decreaseFig. 7. Engineering stress–strain curves for specimens cast at different speeds (as cast): (a)Al–16%Mg 2Si; (b) Al–8%Mg 2Si; (c) AA6061.D.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting (1087)the ductility of the composite. Moreover, the lamellar and coarse eutectic structure encourages the propagation of cracks and thus further restricts the tensile properties of the Al–16%Mg 2Si composite.Fig. 8. Fracture surfaces of specimens cast at different DC casting speeds: Al–16%Mg 2Si cast at (a) 100 mm·min –1 and (b) 300 mm·min –1; Al–8%Mg 2Si cast at (c) 100 mm·min –1 and (d) 300 mm·min –1; and AA6061 cast at (e) 100 mm·min –1 and (f) 300 mm·min –1.The refinement of Mg 2Si particles enhances both the continuity of the Al matrix and the load-bearing ability of the Mg 2Si particles. When the interface-bearing stress is higher than the intrinsic yield strength of Mg 2Si particles, internal cracks occur at the Mg 2Si particles, as shown in Fig. 8(b). Moreover, the fine dimples were observed in the Al matrix, which enhances the ductility of the composite. In this case, the fracture mechanism is both brittle fracture and small extent of ductile fracture. Fine, regular, and homoge-neous Mg 2Si phases can restrain crack initiation to extend along the Mg 2Si particles and strengthen the cohesion of the α-Al matrix.Fig. 8(c) shows the fracture surface of an Al–8%Mg 2Si composite cast at low speed; it exhibits clear cleavage cha-racteristics derived from its intrinsic brittleness. This re-markable brittle fracture results in this sample exhibiting the lowest elongation among the investigated specimens. With increasing casting speed, the fracture surface did not exhibit an obvious cleavage plane in Fig. 8(d); the elongation slightly increased. The increase of the casting speed had no evident effect on the transformation of the fracture characteristic. Fig. 8(e) shows a typical fracture surface of the AA6061 alloy. The excellent ductility of the alloy coincides with the formation of substantial dimples on the formed separation surface. With increasing casting speed, the elongation fur-ther increases because of an increase in the number of dim-ples and the refinement of the dimples, as shown in Fig. 8(f).1088 Int. J. Miner. Metall. Mater., Vol. 25, No. 9, Sep. 20184. ConclusionsThe effect of high-speed DC casting on microstructures and mechanical properties of different Al–Mg–Si alloys was studied. The following conclusions were drawn from the results of the present investigation:(1) A high DC casting speed substantially refines the primary Mg2Si particles and the eutectic Mg2Si structure of the Al–Mg2Si in situ composites. It also effectively decreas-es the grain size and refines the intermetallic phases of AA6061 alloy. With increasing casting speed, the primary Mg2Si transforms from an irregular to a polygonal mor-phology; the eutectic Mg2Si changes from lamellar to rod-like and fibrous morphologies. The morphology of eu-tectic Mg2Si shows the difference in the different Al–Mg2Si composites when the billets were cast at the highest casting speed.(2) The microstructural improvements, including the re-finement of theMg2Si phase, the morphology transformation of the Mg2Si phase, and the decrease of grain size, effec-tively strengthen the mechanical properties of the Al–Mg2Si composites and the AA6061 alloy. 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基金项目：湖南省自然科学基金资助项目（2019JJ40245）；南华大学核燃料循环技术与装备湖南省协同创新中心开放基金项目收稿日期：2020－12－03B 4C/Al 复合材料的组织、力学性能和制备研究进展*石细桥，柏兴旺※，俞雪奇，何鹏（南华大学机械工程学院，湖南衡阳421001）摘要：B 4C 颗粒增强铝基复合材料不仅比强度高、耐磨性能优异，还兼具多种功能特性，是核工业、航空航天、汽车工业等领域中不可缺少的结构材料和功能材料。

综述了B 4C/Al 复合材料的研究现状，总结了搅拌铸造法、粉末冶金法和冷喷涂增材制造法等各种方法的优缺点，对比了不同工艺下制备的B 4C 颗粒增强铝基复合材料的硬度、拉伸强度、屈服强度、抗压强度等力学性能方面及显微组织的表现，并展望了其发展方向。

关键词：B 4C/Al 复合材料；显微组织；力学性能中图分类号：TB333文献标志码：A文章编号：1009－9492(2021)03－0076－03开放科学（资源服务）标识码（OSID ）：Study Progress on the Microstructure,Fabrication and Mechanical Properties ofB 4C/Al CompositesShi Xiqiao ，Bai Xingwang ※，Yu Xueqi ，He Peng（School of Mechanical Engineer,University of South China,Hengyang,Hunan 421001,China ）Abstract:B 4C particle-reinforced aluminum-based composite materials not only have high specific strength,excellent wear resistance,but also have a variety offunctional properties,which is indispensable structural material and functional material in the nuclear industry,aerospace,automobile industry and other fields.The research status of B 4C/Al composite materials was reviewed,the advantages and disadvantages of various methods such as stirred casting method,powder metallurgy method and cold spray additive manufacturing method were summarized,and the mechanical properties such as hardness,tensile strength,yieldstrength,compressive strength and microstructure of B 4C particle reinforced aluminum matrix composites materials prepared under different processes were compared,and its development direction was prospected.Key words:B 4C/Al composite;microstructure;mechanical properties第50卷第03期Vol.50No.03机电工程技术MECHANICAL &ELECTRICAL ENGINEERING TECHNOLOGYDOI:10.3969/j.issn.1009-9492.2021.03.015石细桥，柏兴旺，俞雪奇，等.B4C/Al 复合材料的组织、力学性能和制备研究进展［J ］.机电工程技术，2021，50（03）：76-78.0引言B 4C 颗粒增强铝基复合材料（AMC ），因其具有高比强度、高比刚度、良好的力学性能以及出色的导热性和化学稳定性[1-3]，成为航空航天、汽车工业等领域中不可缺少的结构材料和功能材料。

Microstructure and mechanical properties of twinned Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy processed by mechanical alloying and spark plasmasinteringSicong Fang,Weiping Chen,Zhiqiang Fu ⇑School of Mechanical and Automotive Engineering,South China University of Technology,Guangzhou,Guangdong 510640,Chinaa r t i c l e i n f o Article history:Received 9July 2013Accepted 29August 2013Available online 7September 2013Keywords:High entropy alloys Mechanical alloying Spark plasma sintering Nanoscale twins Microstructurea b s t r a c tMost of multi-component high entropy alloys (HEAs)only consist of metallic elements.In the present paper,by introducing nonmetallic carbon element,non-equiatomic Al 0.5CrFeNiCo 0.3C 0.2HEA has been successfully prepared by mechanical alloying (MA)and spark plasma sintering (SPS)process.Alloying behavior,microstructure,phase evolution and mechanical properties of the alloy were investigated systematically.During the MA process,a supersaturated solid solution with both face-center cubic (FCC)and body-center cubic (BCC)structures was formed within 38h of milling.However,a major FCC phase,a BCC phase,Cr 23C 6carbide and an ordered BCC phase were observed after SPS.The FCC phase is enriched in Fe–Ni,the BCC phase is enriched in Ni–Al and the ordered BCC phase is especially enriched in Al,respectively.In addition,nanoscale deformation twins obviously presented only in partial FCC phase after SPS.The compressive strength and Vickers hardness of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy are 2131MPa and 617±25HV,respectively.Ó2013Elsevier Ltd.All rights reserved.1.IntroductionTraditional alloys are typically composed of one principal ele-ment that occupies at least 50at.%in the composition,with minor additions of other elements to obtain definite microstructure and properties,such as Fe-,Al-,Cu-and Mg-based alloys [1].However,high entropy alloys (HEAs)that proposed by Yeh et al.in 2004have broken this conventional paradigm [2,3].This kind of alloys are de-fined as alloys which consist of at least five principal elements,and the concentration of each constituent element ranges from 5to 35at.%.Previous researches show that HEAs can be processed to form simple solid solution structures instead of intermetallics and other complicated compounds [4–6].This phenomenon is commonly attributed to the high configurational entropy in the solid solution state of HEAs [2,7].Furthermore,HEAs have also exhibited interesting properties such as high hardness and high strength [8,9],good thermal stability [10],outstanding wear and oxidation resistance [3,11],which offer great potential for engi-neering applications.The HEA systems explored in the past decade show that metal-lic elements are the most commonly used,e.g.Al,Cr,Fe,Co,Ni,Cu,Ti,etc.[12–14].It is known that the proper addition of nonmetallic elements like C in some traditional alloys is favorable to theirstructures and properties [15,16].However,to the best of our knowledge,HEAs with addition of C element have been rarely investigated and reported.In contrast,AlCrFeNiCo HEA system pre-pared by arc-melting has been extensively studied in existing liter-atures [5,17,18].Wang et al.[19]have investigated microstructure and mechanical properties of Al x CrFeNiCo (06x 62)HEAs,finding that the as-cast Al x CrFeNiCo alloys can possess FCC and/or BCC structure(s)depending on the aluminum content.Increasing con-centration of Al in this alloy system can lead to the formation of BCC phase,which possess high strength and high hardness while inferior plasticity.Apparently,the concentration of Al element should be in an appropriate range to achieve optimum properties.Hence,considering all the factors discussed above,Al 0.5CrFeNiC-o 0.3C 0.2HEA was designed.Moreover,the most widely studied processing route for HEAs is arc-melting (casting),and only a few reports deal with mechanical alloying (MA)[20–23].As a widely used solid state processing route,MA can easily fabricate nanocrystalline materials with good homogeneity from elemental powders,thus MA can be an ideal way to prepare HEAs [20,24,25].In addition,SPS can rapidly con-solidate alloy powders to high density by applying pressure and passing electric pulse current within short soaking time [26].HEAs synthesized by MA and SPS have been reported to possess good densification characteristics,as well as high strength and high hardness [27–29].Hence,MA and SPS were combined to prepare Al 0.5CrFeNiCo 0.3C 0.2HEA.Alloying behavior and phase evolution0261-3069/$-see front matter Ó2013Elsevier Ltd.All rights reserved./10.1016/j.matdes.2013.08.099Corresponding author.Tel.:+862087113832;fax:+862087112111.E-mail addresses:kopyhit@ ,fu.zhiqiang@ (Z.Fu).during milling and consolidation processesied.Microstructure and mechanical properties ofwere investigated.2.Experiment proceduresHEA with a nominal composition of Al0.5pressed in molar ratio)was prepared by drythen by wet milling for4h in ethanol.Al,Cr,Fe,elemental powders with purity higher than99.9size of645l m(325mesh)were mechanicallywas carried out in a high energy planetaryPlanetary Ball Mill)at300rpm with a ball toof10:1under argon atmosphere.Highvials and tungsten carbide balls(10mm inthe milling media.In order to confirm the alloymilling,powder samples were taken out after27,38,42h respectively.The42h ball milledsubsequently sintered by Dr.Sinter Model SPS-825Spark PlasmaSintering System(Sumitomo Coal Mining Co.Ltd.,Japan)at 1273K for8min at a pressure of30MPa under vacuum(residual cell pressure<8Pa).The samples were heated to873K within 4min,while from873K to1173K and from1173K to1273K, heating rates of75K minÀ1and50K minÀ1were used, respectively.The milled powders and the bulk alloy after sintering were ana-lyzed by a Bruker D8ADVANCE X-ray diffractometer(XRD)with a Cu K a radiation.The microstructure of the alloy was revealed by etching in aqua regia and observed using scanning electron micros-copy(SEM,Zeiss Supra40,Carl Zeiss NTS GmbH,Germany).Thin-foil specimens were prepared by mechanical thinning followed by ion milling at room temperature and were analyzed by a transmis-sion electron microscopy(200kV TEM,JEM-2100,JEOL,Japan) with selected area electron diffraction(SAED)analysis.According to GB/T7314-2005[30],the room-temperature compressive prop-erties of the cylindrical samples(U3mmÂ4.5mm in size)were measured with an Instron5500testing system at a strain rate of 1Â10À3sÀ1.Three compression tests were performed to obtain average value.Hardness measurement was conducted using a Dig-ital Micro Hardness Tester HVS-1000Vickers hardness instrument under a load of300gf.The reported hardness value is an average of at least10measurements.3.Results and discussion3.1.XRD analysis and microstructure during MAFig.1shows the XRD patterns of Al0.5CrFeNiCo0.3C0.2HEA pow-ders prepared under different milling durations.It can be seen that diffraction patterns of all alloy elements are included in the initial blending powder.As the milling time increases,drastic decrement of diffraction intensity is observed after6h of milling.The peaks of Al,Co and C elements dissolve most rapidly.According to Chen et al.[31],the alloying sequence for the Cu0.5NiAlCoCrFeTiMo alloy system correlates best with the melting point of the component elements.Thus rapid dissolution of both Al and Co could be asso-ciated with their lower melting point than other elements.In con-trast,the early disappearance of diffraction peak of C may result from its smallest atomic fraction in the alloy.Many diffraction peaks can hardly be seen when the milling time reaches up to 27h.The complete disappearance of all the elemental peaks and the formation of solid solution are founded within38h of milling. Predominant peaks corresponding to an face-centered cubic(FCC) and a body-centered cubic(BCC)crystal structures are respectively observed.Subsequently,the powder was subjected to wet ball milling in ethanol for4h with the aim to obtainfine metallic pow-der for being conductive to sintering.As the milling time extends to42h,the diffraction peaks exhibit no change except for a minor broadening and a significant increment of diffraction intensity.The crystal size and lattice strain of the BCC and FCC phases with different milling time have been calculated from the X-ray peak broadening using Scherrer’s formula after deducting the instrumental contribution.The calculated results are listed in Table1.The crystalline sizes of both BCC and FCC phases after 42h of milling are slightly refined compared with38h of milling, while the lattice strain of these two phases increase as the milling time prolongs.The crystallite refinement can be attributed to the circulation of crushment and agglomeration during the MA process.Reasons for the increment of lattice strain include size mismatch effect between the constituent elements,increasing grain boundary fraction and high dislocation density imparted by MA[21].The crystallite refinement and high lattice strain may account for the above-mentioned intensity increment and peak broadening in the diffraction.The SEM images of the Al0.5CrFeNiCo0.3C0.2HEA powders of dif-ferent milling time are shown in Fig.2.The primitive powder shows a granular size of less than40l m.In the early period of MA(as shown in Fig.2(b)),the particles cold weld together to form even larger particles than that of primitive powder.Subsequently, when the milling time reaches15h(Fig.2(c)),most of the cold welded agglomerations are crushed down to smaller particles. The27h milled alloy powder reveals an average particle size of approximately5l m as shown in Fig.2(d),and the particles cold weld again when the milling time reaches38h(Fig.2(e)).This cir-culation of crushing and cold welding induced by the ball mill gradually reduces the crystalline size and facilitates the diffusion and alloying among different elements.The38h milled alloy pow-der is subsequently wet ball milled for4h with alcohol as the mill-ing media.It can be seen from Fig.2(f)that the elliptoid particles are fractured,exhibiting a lamellar structure.The particle size for thefinal powder is much smaller than that of38h milled powder.1.XRD patterns of Al0.5CrFeNiCo0.3C0.2HEA powders with different milling time.Table1The crystalline size and lattice strain of Al0.5CrFeNiCo0.3C0.2HEA during MA.Milling time(h)Crystalline size(nm)Lattice strain(%)BCC FCC BCC FCC 3812130.710.734211120.760.77974S.3.2.Phase evolution and microstructure after SPSFig.3illustrates the XRD pattern of Al0.5CrFeNiCo 0.3C 0.2HEA detailed analysis of the XRD pattern suggests a BCC phase,a Cr 23C 6carbide phase and an visible.Except the ordered BCC phase,the other three phases are calculated to be (BCC),10.652Å(Cr 23C 6),respectively.After the FCC phase reduces from 0.77%after MA phase from 0.76%to 0.27%,which confirmsthe nearly annihilation of defects introduced during MA after sin-tering.As hereinbefore mentioned,the main phases of Al 0.5CrFe-NiCo 0.3C 0.2HEA powder after 42h of milling are the BCC and FCC solid-solution phases,demonstrating that densification at 1273K for 8min has resulted in phase evolution.The MA process could in-duce large strain and defects which might lead to the extension of solubility,thus the milled powders are generally in a non-equilib-rium state.It is certain that a reordering process happened during SPS,leading to the metastable supersaturated solid solutions of MA to more stable phases.This reordering can be correlated with the above-mentioned annihilation of defects introduced by severe plastic deformation during MA.Furthermore,the huge pulsed elec-tric current during the SPS process can also facilitate the phase evolution.Fig.4shows the SEM micrographs of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA consolidated by SPS.Few porosities can be seen in low magni-fied image (Fig.4(a)).Actually,relative density of the sintered Al 0.5-CrFeNiCo 0.3C 0.2HEA sample,which is calculated with respect to theoretical density,reaches up to 99.6%.Two distinctive areas are visible in Fig.4(a),viz.,irregular bulk areas and irregular pot hole areas.The irregular bulk areas are most likely to be the main phase because they account for a higher volume fraction.While the pot hole areas can be consist of phases which are removed by corrosive agent.High magnified image (Fig.4(b))shows that some nanoscale black spots are dispersed in irregular pot hole areas.Most of these spots might be numerous nanoscale phases,and a small minority of spots might be ultrafine porosities.In summary,bulk Al 0.5CrFe-NiCo 0.3C 0.2HEA might consist of at least three types of phases.The TEM bright field image and corresponding selected area electron diffraction (SAED)patterns of Al 0.5CrFeNiCo 0.3C 0.2HEAAl 0.5CrFeNiCo 0.3C 0.2HEA powders with different milling time.(a)0h,(b)6h,(c)15h,(d)27h,(e)38h and (f)Fig.3.XRD pattern of Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.after SPS are shown in Fig.5.In order to confirm the phase compo-sition of Al 0.5CrFeNiCo 0.3C 0.2HEA,EDS/TEM method was used to measure chemical compositions (in at.%)of the phases in Fig.5(a)(regions marked A–Y).Results of the chemical composi-tion analysis are listed in Table 2.It is noticed that all the regions can be classified as four different phases.Crystal structures of these phases are measured by their corresponding SAED patterns as shown from Fig.5(b–e).As a result,these four phases are measured to be an FCC phase,a BCC phase,Cr 23C 6carbide and an ordered BCC phase,respectively.It is worth mentioning that diffraction pattern of region C exhibits a FCC twin structure of the FCC phase.Clear microstructures and corresponding SAED patterns of twins will be illustrated hereinafter.As shown in Table 2,the FCC phase isdepleted in Al,Cr compared with the nominal CrFeNiCo 0.3C 0.2HEA.The BCC phase is Ni-rich,Fe-depleted and C-depleted,while the or-greatly rich in Al.The lattice parameters of measured by SAED are 3.700Å(FCC),Å(Cr 23C 6)and 4.320Å(Ordered BCC),Cr 23C 6phase also presents Fe,Ni,Co and Al elements,especially Fe shows a high concentration (12.7±0.4at.%).In conclusion,the bulk Al 0.5CrFeNiCo 0.3C 0.2HEA exhibits a Fe–Ni-rich FCC phase,a Ni ÀAl-rich BCC phase,a Cr 23C 6carbide phase and a Al-rich ordered BCC phase,which is in accordance with the XRD result (Fig.3).It is noticeable that the grain size of these phases shows a wide distri-bution ranging from several hundred nanometers to 1l m.How-ever,the Cr 23C 6and the ordered BCC phases are much finer than the FCC and the BCC phases.It is a reasonable explanation that the Cr 23C 6and the ordered BCC phases are formed in a phase evo-lution and reordering process during SPS.The formation of these four phases is complicated.According to the Gibbs phase rules,the number of equilibrium phases (p )is p =n +1for the alloy that contains n elements.Since phase formation not in equilibrium conditions,the number p >n +1.However,the Al 0.5CrFeNiCo 0.3C 0.2four types of phases,showing much fewer Yeh et al.[2]has proposed that it is possible solutions when HEAs contain five or more 4.SEM micrographs of Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.(a)low-magnified image,and (b)high-magnified image.corresponding SAED patterns of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.(a)bright field image,(b)SAED pattern of FCC [011]1]zone axis (Region D),(d)SAED pattern of Cr 23C 6[112]zone axis (Region B),and (e)SAED pattern of ordered BCC [012]Table 2Chemical composition (in at.%)analysis results of the phases by EDS/TEM.RegionsPhases Cr Fe Ni Al Co C A,C,F,G,I,J,K,M,N,O,P,T,V,W,X,Y FCC 21.8±0.931.8±1.229±1.1 4.4±0.68.5±0.8 4.5±0.5D,E,Q,R BCC 19.7±1.218.6±2.032.8±1.217.5±1.87.2±0.5 4.2±0.4B,L Cr 23C 663.5±2.112.7±0.4 2.9±0.4 1.8±0.3 2.6±0.316.5±1.6H,S,UOrdered BCC 1.6±0.5 2.4±0.42±0.486.8±1.7 1.2±0.46±1.0Nominal composition–25252512.57.55are randomly distributed in the crystal lattice,though the alloy exhibits high mixing entropy.It reveals that high mixing entropy is insufficient to dominate the formation of phases in HEAs sys-tems.Zhang et al.[32]related the formation of simple solid solu-tion to the mixing enthalpy (D H mix )and atomic size difference (d ).Subsequently,Yang and Zhang [33]proposed a solid-solution formation rule for multi-component HEAs based on the calculation of most of reported HEAs.According to Yang et al.,two parameters can be used to estimate the phase formation behavior of HEAs:O is defined as a parameter of the entropy of mixing timing the average melting temperature of the elements over the enthalpy of mixing,d is defined as the mean square deviation of the atomic size of ele-ments.It is proposed that HEAs form simple crystal structures when O P 1.1and d 66.6%.These two parameters are defined by Eqs.(1)and (2),respectively.X ¼T m D S mix j D H mix jð1Þd ¼ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiX n i ¼1c i ð1Àr i = r Þ2q ð2Þwhere c i is the atomic percentage of the i th component, r ¼Pn i ¼1c i r i is the average atomic radius and r i is the atomic radius of the i th component.T m ,D S mix and D H mix are calculated as follows:T m ¼Xn i ¼1c i ðT m Þið3ÞHere,(T m )i is the melting point of the i th component of alloy.D S mix¼ÀR Xn i ¼1ðc i ln c i Þð4Þwhere c i is the mole percent of component,P ni ¼1c i¼1,and R(=8.314JK À1mol À1)is gas constant.D H mix ¼X n i ¼1;i –jX ij c i c j ð5Þwhere X ij ¼4D H mix ijis the regular solution interaction parameter between the i th and j th elements,c i or c j is the atomic percentage of the i th or j th component,and D H mix ijis the enthalpy of mixing of binary liquid alloys.Table 3presents the value of mixing enthal-py of atom-pairs of Al 0.5CrFeNiCo 0.3C 0.2HEA.The calculated results of O and d for the alloy are 1.41and 10.01%,respectively.It is obvi-ous that O well matches while d breaks the solid-solution formation rules for the multi-component high entropy alloys proposed by Yang et al.This phenomenon indicates that the formation of Cr 23C 6carbide and ordered BCC phase is reasonable in this alloy sys-tem and it well conforms to Yang’s research.Firstly,it is obvious that the value of O shows the relative predominance of T m D S mix and D H mix .Since O >1is obtained,T m D S mix should be the predomi-nant part of the free energy.It is known that mixing entropy indi-cates the tendency for the formation of random solution while enthalpy of mixing in any system indicates the tendency for order-ing or clustering,thus the formation of solid-solution phases ought to be much easier than the formation of intermetallic compoundsand other ordered phases in the Al 0.5CrFeNiCo 0.3C 0.2HEA.Obviously,the formation of FCC solid-solution phase as the main phase of Al 0.5-CrFeNiCo 0.3C 0.2HEA is mainly correlated with the alloy’s high mix-ing entropy.Secondly,the relatively large value of d indicates that the atomic size difference between components is too large for this alloy system to form entire solid-solution phases.Because the large atomic size mismatch could lead to serious lattice distortion and subsequently increase the corresponding strain energy which could lower the stability of solid-solution.Actually,the interstitial solubil-ity of carbon element in alloy is quite limited,so that carbon ele-ment has a strong tendency to exist in the form of carbonization or graphite in alloy.Moreover,the diffusion of atoms in the matrix could be suppressed due to large atomic size difference.Thus it facilitates the atomic segregation,even results in the formation of amorphous structures [35].As listed in Table 3,atomic radiuses of C (0.77Å)and Al (1.43Å)are significantly different from the other elements.This could be a reasonable explanation for the formation of the Cr 23C 6carbide and the Al-rich ordered BCC phase.In addition,it can also be found that the mixing enthalpy of most of the atomic pairs is highly negative (shown in Table 3).Al and C are the main contributor to the negative enthalpy of mixing in this alloy,indicat-ing their strong tendency for ordering or clustering.It is interesting to note that the mixing enthalpy of C and Cr (À61kJ/mol)is the most negative.This could be another factor which contributes to the formation of the Cr 23C 6carbide in the alloy.In addition,as above-mentioned,Fe has a high concentration (12.7±0.4at.%)in Cr 23C 6phase,revealing that Fe has high solubility in this type Cr 23C 6phase.According to Table 3,the mixing enthalpy of C and Fe (À50kJ/mol)stays a highly negative level,leading to Fe atoms preferring C sites.Mechanically alloyed powders of Al 0.5CrFeNiCo 0.3-C 0.2HEA exhibit simple solid solution structure within 38h of mill-ing,which can be attributed to the formation of supersaturated solid solution.The solid solubility extension can be ascribed to the high mixing entropy effect as well as the non-equilibrium state of the MA process [36].Interestingly,twinned FCC phase is also ob-served.Fig.6shows TEM images and corresponding SAED patterns of the FCC phase with nanoscale twins.It is worth pointing out that nanoscale twins are found only in the FCC phase,which has been confirmed by EDS/TEM and the corresponding diffraction patterns.The lamella thickness of the nanoscale twins shown in Fig.6(a)is less than 60nm.The corresponding SAED pattern (the matrix axisis [011]M and the twin axis is ½0 1 1 T)which is shown in the inset of the upper right of Fig.6(a)indicates the nanoscale twins belong to the FCC phase.Actually,existing researches in deformation twin-ning of nanocrystalline materials are mainly focus on FCC crystal structure metals,which is probably attributed to their favorable capability of deformation twinning [37].Fig.6(b)illustrates a noticeable twin structure of the FCC phase surrounding by grey phases.These grey phases are found to be the Cr 23C 6carbide through EDS/TEM analysis,revealing that the formation of Cr 23C 6might have effect on deformation twinning.Since Cr 23C 6shows a hard brittle texture [38],it could be presumed that partial FCC phase between the Cr 23C 6carbide is not readily deformed and con-solidated.Under the isostatic pressure of 30MPa during the SPS process,certain parallel crystal faces of partial crystals moving opposite to each other along a direction with a certain value of dis-placement distance.Thus twinning in the FCC phase may occur dur-ing the phase evolution and densification with the aim to be more stable and attaining complete densification [29].Moreover,a cer-tain twin system could be activated by a critical resolved shear stress which can be achieved during crystal deformation and phase evolution process.It is worth pointing out that similar nanoscale twins in CoNiFeAl 0.6Ti 0.4and CrCoNiFeAl 0.6Ti 0.4HEAs prepared by MA and SPS have been observed in previous studies of our research group [28,29].Table 3The chemical mixing enthalpy D H mix ij ;kJ =mol of binary equiatomic alloys calculated by Miedema’s approach [34].Element (atomic sizes,Å)C Ni Cr Co Fe Al C (0.77)–À39À61À42À50À36Ni (1.24)––À70À2À22Cr (1.25)–––À4À1À10Co (1.26)––––À1À19Fe (1.26)–––––À11Al (1.43)––––––S.Fang et al./Materials and Design 54(2014)973–9799773.3.Mechanical propertiesFig.7shows the room-temperature compressive stress–strain curve of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.The compressive strength and compression ratio of the alloy are 2131MPa and 3.0%,respectively.The average Vickers hardness of Al 0.5CrFeNiC-o 0.3C 0.2HEA after SPS has been measured to be 617±25HV.How-ever,plastic deformation behavior is not characterized in theindicating a limited compression plasticity of mechanical properties of Al 0.5CrFeNiCo 0.3C 0.2HEA HEAs of AlCrFeNiCo HEA system are listed The listed HEAs are prepared by casting 0.5CrFeNiCo 0.3C 0.2HEA studied in this paper.hardness of these as-cast HEAs has beenreported.Obviously,Al 0.5CrFeNiCo 0.3C 0.2alloy exhibits the highest hardness of the HEAs listed in Table 4.The high strength and high hardness of Al 0.5CrFeNiCo 0.3C 0.2HEA is possibly due to the formation of the ordered BCC phase and the Cr 23C 6carbide,the nanocrystalline structure,as well as solid solu-tion strengthening mechanism of Al atoms.It is worth pointing out that nanoscale twins might also have a considerable effect on strengthening the alloy.The strengthening mechanism of nano-scale twins is associated with the complicated interaction between dislocations and twin boundaries.It is believed that twin bound-aries are effective in blocking dislocation motion,especially when the thickness of twin/matrix lamellae decreases down to the nano-scale,and a Hall–Petch-type relationship exists for twin boundary strengthening [39,40].The limited compression ratio may owing to the brittle Cr 23C 6carbide,the ordered BCC phase,solid solution strengthening of Al atoms and ultrafine porosities.4.ConclusionsAl 0.5CrFeNiCo 0.3C 0.2high entropy alloy with nanocrystalline has been successfully synthesized by MA and SPS.A supersaturated so-lid solution with both FCC and BCC structures is evidently observed after MA.After SPS,Cr 23C 6carbide and an ordered BCC phase are newly formed.The TEM analysis results confirm that the alloy con-sists of one FCC phase,one BCC phase,Cr 23C 6carbide and one or-dered BCC pared with the nominal composition of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy,the FCC phase is enriched in Fe–Ni,the BCC phase is enriched in Ni–Al,while the ordered BCC phase is especially enriched in Al.The grain size of these phases shows a wide distribution ranging from several hundred nanometers to 1l m.Nanoscale deformation twins present only in partial FCC phase.The compressive strength,compression ratio and Vickers hardness of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy are 2131MPa,3.0%and 617±25HV,respectively.AcknowledgementsThe authors wish to acknowledge the financial support by National Natural Science Foundation of China (Grant No.51271080.)and Project supported by Guangdong Provincial Natural Science Foundation of China (Grant No.S2011010002227.).References[1]Zhao ZD,Chen Q,Chao HY,Huang SH.Microstructural evolution and tensilemechanical properties of thixoforged ZK60-Y magnesium alloys produced by two different routes.Mater Des 2010;31(4):1906–16.[2]Yeh JW,Chen SK,Lin SJ,Gan JY,Chin TS,Shun TT,et al.Nanostructured high-entropy alloys with multiple principal elements:novel alloy design concepts and outcomes.Adv Eng Mater 2004;6(5):299–303.and SAED pattern of twinned FCC phase.(a)a twinned FCC grain with 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Microstructure evolution and mechanical properties of1 000 MPa cold rolled dual-phase steelZHAO Zheng-zhi(赵征志), JIN Guang-can(金光灿), NIU Feng(牛枫), TANG Di(唐荻), ZHAO Ai-min(赵爱民) Engineering Research Institute, University of Science and Technology Beijing, Beijing 100083, ChinaReceived 10 August 2009; accepted 15 September 2009Abstract: The microstructure evolution of 1 000 MPa cold rolled dual-phase (DP) steel at the initial heating stages of the continuous annealing process was analyzed. The effects of different overaging temperatures on the microstructures and mechanical properties of 1 000 MPa cold rolled DP steel were investigated using a Gleeble−3500 thermal/mechanical simulator. The experimental results show that ferrite recovery and recrystallization, pearlite dissolution and austenite nucleation and growth take place in the annealing process of ultra-high strength cold rolled DP steel. When being annealed at 800 ℃ for 80 s, the tensile strength and total elongation of DP steel can reach 1 150 MPa and 13%, respectively. The microstructure of DP steel mainly consists of a mixture of ferrite and martensite. The steel exhibits low yield strength and continuous yielding which is commonly attributed to mobile dislocations introduced during cooling process from the intercritical annealing temperature.Key words: cold rolled dual-phase steel; microstructure evolution; recrystallization; mechanical property; overaging temperature1 IntroductionAdvanced high-strength steels (AHSS) have been used in the automotive industry as a solution for the weight reduction, safety performance improvement and cost saving. Among them, the dual-phase (DP) steels, whose microstructure mainly consists of ferrite and martensite, are an excellent choice for applications where low yield strength, high tensile strength, continuous yielding, and good uniform elongation are required [1−4].The continuous annealing process to produce cold rolled DP steels typically has the following stages: heating to the intercritical temperature region, soaking in order to allow the nucleation and growth of austenite, slow cooling to the quench temperature, rapid cooling to transform the austenite into martensite, overaging, and air cooling. The amount and morphology of the constituents formed depend on such annealing parameters. The effects of the retained austenite, ferrite, and martensite morphologies on the mechanical behavior of DP steels have been intensively investigated[5−9]. As we all known, overaging treatment is an important process during the production of dual-phase steel. It can reduce the hardness of martensite and improve the comprehensive mechanical properties of DP steel [10−14].The purpose of the present research was to study the microstructure evolution of cold rolled DP steel at the initial heating stages of the continuous annealing process using a Gleeble simulator. At the same time, the effects of overaging temperature on the mechanical properties of DP steel were also studied. The microstructures of specimens simulated on a Gleeble simulator, were analyzed using scanning electron microscopy (SEM) and transmission electron microscopy (TEM).2 ExperimentalThe chemical compositions of the experimental steel (mass fraction, %) were: 0.14−0.17C, 0.40−0.60Si, 1.70−1.90Mn, 0.02−0.04Nb, 0.40−0.60Cr, ≤0.010P, ≤0.010S, 0.02−0.06Al and balance Fe. Firstly, experimental steels were smelted in a 50 kg vacuum induction furnace. After smelting, experimental steels were forged into 35 mm×100 mm×100 mm cubic samples. The forged slabs were reheated to 1 200 ℃and soaked for 1 h. The hot rolled thickness was 3.5 mm after 6 passes rolling. The finish rolling temperature was about 880 ℃. The coiling temperature was 620 ℃. After being pickled in hydrochloric acid, the hot rolledFoundation item: Project(2006BAE03A06) supported by the National Key Technology R&D Program during the 11th Five-Year Plan Period Corresponding author: ZHAO Zheng-zhi; Tel: +86-10-62332617; E-mail: zhaozhzhi@ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s564bands were cold rolled to the final thickness of 1.0 mm, and the reduction was about 70%. Finally, the cold rolled sheets were cut into the samples for the simulation of continuous annealing experiment.The microstructure evolution at the initial steps of the continuous annealing process was studied using a Gleeble 1500 simulator. The steel was heated at 10 ℃/sto the different heating temperatures (550, 630, 670, 710, 730, 750 and 780 ℃) and held for 20 s followed by water-quenching. The effects of different overaging temperatures on the microstructures and mechanical properties of DP steel were investigated using a Gleeble 3500 simulator. The processing schedules and parameters used are shown in Fig.1. The soaking temperature of intercritical region was set at 800 ℃, soaking time is 80 s; after a slow cooling, the samples were rapidly cooled to 240, 280, 320 and 360 ℃, respectively and soaked for 300 s; at last, the samples were air cooled to the room temperature.Fig.1 Continuous annealing process of DP steelAfter heat treatment, the steel sheet would be cut into standard tensile specimens (length 200 mm, gauge length 50 mm). The tensile test was performed with CMT4105-type tensile test machine to test mechanical properties. The longitudinal cold rolling plane sections of samples after annealing were prepared and etched with 4% natal. The microstructure was analyzed by scanning electron microscopy (SEM). Some samples were analyzed using transmission electron microscopy (TEM).3 Results and discussion3.1 Mechanical properties and microstructures ofsamples after hot-rolling and continuousannealingTable 1 shows the tensile test data for the two samples after hot-rolling and continuous annealing in terms of yield strength, ultimate tensile strength and total elongation. When the annealing temperature is 800 ℃and soaking time is 60 s, the tensile strength reaches 1 110 MPa and the total elongation reaches 12%. Compared with the hot-rolled samples, the yield strength and total elongation of sample after annealing are similar, but the tensile strength increases by about 450 MPa. The yield ratio decreases obviously. The engineering uniaxial tensile stress—strain curve of the sample after continuous annealing is characterized by very uniform plastic flow until necking. There is no physical yield point and yield point extension, that is, the steel exhibits continuous yielding which is commonly attributed to mobile dislocations introduced during cooling from the intercritical annealing temperature. Many dislocation sources come into action at low strain and plastic flow begins simultaneously through the specimen, thereby suppressing discontinuous yielding[15].Table 1 Mechanical properties of samples after hot rolling and annealingConditionYieldstrength/MPaTensilestrength/MPaYieldratio*Totalelongation/% Hot rolling555 665 0.83 16 Annealing540 1110 0.49 12* Yield ratio is defined as the ratio of yield strength to tensile strength.The microstructures of the hot-rolled and cold-rolled samples are shown in Fig.2. It can be observed that hot rolled steel features a band microstructure, i.e. pearlite band in a ferrite grain matrix. The ferrite grain size is measured to be 5.0−9.0 µm. After cold rolling, the microstructure consists of elongated grains of ferrite and deformed colonies of pearlite (Fig.2(b)). After cold-rolling, there is an increase in the stored energy of the steel due to the high dislocation density and this provides the driving pressure for the ferrite recrystallization during annealing process. The total ferrite grain boundary area increases and the cementite laminar structure in pearlite is broken down. The latter has been shown to promote spheroidization of cementite during subsequent annealing process.The SEM micrograph of the sample after annealing is given in Fig.3(a). The microstructure of DP steel consists of a mixture of ferrite, martensite, martensite/austenite constituent. There is also some bainite in the microstructure. The martensite islands are homogeneously distributed in ferrite matrix. The DP steel has finer grain size and the size of ferrite grain and martensite island are about 1.0−2.0 µm. Some martensite islands have a bright white circle around the edge, and the center of martensite is of irregular black structure.ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s565Fig.2 Microstructures of steel after hot rolling (a) and cold rolling (b)Fig.3 SEM images (a) and TEM micrograph (b) of steel after continuous annealingThe main reason is the manganese partitioning will occur during the continuous annealing process. During the heating process, a high-Mn side lap forms around austenite, which makes the hardenability of austenite island edge higher than that of the center. So, it makes high-Mn side lap form around martensite in the cooling process. The volume fraction of martensite is about 40%, which is the main reason for DP steel with a higher strength. After the continuous annealing process, band structure is significantly improved, which plays an important role in improving the performance of DP steel.The fine structures of martensite and ferrite are shown in Fig.3(b) by the TEM observation. The lath martensite is fine, and is relatively clean; at the same time, a very high density of dislocations can be observed in the ferrite grain adjacent to martensite. These dislocations are generated in order to accommodate transformation induced strain built between martensite transformed by quenching and retained ferrite. In addition, they are known to be mobile and play an important role on rapid, extensive strain hardening of DP steel from the onset of its plastic deformation.3.2 Microstructure evolution at initial steps ofcontinuous annealing processThe microstructure evolution at the initial stages of the continuous annealing process is very important for producing the ultra-high strength DP steel. During the annealing process of high strength DP steel, ferrite recovery and recrystallization, pearlite dissolution and austenite nucleation and growth will occur. When the sample is heated to 550 , the℃microstructure has no visible change as compared with the cold rolled sample. The ferrite grain is stretched along the rolling direction significantly; lamellar pearlite is stretched along the rolling direction too. At the same time, there are some carbide particles in the ferrite matrix, as shown in Fig.4(a). At this temperature, the recrystallization nucleus was not found in the structure. So, at this stage the sample is still at the recovery stage. When the heating temperature is 630 , the℃recrystallization nucleus begins to appear in the microstructure. The nucleus of crystal appears mainly nearby the large deformation ferrite (Fig.4(b)). The recrystallization nucleus is fine and equiaxed. Large deformation storage power is present in the large deformation region. So, recrystallization nucleus forms in this region firstly. With the heating temperature increasing, the recrystallization nucleus begins to grow. Therefore, the size of recrystallization is uneven at this stage, as shown in Fig.4(c). When the heating temperature is 670 ℃, the deformation structure still exists in the microstructure. With the temperature increasing, the deformed ferrite grains are replaced by recrystallization ferrite grains. When the heating temperature is 710 , the d℃eformation structure has already vanished, which is replaced by theZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s566equiaxed recrystallization grain. So, the process of recrystallization completes basically. In the ferrite recrystallization process, the pearlite transforms to granular from lamellar gradually.When the heating temperature is 730 ,℃it begins to enter the two-phase region; and the ferrite and spheroidised carbides begin to transform to austenite. A small amount of austenite nucleates in the original pearlite region, as shown in Fig.4(e). Austenite nucleates mainly in the ferrite and pearlite grain boundary; and a part of austenite also nucleates in the carbide particles of ferrite. After austenite nucleation, it begins to grow rapidly. At this stage, the pearlite dissolves rapidly. When the temperature reaches 750 , the austenite℃transformation occurs obviously. The bright white particle which distributes in the ferrite matrix is the martensite island. The martensite transforms from austenite during the rapid cooling process. At the same time, a small amount of martensite particles can also be observed in ferrite; and there are still some non-dissolved carbide particles in the ferrite matrix. The initial austenite growing-up is mainly controlled by the carbon Fig.4Microstructure evolutions duringcontinuous heating process: (a) 550 ℃; (b)630 ℃; (c) 670 ℃; (d) 710 ℃; (e) 730 ℃; (f)750 ℃; (g) 780 ℃ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s567diffusion in the austenite, and the diffusion path is along the pearlite/austenite interface. When the annealing temperature is 780 , the austenite volume increase℃s, and the number of carbide particles is reduced gradually. There is only a very small amount of carbide particles distributing in ferrite matrix.3.3 Effect of overaging temperature onmicrostructure and mechanical properties ofDP steelThe overaging is a temper treatment to harden martensite in the dual-phase steel, reduce the hardness of martensite and improve the comprehensive mechanical properties[16]. Fig.5 shows the effect of overaging temperature on the mechanical properties of dual-phase steel. All the samples are intercritically annealed at 800℃ with different overaging temperatures. As can be seen from Fig.5, the highest tensile strength is achieved in the sample overaged at 280 ℃. The yield strength is 560 MPa, the tensile strength is 1 150 MPa, and the total elongation reaches 13%. The good combination of high strength and toughness properties is obtained. And then, with the increase of overaging temperature, the yield strength and tensile strength of samples decrease, while the total elongation increases. When the overaging temperature reaches 360 ℃, the tensile strength decreases, the yield strength does not change significantly. The mechanical properties of sample cannot meet the necessary requirements of CR980DP. At the same time, the stress—strain curve of the steel shows discontinuous yielding behaviour and develops yield plateaus.Fig.6 shows the SEM microstructures with different overaging temperatures. It can be seen that the microstructure mainly consists of dark grey ferrite grains and white martensite. When the overaging temperature is 360 ℃, the martensite boundary is fuzzier than that of sample overaged at 320 ℃, and there are more carbides, which is due to the effects of tempering on the martensite, such as the volume contraction of martensite during the tempering, the changes of the martensite strength and additional carbon clustering or precipitation near the ferrite and martensite interfaces.Fig.5 Effects of different overaging temperatures on mechanical propertiesFig.6 SEM images of microstructures of DP steel overaged at different temperatures: (a) 240 ℃; (b) 280 ℃; (c) 320 ℃; (d) 360 ℃ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s5684 Conclusions1) When the DP steel is annealed at 800 ℃ for 80 s and overaged at 280 ℃, the tensile strength and total elongation of ultra-high strength dual-phase steel can reach 1 150 MPa and 13%, respectively.2) The microstructure of DP steel consists of a mixture of ferrite, martensite, martensite/austenite constituent. There are also some bainites in the microstructure. The martensite islands are homogeneously distributed in ferrite matrix. The ferrite and martensite island grain size are about 1.0−2.0 µm. When the overaging temperature reaches 360 ℃, the tensile strength decreases, the yield strength does not change significantly. The mechanical properties of sample cannot meet the necessary requirements of CR980DP. At the same time, the steel shows discontinuous yielding behaviour and develops yield plateaus.References[1]KANG Yong-lin. Quality control and formability of the mordernMotor plate [M]. 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第52卷第10期表面技术2023年10月SURFACE TECHNOLOGY·241·Na2WO4含量对镁合金微弧氧化膜层颜色和耐蚀性的影响祝海涛1,2，孙金峰1,3，孟永强1,3，王立伟1,2，彭珍珍1,2，汪殿龙1,2*（1.河北科技大学 材料科学与工程学院，石家庄 050018；2.河北省材料近净成形技术重点 实验室，石家庄 050018；3.河北省柔性功能材料重点实验室，石家庄 050000）摘要：目的研究电解液中的Na2WO4含量对AZ31B镁合金微弧氧化膜层的形成过程、颜色、微观结构、耐蚀性能的影响。

方法通过添加不同含量的NH4VO3和Na2WO4的碱性铝酸盐电解液体系，在AZ31B镁合金表面制备黑色的微弧氧化膜层。

采用SEM、EDS分析加入不同含量的Na2WO4后膜层表面的微观形貌及元素组成，采用XRD分析物相组成，通过电化学实验表征膜层的耐腐蚀性能。

结果随着Na2WO4含量的增加，微弧氧化过程中的起弧电压下降，膜层的致密性提高，厚度呈先增加后减小的趋势。

当Na2WO4的质量浓度为0.5 g/L时，膜层的厚度最大，且此时膜层表面微孔分布均匀，色度最低，耐蚀性最好，自腐蚀电位为−0.138 V，自腐蚀电流密度为2.36×10−7 A/cm2，相较于基体降低了3个数量级。

结论增加Na2WO4含量会使微弧氧化成膜过程中的电弧发生变化，适当增加Na2WO4会提高膜层的厚度，降低膜层的CIE色度，使陶瓷膜层表面的微孔分布得更加均匀致密，从而提高膜层的耐蚀性能。

当Na2WO4含量过高时，会使膜层的离子浓度升高，电阻增大，介电击穿电压上升，导致膜层表面被烧蚀，耐腐蚀性能降低。

关键词：镁合金；黑色膜层；微弧氧化；Na2WO4；耐腐蚀性；微观结构中图分类号：TG174.4文献标识码：A 文章编号：1001-3660(2023)10-0241-09DOI：10.16490/ki.issn.1001-3660.2023.10.019Effect of Na2WO4 Content on Color and Corrosion Resistance ofMicro-arc Oxidation Coating on Magnesium AlloyZHU Hai-tao1,2, SUN Jin-feng1,3, MENG Yong-qiang1,3, WANG Li-wei1,2,PENG Zhen-zhen1,2, WANG Dian-long1,2*(1. School of Materials Science and Engineering, Hebei University of Science and Technology, Shijiazhuang 050018, China;2. Key Laboratory of Material Near-net Forming Technology in Hebei Province, Shijiazhuang 050018, China;3. Hebei Key Laboratory of Flexible Functional Materials, Shijiazhuang 050000, China;)ABSTRACT: In this work, the method preparing micro-arc oxide coating on AZ31B magnesium alloy surface was studied, the formation process of AZ31B magnesium alloy micro-arc oxide ceramic layer was explored by the content of Na2WO4 in electrolyte, and the microstructure and corrosion resistance of AZ31B magnesium alloy were characterized. By using NaAlO2 as收稿日期：2022-09-23；修订日期：2023-02-10Received：2022-09-23；Revised：2023-02-10基金项目：国家自然科学基金（52101015）；河北省高等学校科学技术研究项目（BJK2022020）；河北省自然科学基金（E2022208070/E2021208005）Fund：National Natural Science Foundation of China (52101015)；Science and Technology Project of Hebei Education Department (BJK2022020)；Natural Science Foundation of Hebei Province (E2022208070/E2021208005)引文格式：祝海涛, 孙金峰, 孟永强, 等. Na2WO4含量对镁合金微弧氧化膜层颜色和耐蚀性的影响[J]. 表面技术, 2023, 52(10): 241-249. ZHU Hai-tao, SUN Jin-feng, MENG Yong-qiang, et al. Effect of Na2WO4 Content on Color and Corrosion Resistance of Micro-arc Oxidation Coating on Magnesium Alloy[J]. Surface Technology, 2023, 52(10): 241-249.*通信作者（Corresponding author）·242·表面技术 2023年10月the basic electrolyte and adding a small amount of NH4VO3 and Na2WO4 on the surface of AZ31B magnesium alloy, a black micro-arc oxidation ceramic layer was prepared. Scanning electron microscope (SEM) and EDS were adopted to analyze the surface morphology and composition of the coatings with different contents of Na2WO4 and XRD was applied to analyze the phase structure. Electrochemical workstation was used to test the self-corrosion potential and self-corrosion current density of the micro-arc oxidation ceramic layer, and the CIE chromaticity of the coating layer was measured by 3nh chromaticity meter.After experimental verification, NH4VO3 with interaction to Na2WO4 contributed to the successful preparation of the black micro-arc oxidation ceramic layer. Through observation, it was found that with the increase of Na2WO4 content, the arc started to drop.The arc starting voltage was reduced from 320 V to 210 V. In the process of preparation, the micro-arc oxidation coating density was significantly increased and coating thickness increased firstly and then decreased. With the increase of Na2WO4 content, CIE chromaticity decreased firstly and then increased, and the corrosion resistance of the sample increased.When the content of Na2WO4 was 0.5 g/L, the micropores on the surface of the coating were evenly distributed, uniform in size and low in porosity, and the thickness of the coating was the largest at this time. The thickness was 23.8 μm. The CIE chromatic value was the smallest, which was 27.99, and the corrosion resistance was the best. The electrochemical test in 3.5% NaCl solution showed that the self-corrosion potential was −0.138 V and the self-corrosion current density was 2.36×10–7 A/cm2, which decreased by 3 orders of magnitude compared with the matrix. At the same time, compared with the corrosion current density of 1.15×10–5A/cm2, its corrosion current density under the condition of no Na2WO4, decreased by 2 orders of magnitude.However, when sodium tungstate was added to 1 g/L, the coating appeared ablative phenomenon and the corrosion resistance decreased. The self-corrosion current density is not different from that of the matrix. It is concluded that the increase of Na2WO4 content will change the arc in the process of micro-arc oxidation and reduce the arc starting voltage. An appropriate increase of Na2WO4 will increase the thickness of the coating layer, reduce the CIE value of the coating layer, make the distribution of micropores on the surface of the ceramic coating layer more uniform and compact, improve the compactness of the coating layer, and thus improve the corrosion resistance of the coating layer. However, when the content of Na2WO4 is too high, the concentration of ions in the coating layer will increase, and more ions will adsorb to the surface of the coating layer for reaction.The resistance of the coating layer increases, leading to the rise of dielectric breakdown voltage, resulting in the surface ablation of the coating layer and decrease in coating densification and corrosion resistance.KEY WORDS: magnesium alloy; black coating; micro-arc oxidation; Na2WO4; corrosion; microstructure近年来，随着镁合金材料的广泛应用，对其性能的要求越来越高，镁合金作为轻质的金属结构材料，具有电磁屏蔽性好、抗震性好、比强度高、比刚度高、易回收等特点[1-2]，在汽车、航空航天、3C产品等领域具有广阔的应用前景。

表面技术第52卷第3期Al2O3-13%TiO2绝缘防护复合涂层组织及电偶腐蚀性能吴护林1，张智峰2，彭冬1，杨钊2，宋凯强1，丛大龙1，李毅2，谢杨2，陈爽2，黄安畏1，李忠盛1（1.西南技术工程研究所，重庆 400039；2.中国核动力研究设计院，成都 610094）摘要：目的研究等离子喷涂的Al2O3-13%TiO2涂层和封孔处理后的Al2O3-13%TiO2复合涂层对TC4-H70异种金属电偶对的腐蚀防护效果。

方法采用X射线衍射仪、扫描电镜、能谱仪对涂层的物相组成、组织形貌、元素分布进行表征分析，使用电化学工作站和电偶腐蚀测量仪对涂层及对比试样的耐蚀性能进行分析研究。

结果等离子喷涂的Al2O3-13%TiO2涂层由α-Al2O3和γ-Al2O3两相组成，以γ-Al2O3相为主。

Al2O3-13%TiO2涂层中存在微孔与微裂纹等缺陷，腐蚀介质易渗入，因此Al2O3-13%TiO2涂层的耐蚀性较差。

经过封孔处理后的Al2O3-13%TiO2涂层，表层缺陷被充分填充，同时在陶瓷表层形成厚度为20~40 μm的致密阻挡层，有效阻隔了NaCl腐蚀介质的渗入，涂层腐蚀电流密度相比于H70基体和Al2O3-13%TiO2涂层试样减小了4个数量级，基底与涂层间的界面电荷转移电阻值相较于H70基体和Al2O3-13%TiO2涂层提高了5个数量级，涂层耐蚀性和绝缘性显著提升。

TC4-H70电偶对经15 d电偶腐蚀试验后，H70表面发生腐蚀，封孔处理后的Al2O3-13%TiO2涂层可有效降低TC4-H70电偶对间的电偶腐蚀作用，经15 d电偶腐蚀试验后，试样未腐蚀，且涂层完整。

结论封孔处理后的Al2O3-13%TiO2涂层优异的电偶腐蚀防护效果主要得益于其高电阻和表层的高致密性，几乎阻隔了异种金属间的电子传输，使得异种金属间的电偶电池作用极其微弱，可有效延长异种金属海水管路的使用寿命，在电偶腐蚀防护领域具有巨大的应用前景。

Microstructure,mechanical properties and thermal shock resistance of nano-TiN modified TiC-based cermets with different bindersXiaobo Zhang,Ning Liu *Department of Materials Science and Engineering,Hefei University of Technology,Hefei 230009,ChinaReceived 24December 2007;accepted 27January 2008AbstractMicrostructure,mechanical properties and thermal shock resistance of TiC-based cermets modified by nano-TiN were studied.The highest transverse rupture strength,fracture toughness and hardness were obtained for the cermets with the compositions 14Mo–20Ni (cermet A),4Mo–10Co–10Ni (cermet B)and 4Mo–20Co (cermet C),respectively.The quenching-strength method and indentation-quench method were used for studying the thermal shock resistance of nano-TiN modified TiC-based cermets.The results of quench-ing-strength test denote that the critical temperature differences for sharp decline of residual strength are about 410,370and 330°C for cermets A,B and C,respectively.The results of indentation-quench test show that the micro crack length increases and the micro cracks propagate fast with increasing thermal shock temperature and cycles.The propagation rate of the crack is controlled by the stress intensity of crack tip.Both tests show that cermets with high transverse rupture strength have excellent thermal shock resistance.Ther-mal shock resistance parameter R and R st calculated are in good agreement with the experimental results.Ó2008Published by Elsevier Ltd.Keywords:Nano-TiN modified TiC-based cermets;Binder phases;Thermal shock resistance;Mechanical properties1.IntroductionThe unique combination of mechanical properties such as excellent wear resistance and good chemical stability at elevated temperature makes Ti(C,N)-based cermets being of great importance in metal cutting operations.Nowa-days,cermets cutting tools are widely used for semi-finish-ing and finishing works on steel and cast iron [1–6].However,because of their relatively low fracture toughness and thermal shock resistance,Ti(C,N)-based cermets are still limited for wider utilization.With the development of nano-technology,nano-modified cermets have been received more and more attentions due to the available high properties [7–10].The initial grain size plays a crucial role in the microstructure development during sintering,particularly when the grains are of nano-crystalline nature because the driving force for densification during sinteringis the lowering of surface energy.The evolution of the microstructure during sintering in the presence of nano-powder is expected to reflect the impact of extremely high rate of diffusion caused by the large grain boundary frac-tion [11].Ti(C,N)-based cermets used in high temperature applica-tion as a material of cutting tools are often exposed to rapid temperature changes in which case will cause thermal dam-age.Therefore,it is necessary to clarify the repeated thermal shock behavior of the cutting materials.The quenching-strength test is a common method to determine the thermal shock resistance of ceramics.Another method to evaluate the thermal shock resistance of brittle materials is indenta-tion-quench test.Thermal shock resistance of ceramics has been widely studied [12–15].However the research about the thermal behavior of cemented carbides and cermets lacks of reports.In this paper,besides the microstructure and mechanical properties,the thermal shock resistance of nano-TiN modified TiC-based cermets was studied both by indentation-quench test and quenching-strength test,0263-4368/$-see front matter Ó2008Published by Elsevier Ltd.doi:10.1016/j.ijrmhm.2008.01.008*Corresponding author.Tel.:+865512909865;fax:+865512905383.E-mail address:ningliu@ (N.Liu)./locate/IJRMHMAvailable online at International Journal of Refractory Metals &Hard Materials 26(2008)575–582and the effect of temperature and metallic binder on the ther-mal shock resistance were also discussed.Furthermore,the thermal shock resistance parameter R(D T c)and R st were cal-culated and compared with the experimental results in order to better understand the thermal behavior of cermets.2.Experimental procedures2.1.Samples preparationCommercially available TiC(2.56l m),nano-TiN (0.04l m),WC(3.52l m),Mo(2.33l m),Ni(2.95l m), Co(2.46l m)and C(3.25l m)were used as raw materials. The theoretical density and chemical composition of raw powders are listed in Table1,and the composition design of the experimental cermets is listed in Table2.The SEM images of TiC and TiN raw powders are shown in Fig.1.Powder mixtures were wet milled with WC–Co balls in ethanol bath for24h by a planetary ball mill and then dried.Green compacts were prepared by pressing at the uniaxial pressure of200MPa,dewaxed at800°C at a heat-ing rate of0.5°C/min and vacuum sintering(0.1Pa)was conducted at1430°C for1h.2.2.Experimental methodsThe mechanical properties of cermets at room tempera-ture,such as transverse rupture strength,fracture tough-ness and hardness,were measured.Transverse rupture strength test was conducted on a CMT-5105electron test-ing machine by three-point bending method(span20mm, crosshead speed0.5mm/min).The geometric sizes of sam-ples are5Â5Â30mm.Vickers hardness,H V,was exam-ined under an indentation load of30kg,and the fracturetoughness was calculated from H V,according to the for-mula proposed by Shetty et al.[16].The theoretical densi-ties of cermets,d t,were calculated by the formula:d t¼100%=ða%=d aþb%=d bþÁÁÁÞð1Þwhere a%,b%etc.are the mass of the starting compositions (wt%),d a,d b etc.are the theoretical densities of the starting powders.The actual densities of cermets,d cer,were measured by using Archimedes method with distilled water and calcu-lated by the formula:d cer¼G1d=ðG1þG2ÀG3Þð2Þwhere G1is the weight of the sample,G2is the weight of the thin copper line,G3is the weight of the sample and the cop-per line in distilled water and d=1.0g/cm3is the density of the distilled water.The relative density of each cermet is the ratio of their actual density with theoretical density,i.e.d cer/d t.Samples for the quenching-strength test were heated at 250,350,450and550°C,respectively,for10min and then quenched into water at room temperature(20°C).The geo-metric sizes of samples are5Â5Â30mm.The residual strength was measured by three-point bending method. The supplementary test was conducted at390,410andTable1Theoretical density and chemical compositions of starting powders Powder Theoretical density(g/cm3)Chemical composition(wt%) TiC 4.93C free:0.179,O:0.13TiN 5.20C<1.0,O<2.0WC15.55C free:0.02Mo10.20C:0.0036,Fe:0.002,O:0.095 Ni8.90C<0.15,S<0.001,O<0.015 Co8.90C:0.058,S:0.0045,O:0.008C 2.25N:0.00015,O:0.3Table2Composition design of the cermets(wt%)Cermet TiC TiN WC Mo Ni Co C A401015142001 B501015410101 C50101540201Fig.1.SEM images of raw powders(a)micro-sized TiC,(b)nano-sized TiN.576X.Zhang,N.Liu/International Journal of Refractory Metals&Hard Materials26(2008)575–582430°C for evaluating the critical temperature exactly according to the obtained results.The indentation-quench method was also used for studying the thermal shock resis-tance.Indentation was made by Vicker’s hardometer on the samples.Samples with indentation for indentation-quench test were heated at300,400,500and600°C for 10min,respectively,and then quenched into water at room temperature.The micro crack length of indentation was measured by an optical microscope after polishing the sur-faces.The thermal shock test was repeated at least25cycles according to the requirement of the test.Microstructures and fracture surfaces were observed by a scanning electron microscope(LEO-1530VP,LEO,Ger-many)coupled with an X-ray energy-dispersive spectrome-ter(OXFORD INCA X-Sight,UK)in back-scattered electron mode(BSE)and secondary electron model(SE). The phase identification of each system was carried out by XRD(D/max-rB,Rigaku,Japan).Thin foils for trans-mission electron microscopy(TEM)were prepared by ion milling using5kV argon ions in a dual-ion mill model Gadan600.TEM analysis was performed on a Hitachi transmission electron microscope(Japan)model H-800 operated at200kV.3.Results and discussion3.1.Microstructure and mechanical properties of cermets at room temperatureThe microstructures of cermets at room temperature are shown in Fig.2.It is denoted that the typical black core/ grey rim ceramic grains embed into white binder,some rims show two parts:bright inner rim and grey outer rim. Moreover,some coreless grains are also observed.Black core is the remnant of undissolved TiC,grey rim is the solid solution(Ti,Mo,W)C that formed in the stage of liquid sintering,white binder is mainly Ni and coreless grains are(Ti,Mo,W)(C,N).The formation mechanism of these phases has been discussed in references[8,10,17].More-over,the grain size of cermet A isfiner than those of cer-mets B and C.This can be due to the contribution of 14wt%Mo,which can inhibit grain growth effectively [18].When the content of Mo is just4wt%,the grains of cermets B and C become coarser.Even though the total content of binder is the same and the average grain size is also close in cermets B and C;there is still something dif-ferent:the distribution of grain size of cermet B is more average than that of cermet C.The microstructure observed in cermet C shows that some grains size is more than4l m while some grains size is less than1l m.This asymmetric distribution of grains will harm the mechanicalproperties.Fig.3shows the XRD profiles of three cermets. The results denote that the cermets consist of ceramics and binders.TEM micrographs of microstructure of cermets are shown in Fig.4.Some nano-TiN particles about100nm are observed at the interfaces of TiC/TiC grains.The impedance of nano-TiN particles on the growth of TiC matrix makes the microstructure of cermetsfiner[19].Mechanical properties and relative densities of cermets at room temperature are listed in Table3.Cermet A shows the higher strength and hardness compared with that of cermet B.Addition of14wt%Mo contributes a lotto Fig.2.SEM micrographs showing microstructures of cermets(a)cermet A,(b)cermet B,(c)cermet C.X.Zhang,N.Liu/International Journal of Refractory Metals&Hard Materials26(2008)575–582577the properties.Some papers have indicated that adding Mo or Mo 2C into Ti(C,N)-based cermets can improve the wettability between ceramic phase and metallic phase and decrease the grain size of cermets [18,20].Based on Hall–Petch formula,cermet A with the finer grain size should have higher strength and hardness than that of cermet B as shown in Fig.2.The second reason for higher strength and hardness of cermet A is the solid solution strengthen-ing effect of binder caused by molybdenum.It is a general phenomenon that cermet with coarse grains has higher fracture toughness than that of cermet with finer grains [8,10].The reasons are due to the higher bonding force inthe transcrystalline fracture of coarse grains than that of in the intergranular fracture of fine grains and increasing the propagation path caused by part of intergranular frac-ture of coarse grains.The content of Mo in cermets B and C is the same,i.e.4wt%,so the grain sizes of cermets B and C are nearly the same.Cermet B has higher fracture toughness and trans-verse rupture strength than that of cermet C should be related to the better distribution of grain size as shown in Fig.2.Nevertheless,cermet C has the higher hardness which may be due to the better wettability of cobalt in which case improves the bonding force between the ceram-ics and metals.The results of this experiment are in good agreement with references [21].The theoretical densities of cermets A,B and C calculated by Eq.(1)are 6.638,6.206and 6.206g/cm 3,respectively,and the actual densi-ties of them are 6.637,6.203and 6.204g/cm 3,respectively.The relative densities of the cermets are listed in Table 3.The results show that almost full densities of nano-TiN modified TiC-based cermets are obtained.3.2.Thermal shock resistance of quenching-strength method The residual strength of cermets for single thermal shock is plotted in Fig.5.The residual strength behavior of the three kinds of cermets follows Hasselman’s theory [22,23]that predicts a discontinuity with considerable strength degradation at a critical thermal shock temperature differ-ence D T c .D T c values are about 410,370and 330°C for the cermets A,B and C,respectively.Several parameters of thermal shock resistance were defined to relate thermo-physical and thermo-mechanical properties of the materials,such as R ,R 0,R 0000and R st [22–24].R and R 0parameters predict the critical tempera-ture difference in a body under conditions of sharp and moderate heat flow,respectively,which are expressed as the resistance to crack initiation.R 0000and R st parameters express the ability of a material to resist crackpropagationFig.3.XRD profiles of threecermets.Fig. 4.TEM micrographs showing the nano-TiN distributed at the interface of TiC/TiC (cermet A).(a)TEM bright field,(b)SAED pattern and index of nano-TiN.Table 3Mechanical properties and relative densities of cermets Cermet r bb (MPa)H V (GPa)K IC (MPa m 1/2)Relative density (%)A 152117.512.399.98B 143917.112.799.95C125118.111.699.96Fig. 5.Residual strength of cermets as a function of thermal shock temperature difference (D T )for single thermal shock.578X.Zhang,N.Liu /International Journal of Refractory Metals &Hard Materials 26(2008)575–582and further damage and loss of strength with increasing severity of thermal shock [12].In this experiment,samples were heated and cooled rapidly;furthermore,the relative density of cermets is over 99.95%,so parameter R ,which predicts the critical temperature difference (D T c )in a body under condition of sharp heat and cool flow,is used to rep-resent the thermal shock resistance:D T c ¼R ¼r ð1Àm ÞE að3Þwhere r is the strength,m is the Poisson ratio,E is the Young’s modules and a the thermal expansion coefficient.It indicates that D T c is determined by strength (r ),Poisson ratio (m ),Young’s modules (E )and thermal expansion coef-ficient (a )of materials.The values of E ,m and a of compos-ites can be calculated according to the following equations [25,26]:E ¼E 1V 1þE 2V 2þÁÁÁð4Þm ¼m 1V 1þm 2V 2þÁÁÁÁð5Þa ¼a 1K 1F 1=q 1þa 2K 2F 2=q 2þÁÁÁK 1F 1=q 1þK 2F 2=q 2þÁÁÁð6Þwhere V is the volume fraction of material,F is the weight fraction,q is the density and K =E /3(1À2m )is the volume module.The physical parameters of raw materials used in cer-mets are listed in Table 4[26–28],the physical parameters and D T c of cermets calculated according to the equations above are shown in Table 5.It indicates that E ,m and a of cermets with different binders change slightly,hence D T c is mainly determined by r .The values of D T c calcu-lated are close to those measured by quenching-strength test.The good agreement relationship of D T c obtained by experiment and calculation were also reported in ceramics [29,30].Fig.6shows the fracture surfaces of cermets thermal shocked at temperature difference D T =370°C.It is denoted that the fracture surface of cermet A shows finer morphology and more intergranular fracture than that ofcermet B,while the fracture surface of cermet C is nearly all rough transgranular fracture.3.3.Thermal shock resistance of indentation-quench method The crack length with the quenching cycles is plotted in Fig.7.The pattern of crack propagation is similar in threeTable 4Physical properties of raw materials in cermets Parameters TiC TiN WC Mo Ni Co C E (GPa)379350731336197219690m0.190.220.220.320.290.310.2a (Â10À6K À1)7.749.353.846.2013.213.52.1Table 5Calculated physical properties and critical thermal shock temperature difference (D T c )of cermets Cermet E (GPa)m a (Â10À6K À1)D T c (°C)A 3760.2237.54417B 3800.2147.76383C3820.2167.83328Fig.6.SEM fracture surfaces of cermets thermal shocked at D T =370°C (a)cermet A,(b)cermet B,(c)cermet C.X.Zhang,N.Liu /International Journal of Refractory Metals &Hard Materials 26(2008)575–582579tested cermets and the process has the following rules:At low temperature(about300°C),no obvious crack propa-gation can be detected even though the thermal shock cycles are more than100times.At higher temperature range(from400to600°C),the crack propagation rate increases with the increase of quenching temperature and times.Some cracks can be seen clearly without an optical microscope after several thermal shock times at higher tem-perature and propagate to the edges of samples.Stress intensity is used for explaining the character of thermal shock resistance[31,32].At low temperature,the stress intensity of crack tip is not big enough to cause the crack to propagate.So no obvious crack propagation can be detected even though the thermal shock cycles are more than100times.However,the damage accumulation will be made on the microstructure located at the crack tip to cer-tain extent with increasing the thermal shock times.As a consequence,the crack propagates with the further increase of thermal shock times.The higher the thermal shock temperature is,the higher the stress intensity of crack tip is.Therefore,the crack is much easier to propa-gate at high temperature.TiC-based cermets with different binder phases exhibit different thermal shock resistances as shown in Fig.7.It is denoted that the propagation rate of the crack in cermets depends on the binder and decreases according to the fol-lowing order:cermet C?cermet B?cermet A.The rea-son is maybe due to the higher oxidization resistance of Ni during the thermal quenching and the better mechanical properties of cermet A.The thermal stress crack stability parameter,R st,which is usually used to evaluate the ther-mal shock properties of ceramics,is calculated by the fol-lowing equation:R st¼½c f=ða2EÞ 1=2ð7ÞIn the calculation of R st,the effective surface energy,c f,was converted according to the Irwin equation:K IC=(2c E)1/2.Fig.7.Curves of micro crack length of cermets and thermal shock repeated cycles at different temperatures.Table6Calculated thermal shock resistance parameter(R st)of cermetsParameter A B CR st(l m1/2°C)306830453002 580X.Zhang,N.Liu/International Journal of Refractory Metals&Hard Materials26(2008)575–582The calculated results(as shown in Table6)show that R st decreases from cermet A to C.The calculated results corre-late well with the experimental results.The same result was obtained in SiC whisker reinforced Si3N4ceramic compos-ites[33].4.Conclusions(1)Cermet with14Mo–20Ni as the binder shows the besttransverse rupture strength,cermet with4Mo–10Co–10Ni presents the highest fracture toughness while cermet with4Mo–20Co exhibits the highest hardness.(2)The results of quenching-strength test show that thecritical temperature differences for sharply decline of residual strength are about410,370and330°C for the binder14Mo–20Ni,4Mo–10Co–10Ni and 4Mo–20Co,respectively.(3)The results of indentation-quench test show that thecrack propagates with the increase of thermal shock temperature and cycles.The propagation rate of the crack is controlled by the stress intensity of crack pared with the cermets with4Mo–20Co, the thermal shock resistance of cermet with4Mo–10Ni–10Co is better and that of cermet with14Mo–20Ni is the best.(4)The results of thermal shock resistance of nano TiNmodified TiC-based cermets calculated correlate well with the results tested by quenching-strength and indentation-quench experiments. 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