MECHANICAL PROPERTIES AND MICROSTRUCTURE OF Mg-5Li-5Al-3Zn-xCd 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 12 期第 113-122 页材料工程Vol.51Dec. 2023Journal of Materials EngineeringNo.12pp.113-122第 51 卷2023 年 12 月基于DSCCR 生产工艺的终轧温度对轧制过程中低碳钢组织与性能影响分析Effect of finish rolling temperature on microstructure and properties of low carbon steel based on DSCCR production process李朝阳1，赵志鹏2*，田鹏3，梁晓慧3，王书桓1*，康永林3（1 华北理工大学 冶金与能源学院，河北 唐山 063210；2 北京科技大学 钢铁共性技术协同创新中心，北京100083；3 北京科技大学 材料科学与工程学院，北京100083）LI Chaoyang 1，ZHAO Zhipeng 2*，TIAN Peng 3，LIANG Xiaohui 3，WANG Shuhuan 1*，KANG Yonglin 3（1 School of Metallurgy and Energy ，North China University of Scienceand Technology ，Tangshan 063210，Hebei ，China ；2 Collaborative Innovation Center of Steel Technology ，University of Science and Technology Beijing ，Beijing 100083，China ；3 School of Materials Science and Engineering ，University of Science and TechnologyBeijing ，Beijing 100083，China ）摘要：基于东华连铸连轧生产线（Donghua steel continuous casting rolling ，DSCCR ）分别进行终轧温度为880，820 ℃和780 ℃的热轧实验，研究终轧温度对低碳钢组织和性能的影响。

Microstructures and properties of high-entropyalloysYong Zhang a ,⇑,Ting Ting Zuo a ,Zhi Tang b ,Michael C.Gao c ,d ,Karin A.Dahmen e ,Peter K.Liaw b ,Zhao Ping Lu aa State Key Laboratory for Advanced Metals and Materials,University of Science and Technology Beijing,Beijing 100083,Chinab Department of Materials Science and Engineering,The University of Tennessee,Knoxville,TN 37996,USAc National Energy Technology Laboratory,1450Queen Ave SW,Albany,OR 97321,USAd URS Corporation,PO Box 1959,Albany,OR 97321-2198,USAe Department of Physics,University of Illinois at Urbana-Champaign,1110West Green Street,Urbana,IL 61801-3080,USA a r t i c l e i n f o Article history:Received 26September 2013Accepted 8October 2013Available online 1November 2013a b s t r a c tThis paper reviews the recent research and development of high-entropy alloys (HEAs).HEAs are loosely defined as solid solutionalloys that contain more than five principal elements in equal ornear equal atomic percent (at.%).The concept of high entropyintroduces a new path of developing advanced materials withunique properties,which cannot be achieved by the conventionalmicro-alloying approach based on only one dominant element.Up to date,many HEAs with promising properties have beenreported, e.g.,high wear-resistant HEAs,Co 1.5CrFeNi 1.5Ti andAl 0.2Co 1.5CrFeNi 1.5Ti alloys;high-strength body-centered-cubic(BCC)AlCoCrFeNi HEAs at room temperature,and NbMoTaV HEAat elevated temperatures.Furthermore,the general corrosion resis-tance of the Cu 0.5NiAlCoCrFeSi HEA is much better than that of theconventional 304-stainless steel.This paper first reviews HEA for-mation in relation to thermodynamics,kinetics,and processing.Physical,magnetic,chemical,and mechanical properties are thendiscussed.Great details are provided on the plastic deformation,fracture,and magnetization from the perspectives of cracklingnoise and Barkhausen noise measurements,and the analysis of ser-rations on stress–strain curves at specific strain rates or testingtemperatures,as well as the serrations of the magnetizationhysteresis loops.The comparison between conventional andhigh-entropy bulk metallic glasses is analyzed from the viewpointsof eutectic composition,dense atomic packing,and entropy of 0079-6425/$-see front matter Ó2013Elsevier Ltd.All rights reserved./10.1016/j.pmatsci.2013.10.001⇑Corresponding author.Tel.:+8601062333073;fax:+8601062333447.E-mail address:drzhangy@ (Y.Zhang).2Y.Zhang et al./Progress in Materials Science61(2014)1–93mixing.Glass forming ability and plastic properties of high-entropy bulk metallic glasses are also discussed.Modeling tech-niques applicable to HEAs are introduced and discussed,such asab initio molecular dynamics simulations and CALPHAD modeling.Finally,future developments and potential new research directionsfor HEAs are proposed.Ó2013Elsevier Ltd.All rights reserved. Contents1.Introduction (3)1.1.Four core effects (4)1.1.1.High-entropy effect (4)1.1.2.Sluggish diffusion effect (5)1.1.3.Severe lattice-distortion effect (6)1.1.4.Cocktail effect (7)1.2.Key research topics (9)1.2.1.Mechanical properties compared with other alloys (10)1.2.2.Underlying mechanisms for mechanical properties (11)1.2.3.Alloy design and preparation for HEAs (11)1.2.4.Theoretical simulations for HEAs (12)2.Thermodynamics (12)2.1.Entropy (13)2.2.Thermodynamic considerations of phase formation (15)2.3.Microstructures of HEAs (18)3.Kinetics and alloy preparation (23)3.1.Preparation from the liquid state (24)3.2.Preparation from the solid state (29)3.3.Preparation from the gas state (30)3.4.Electrochemical preparation (34)4.Properties (34)4.1.Mechanical behavior (34)4.1.1.Mechanical behavior at room temperature (35)4.1.2.Mechanical behavior at elevated temperatures (38)4.1.3.Mechanical behavior at cryogenic temperatures (45)4.1.4.Fatigue behavior (46)4.1.5.Wear behavior (48)4.1.6.Summary (49)4.2.Physical behavior (50)4.3.Biomedical,chemical and other behaviors (53)5.Serrations and deformation mechanisms (55)5.1.Serrations for HEAs (56)5.2.Barkhausen noise for HEAs (58)5.3.Modeling the Serrations of HEAs (61)5.4.Deformation mechanisms for HEAs (66)6.Glass formation in high-entropy alloys (67)6.1.High-entropy effects on glass formation (67)6.1.1.The best glass former is located at the eutectic compositions (67)6.1.2.The best glass former is the composition with dense atomic packing (67)6.1.3.The best glass former has high entropy of mixing (67)6.2.GFA for HEAs (68)6.3.Properties of high-entropy BMGs (70)7.Modeling and simulations (72)7.1.DFT calculations (73)7.2.AIMD simulations (75)7.3.CALPHAD modeling (80)8.Future development and research (81)Y.Zhang et al./Progress in Materials Science61(2014)1–9338.1.Fundamental understanding of HEAs (82)8.2.Processing and characterization of HEAs (83)8.3.Applications of HEAs (83)9.Summary (84)Disclaimer (85)Acknowledgements (85)References (85)1.IntroductionRecently,high-entropy alloys(HEAs)have attracted increasing attentions because of their unique compositions,microstructures,and adjustable properties[1–31].They are loosely defined as solid solution alloys that contain more thanfive principal elements in equal or near equal atomic percent (at.%)[32].Normally,the atomic fraction of each component is greater than5at.%.The multi-compo-nent equi-molar alloys should be located at the center of a multi-component phase diagram,and their configuration entropy of mixing reaches its maximum(R Ln N;R is the gas constant and N the number of component in the system)for a solution phase.These alloys are defined as HEAs by Yeh et al.[2], and named by Cantor et al.[1,33]as multi-component alloys.Both refer to the same concept.There are also some other names,such as multi-principal-elements alloys,equi-molar alloys,equi-atomic ratio alloys,substitutional alloys,and multi-component alloys.Cantor et al.[1,33]pointed out that a conventional alloy development strategy leads to an enor-mous amount of knowledge about alloys based on one or two components,but little or no knowledge about alloys containing several main components in near-equal proportions.Theoretical and experi-mental works on the occurrence,structure,and properties of crystalline phases have been restricted to alloys based on one or two main components.Thus,the information and understanding are highly developed on alloys close to the corners and edges of a multi-component phase diagram,with much less knowledge about alloys located at the center of the phase diagram,as shown schematically for ternary and quaternary alloy systems in Fig.1.1.This imbalance is significant for ternary alloys but becomes rapidly much more pronounced as the number of components increases.For most quater-nary and other higher-order systems,information about alloys at the center of the phase diagram is virtually nonexistent except those HEA systems that have been reported very recently.In the1990s,researchers began to explore for metallic alloys with super-high glass-forming ability (GFA).Greer[29]proposed a confusion principle,which states that the more elements involved,the lower the chance that the alloy can select viable crystal structures,and thus the greater the chanceand quaternary alloy systems,showing regions of the phase diagram thatand relatively less well known(white)near the center[33].4Y.Zhang et al./Progress in Materials Science61(2014)1–93solid-solutions even though the cooling rate is very high,e.g.,alloys of CuCoNiCrAlFeTiV,FeCrMnNiCo, CoCrFeNiCu,AlCoCrFeNi,NbMoTaWV,etc.[1,2,12–14].The yield strength of the body-centered cubic(BCC)HEAs can be rather high[12],usually compa-rable to BMGs[12].Moreover,the high strength can be kept up to800K or higher for some HEAs based on3d transition metals[14].In contrast,BMGs can only keep their high strength below their glass-transition temperature.1.1.Four core effectsBeing different from the conventional alloys,compositions in HEAs are complex due to the equi-molar concentration of each component.Yeh[37]summarized mainly four core effects for HEAs,that is:(1)Thermodynamics:high-entropy effects;(2)Kinetics:sluggish diffusion;(3)Structures:severe lattice distortion;and(4)Properties:cocktail effects.We will discuss these four core effects separately.1.1.1.High-entropy effectThe high-entropy effects,which tend to stabilize the high-entropy phases,e.g.,solid-solution phases,werefirstly proposed by Yeh[9].The effects were very counterintuitive because it was ex-pected that intermetallic compound phases may form for those equi-or near equi-atomic alloy com-positions which are located at the center of the phase diagrams(for example,a monoclinic compound AlCeCo forms in the center of Al–Ce–Co system[38]).According to the Gibbs phase rule,the number of phases(P)in a given alloy at constant pressure in equilibrium condition is:P¼Cþ1ÀFð1-1Þwhere C is the number of components and F is the maximum number of thermodynamic degrees of freedom in the system.In the case of a6-component system at given pressure,one might expect a maximum of7equilibrium phases at an invariant reaction.However,to our surprise,HEAs form so-lid-solution phases rather than intermetallic phases[1,2,4,17].This is not to say that all multi-compo-nents in equal molar ratio will form solid solution phases at the center of the phase diagram.In fact, only carefully chosen compositions that satisfy the HEA-formation criteria will form solid solutions instead of intermetallic compounds.The solid-solution phase,according to the classical physical-metallurgy theory,is also called a ter-minal solid solution.The solid-solution phase is based on one element,which is called the solvent,and contains other minor elements,which are called the solutes.In HEAs,it is very difficult to differentiate the solvent from the solute because of their equi-molar portions.Many researchers reported that the multi-principal-element alloys can only form simple phases of body-centered-cubic(BCC)or face-cen-tered-cubic(FCC)solid solutions,and the number of phases formed is much fewer than the maximum number of phases that the Gibbs phase rule allows[9,23].This feature also indicates that the high en-tropy of the alloys tends to expand the solution limits between the elements,which may further con-firm the high-entropy effects.The high-entropy effect is mainly used to explain the multi-principal-element solid solution. According to the maximum entropy production principle(MEPP)[39],high entropy tends to stabilize the high-entropy phases,i.e.,solid-solution phases,rather than intermetallic phases.Intermetallics are usually ordered phases with lower configurational entropy.For stoichiometric intermetallic com-pounds,their configurational entropy is zero.Whether a HEA of single solid solution phase is in its equilibrium has been questioned in the sci-entific community.There have been accumulated evidences to show that the high entropy of mixing truly extends the solubility limits of solid solution.For example,Lucas et al.[40]recently reported ab-sence of long-range chemical ordering in equi-molar FeCoCrNi alloy that forms a disordered FCC struc-ture.On the other hand,it was reported that some equi-atomic compositions such as AlCoCrCuFeNi contain several phases of different compositions when cooling slowly from the melt[15],and thus it is controversial whether they can be still classified as HEA.The empirical rules in guiding HEA for-mation are addressed in Section2,which includes atomic size difference and heat of mixing.Y.Zhang et al./Progress in Materials Science61(2014)1–935 1.1.2.Sluggish diffusion effectThe sluggish diffusion effect here is compared with that of the conventional alloys rather than the bulk-glass-forming alloys.Recently,Yeh[9]studied the vacancy formation and the composition par-tition in HEAs,and compared the diffusion coefficients for the elements in pure metals,stainless steels, and HEAs,and found that the order of diffusion rates in the three types of alloy systems is shown be-low:Microstructures of an as-cast CuCoNiCrAlFe alloy.(A)SEM micrograph of an etched alloy withBCC and ordered BCC phases)and interdendrite(an FCC phase)structures.(B)TEMplate,70-nm wide,a disordered BCC phase(A2),lattice constant,2.89A;(B-b)aphase(B2),lattice constant,2.89A;(B-c)nanoprecipitation in a spinodal plate,7nm(B-d)nanoprecipitation in an interspinodal plate,3nm in diameter,a disorderedarea diffraction(SAD)patterns of B,Ba,and Bb with zone axes of BCC[01[011],respectively[2].illustration of intrinsic lattice distortion effects on Bragg diffraction:(a)perfect latticewith solid solutions of different-sized atoms,which are expected to randomly distribute statistical average probability of occupancy;(c)temperature and distortion effectsY.Zhang et al./Progress in Materials Science61(2014)1–937 the intensities further drop beyond the thermal effect with increasing the number of constituent prin-cipal elements.An intrinsic lattice distortion effect caused by the addition of multi-principal elements with different atomic sizes is expected for the anomalous decrease in the XRD intensities.The math-ematical treatment of this distortion effect for the modification of the XRD structure factor is formu-lated to be similar to that of the thermal effect,as shown in Fig.1.3[41].The larger roughness of the atomic planes makes the intensity of the XRD for HEAs much lower than that for the single-element solid.The severe lattice distortion is also used to explain the high strength of HEAs,especially the BCC-structured HEAs[4,12,23].The severe lattice-distortion effect is also related to the tensileFCC-structured HEAs have very low strength[7],which certainly cannot be explained by thelattice distortion argument.Fundamental studies in quantification of lattice distortion of HEAs are needed.1.1.4.Cocktail effectThe cocktail-party effect was usually used as a term in the acousticsfield,which have been used to describe the ability to focus one’s listening attention on a single talker among a mixture of conversa-tions and background noises,ignoring other conversations.For metallic alloys,the effect indicates that the unexpected properties can be obtained after mixing many elements,which could not be obtained from any one independent element.The cocktail effect for metallic alloys wasfirst mentioned by Ranganathan[42],which has been subsequently confirmed in the mechanical and physical properties [12,13,15,18,35,43].The cocktail effect implies that the alloy properties can be greatly adjusted by the composition change and alloying,as shown in Fig.1.4,which indicates that the hardness of HEAs can be dramat-ically changed by adjusting the Al content in the CoCrCuNiAl x HEAs.With the increase of the Al con-lattice constants of a CuCoNiCrAl x Fe alloy system with different x values:(A)hardnessconstants of an FCC phase,(C)lattice constants of a BCC phase[2].CoNiCrAl x Fe alloy system with different x values,the Cu-free alloy has lower hardness.CoCrCuFeNiAl x[15,45].Cu forms isomorphous solid solution with Ni but it is insoluble in Co,Cr and Fe;it dissolves about20at.%Al but also forms various stable intermetallic compounds with Al.Fig.1.6exhibits the hardness of some reported HEAs in the descending order with stainless steels as benchmark.The MoTiVFeNiZrCoCr alloy has a very high value of hardness of over800HV while CoCrFeNiCu is very soft with a value of less than200HV.Fig.1.7compares the specific strength,which yield strength over the density of the materials,and the density amongalloys,polymers and foam materials[5].We can see that HEAs have densitieshigh values of specific strength(yield strength/density).This is partiallyHEAs usually contain mainly the late transitional elements whoselightweight HEAs have much more potential because lightweightdensity of the resultant alloys will be lowered significantly.Fig.1.8strength of HEAs vs.Young’s modulus compared with conventional alloys.highest specific strength and their Young’s modulus can be variedrange of hardness for HEAs,compared with17–4PH stainless steel,Hastelloy,andYield strength,r y,vs.density,q.HEAs(dark dashed circle)compared with other materials,particularly structural Grey dashed contours(arrow indication)label the specific strength,r y/q,from low(right bottom)to high(left top).among the materials with highest strength and specific strength[5].Specific-yield strength vs.Young’s modulus:HEAs compared with other materials,particularly structural alloys.among the materials with highest specific strength and with a wide range of Young’s modulus[5].range.This observation may indicate that the modulus of HEAs can be more easily adjusted than con-ventional alloys.In addition to the high specific strength,other properties such as high hydrogen stor-age property are also reported[46].1.2.Key research topicsTo understand the fundamentals of HEAs is a challenge to the scientists in materials science and relatedfields because of lack of thermodynamic and kinetic data for multi-component systems in the center of phase diagrams.The phase diagrams are usually available only for the binary and ternary alloys.For HEAs,no complete phase diagrams are currently available to directly assist designing thealloy with desirable micro-and nanostructures.Recently,Yang and Zhang [28]proposed the Xparam-eter to design the solid-solution phase HEAs,which should be used combing with the parameter of atomic-size difference.This strategy may provide a starting point prior to actual experiments.The plastic deformation and fracture mechanisms of HEAs are also new because the high-entropy solid solutions contain high contents of multi-principal elements.In single principal-element alloys,dislo-cations dominate the plastic behavior.However,how dislocations interact with highly-disordered crystal lattices and/or chemical disordering/ordering will be an important factor responsible for plastic properties of HEAs.Interactions between the other crystal defects,such as twinning and stacking faults,with chemical/crystal disordering/ordering in HEAs will be important as well.1.2.1.Mechanical properties compared with other alloysFor conventional alloys that contain a single principal element,the main mechanical behavior is dictated by the dominant element.The other minor alloying elements are used to enhance some spe-cial properties.For example,in the low-carbon ferritic steels [47–59],the main mechanical properties are from the BCC Fe.Carbon,which is an interstitial solute element,is used for solid-solution strength-ened steels,and also to enhance the martensite-quenching ability which is the phase-transformation strengthening.The main properties of steels are still from Fe.For aluminum alloys [60]and titanium alloys [61],their properties are mainly related to the dominance of the elemental aluminum and tita-nium,respectively.Intermetallic compounds are usually based on two elements,e.g.,Ti–Al,Fe 3Al,and Fe 3Si.Interme-tallic compounds are typically ordered phases and some may have strict compositional range.The Burgers vectors of the ordered phases are too large for the dislocations to move,which is the main reason why intermetallic phases are usually brittle.However,there are many successful case studies to improve the ductility of intermetallic compound by micro-alloying,e.g.,micro-alloying of B in Ni 3Al[62],and micro-alloying of Cr in Fe 3Al [63,64].Amorphous metals usually contain at least three elements although binary metallic glasses are also reported,and higher GFA can be obtained with addition of more elements,e.g.,ZrTiCuNiBe (Vit-1),PdNiCuP,LaAlNiCu,and CuZrAlY alloys [65–69].Amorphous metals usually exhibit ultrahigh yield strength,because they do not contain conventional any weakening factors,such as dislocations and grain boundaries,and their yield strengths are usually three to five times of their corresponding crys-talline counterpart alloys.There are several models that are proposed to explain the plastic deforma-tion of the amorphous metal,including the free volume [70],a shear-transformation-zone (STZ)[71],more recently a tension-transition zone (TTZ)[72],and the atomic-level stress [73,74].The micro-mechanisms of the plastic deformation of amorphous metals are usually by forming shear bands,which is still an active research area till today.However,the high strength of amorphous alloys can be sustained only below the glass-transition temperature (T g ).At temperatures immediately above T g ,the amorphous metals will transit to be viscous liquids [68]and will crystallize at temperatures above the first crystallization onset temperature.This trend may limit the high-temperature applica-tions of amorphous metals.The glass forming alloys often are chemically located close to the eutectic composition,which further facilitates the formation of the amorphous metal–matrix composite.The development of the amorphous metal–matrix composite can enhance the room-temperature plastic-ity of amorphous metals,and extend application temperatures [75–78].For HEAs,their properties can be different from any of the constituent elements.The structure types are the dominant factor for controlling the strength or hardness of HEAs [5,12,13].The BCC-structured HEAs usually have very high yield strengths and limited plasticity,while the FCC-structured HEAs have low yield strength and high plasticity.The mixture of BCC +FCC is expected to possess balanced mechanical properties,e.g.,both high strength and good ductility.Recent studies show that the microstructures of certain ‘‘HEAs’’can be very complicated since they often undergo the spinodal decomposition,and ordered,and disordered phase precipitates at lower temperatures.Solution-strengthening mechanisms for HEAs would be much different from conventional alloys.HEAs usually have high melting points,and the high yield strength can usually be sustained to ultrahigh temperatures,which is shown in Fig.1.9for refractory metal HEAs.The strength of HEAs are sometimes better than those of conventional superalloys [14].10Y.Zhang et al./Progress in Materials Science 61(2014)1–931.2.2.Underlying mechanisms for mechanical propertiesMechanical properties include the Young’s modulus,yield strength,plastic elongation,fracture toughness,and fatigue properties.For the conventional one-element principal alloys,the Young’s modulus is mainly controlled by the dominant element,e.g.,the Young’s modulus of Fe-based alloys is about 200GPa,that of Ti-based alloys is approximately 110GPa,and that of Al-based alloys is about 75GPa,as shown in Fig.1.8.In contrast,for HEAs,the modulus can be very different from any of the constituent elements in the alloys [79],and the moduli of HEAs are scattered in a wide range,as shown in Fig.1.8.Wang et al.[79]reported that the Young’s modulus of the CoCrFeNiCuAl 0.5HEA is about 24.5GPa,which is much lower than the modulus of any of the constituent elements in the alloy.It is even lower than the Young’s modulus of pure Al,about 69GPa [80].On the other hand,this value needs to be verified using other methods including impulse excitation of vibration.It has been reported that the FCC-structured HEAs exhibit low strength and high plasticity [13],while the BCC-structured HEAs show high strength and low plasticity at room temperature [12].Thus,the structure types are the dominant factor for controlling the strength or hardness of HEAs.For the fracture toughness of the HEAs,there is no report up to date.1.2.3.Alloy design and preparation for HEAsIt has been verified that not all the alloys with five-principal elements and with equi-atomic ratio compositions can form HEA solid solutions.Only carefully chosen compositions can form FCC and BCC solid solutions.Till today there is no report on hexagonal close-packed (HCP)-structured HEAs.One reason is probably due to the fact that a HCP structure is often the stable structure at low tempera-tures for pure elements (applicable)in the periodic table,and that it may transform to either BCC or FCC at high temperatures.Most of the HEA solid solutions are identified by trial-and-error exper-iments because there is no phase diagram on quaternary and higher systems.Hence,the trial-and er-ror approach is the main way to develop high-performance HEAs.However,some parameters have been proposed to predict the phase formation of HEAs [17,22,28]in analogy to the Hume-Rothery rule for conventional solid solution.The fundamental thermodynamic equation states:G ¼H ÀTS ð1-2Þwhere H is the enthalpy,S is the entropy,G is the Gibbs free energy,and T is the absolute temperature.From Eq.(1-2),the TS term will become significant at high temperatures.Hence,preparing HEAs from the liquid and gas would provide different kinds of information.These techniques may include sput-Temperature dependence of NbMoTaW,VNbMoTaW,Inconel 718,and Haynes 230tering,laser cladding,plasma coating,and arc melting,which will be discussed in detail in the next chapter.For the atomic-level structures of HEAs,the neutron and synchrotron diffraction methods are useful to detect ordering parameters,long-range order,and short-range ordering[81].1.2.4.Theoretical simulations for HEAsFor HEAs,entropy effects are the core to their formation and properties.Some immediate questions are:(1)How can we accurately predict the total entropy of HEA phase?(2)How can we predict the phasefield of a HEA phase as a function of compositions and temperatures?(3)What are the proper modeling and experimental methods to study HEAs?To address the phase-stability issue,thermody-namic modeling is necessary as thefirst step to understand the fundamental of HEAs.The typical mod-eling techniques to address thermodynamics include the calculation of phase diagram(CALPHAD) modeling,first-principle calculations,molecular-dynamics(MD)simulations,and Monte Carlo simulations.Kao et al.[82]using MD to study the structure of HEAs,and their modeling efforts can well explain the liquid-like structure of HEAs,as shown in Fig.1.10.Grosso et al.[83]studied refractory HEAs using atomistic modeling,clarified the role of each element and their interactions,and concluded that4-and 5-elements alloys are possible to quantify the transition to a high-entropy regime characterized by the formation of a continuous solid solution.2.Thermodynamicsof a liquid-like atomic-packing structure using multiple elementsthird,fourth,andfifth shells,respectively,but the second and third shellsdifference and thus the largefluctuation in occupation of different atoms.2.1.EntropyEntropy is a thermodynamic property that can be used to determine the energy available for the useful work in a thermodynamic process,such as in energy-conversion devices,engines,or machines. The following equation is the definition of entropy:dS¼D QTð2-1Þwhere S is the entropy,Q is the heatflow,and T is the absolute temperature.Thermodynamic entropy has the dimension of energy divided by temperature,and a unit of Joules per Kelvin(J/K)in the Inter-national System of Units.The statistical-mechanics definition of entropy was developed by Ludwig Boltzmann in the1870s [85]and by analyzing the statistical behavior of the microscopic components of the system[86].Boltz-mann’s hypothesis states that the entropy of a system is linearly related to the logarithm of the fre-quency of occurrence of a macro-state or,more precisely,the number,W,of possible micro-states corresponding to the macroscopic state of a system:Fig.2.1.Illustration of the D S mix for ternary alloy system with the composition change[17].。

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。

International Journal of Minerals, Metallurgy and Materials Volume 25, Number 11, November 2018, Page 1294https:///10.1007/s12613-018-1682-8Corresponding author: Hamed Jamshidi Aval E-mail:h.jamshidi@nit.ac.ir© University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018Microstructural evolution and mechanical properties of friction stir-weldedC71000 copper–nickel alloy and 304 austenitic stainless steelHamed Jamshidi AvalDepartment of Materials and Industrial Engineering, Babol Noshirvani University of Technology, Shariati Avenue, Babol, 47148-71167, Iran(Received: 20 February 2018; revised: 29 May 2018; accepted: 11 June 2018)Abstract: Dissimilar joints comprised of copper–nickel and steel alloys are a challenge for manufacturers in modern industries, as these met-als are not thermomechanically or chemically well matched. The present study investigated the effects of tool rotational speed and linear speed on the microstructure and mechanical properties of friction stir-welded C71000 copper–nickel and 340 stainless steel alloys using a tungsten carbide tool with a cylindrical pin. The results indicated that a rotational-to-linear speed ratio of 12.5 r/mm did not cause any macro defects, whereas some tunneling defects and longitudinal cracks were found at other ratios that were lower and higher. Furthermore, chro-mium carbide was formed on the grain boundaries of the 304 stainless steel near the shoulder zone and inside the joint zone, directing carbon and chromium penetration toward the grain boundaries. Tensile strength and elongation percentages were 84% and 65% of the corresponding values in the copper–nickel base metal, respectively.Keywords: dissimilar friction stir welding; copper–nickel alloy; austenitic stainless steel; microstructure; mechanical properties1. IntroductionCopper–nickel alloys exhibit substantial corrosion resis-tance and anti-algae properties against biological sediments. Pure copper is not stable in oxygenated electrolytes, espe-cially in marine and chlorine ion environments where cop-per–nickel alloys are widely used, with copper as the main component [1]. The addition of nickel to copper improves the mechanical strength, durability, and resistance to corro-sion, abrasion, and cavitation in sea and polluted water. This alloy also exhibits significant stress corrosion cracking and corrosion fatigue resistance. Corrosion resistance can be in-creased by adding more nickel to copper–nickel alloys [2]. Since these alloys can be easily assembled and welded, they are prime candidates for plumbing systems, ship bodies, and other marine structures.Generally, stainless steel plays a major role in the modern world. Welding of austenite stainless steel is known for two important properties: maintenance of corrosion resistance and prevention of crack formation. Dissimilar joints of coatings on offshore platform insulators, achieved by different tech-niques, are among copper–nickel plate applications for corro-sion prevention. Other applications include the joining of copper–nickel pipes with steel flanges and/or direct joining of these pipes with steel pipes in marine industries [3].Nevertheless, welding of dissimilar metals is always challenging because of numerous factors. These include different melting points, thermal conductivity, and thermal expansion coefficients; galvanic corrosion; the high solidi-fication rate of molten copper; entry of molten copper into steel grain boundaries (especially in the heat-affected zone (HAZ)); formation of hot cracks; high copper oxidation at high temperatures; and type of filler metal [4−9]. It is essen-tial to select the appropriate filler metal and welding para-meters for dissimilar-metal fusion welding of copper–nickel and stainless steel alloys in order to reduce probable defects (e.g., cavitation and gas cavities).Recent developments in solid-state welding have made it an alternative to fusion welding. In comparison to other welding techniques, friction-stir welding is a solid-state technique with an outstanding combination of high speed, precision, and variety. Among different welding methods, friction-stir welding of dissimilar alloys is important due to the ability to join alloys with different properties. In addition,H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1295)different welding configurations in this method (e.g., lap and butt joints) make it applicable in different situations.Few studies have been conducted on friction-stir welding of dissimilar copper and stainless steel alloys. In this regard, Imani et al. [10] investigated a pure copper and stainless steel joint with a thickness of 3 mm using friction-stir weld-ing. It was found that the tool offset toward the copper side played a significant role in eliminating defects in the joints. In addition, Ramirez et al. [11] examined the effects of tool offset on the microstructure and mechanical properties of joints in friction-stir welding of pure copper and 316 stain-less steel with a thickness of 2 mm. They studied 0, 0.6, and 1.6 mm offsets relative to the joint interface. When a major part of the tool was on the steel side, the joint efficiency was 55% of copper base metal. Maximum joint efficiency, i.e., 87% of copper base metal, was reported in the 0.6-mm off-set relative to the joint interface.Furthermore, Najafkhani et al. [12] studied the joint of pure copper and 316 stainless steel with a thickness of 5 mm using friction-stir welding. In their study, all joints cracked from the heat-affected zone of the copper base metal. The highest tensile strength and elongation percentage were 220 MPa and 7%, respectively. In addition, Shamsujjoha et al. [13] studied the lap joint of pure copper with 1018 carbon steel using friction-stir welding. They found that the joining process at the interface was both mechanical and metallur-gical. Jafari et al. [14] also studied the friction-stir welding of pure copper and 304 stainless steel with a thickness of 3 mm. The heat input from the welding increased the grain size in the heat-affected zone and decreased joint ductility by increasing the number of welding passes.According to the literature, there are no studies on the friction-stir welding of copper−nickel and austenite stainless steel alloys. Accordingly, the present study investigated the effects of process parameters on the microstructure and me-chanical properties of friction stir-welded C71000 cop-per−nickel and 304 stainless steel alloys using a tungsten carbide tool with a cylindrical pin. Optical microscopy and scanning electron microscopy (SEM) were used to study the microstructure and detect the created phases in different zones. The mechanical properties of joints were also eva-luated by tensile and microhardness tests.2. ExperimentalIn the present study, C71000 copper−nickel and 304 aus-tenite stainless steel plates with thicknesses of 2 mm were used. Both plates were cut perpendicular to the rolled metal direction and had a dimension of 50 mm × 100 mm. The chemical compositions and mechanical properties of alloysare listed in Tables 1 and 2. The plates were welded in a buttjoint configuration. The copper−nickel alloy was on the re-treating side, while the stainless steel alloy was on the ad-vancing side. According to the literatures [10−11], 0.75 mmof the tool axis was offset to the copper−nickel alloy relativeto the joint interface. Fig. 1 shows the schematic of the tool offsetting procedure. A tungsten carbide tool with a cylin-drical pin with a height of 1.8 mm was used for welding. Fig.2 demonstrates the dimensions and geometry of the appliedtool in welding and Table 3 indicates the welding parameters.The present study selected two rotational speeds of 800 and1000 r/min and three linear speeds of 40, 60, and 80mm/min.Table 1. Chemical composition of alloy wt%ZnMnFeCrCuCNiAlloy0.90.010.05―Base0.0519.12C71000―1.20Base18.500.440.058.10SS304Table 2. Mechanical properties of alloysAlloyUltimate tensilestrength / MPaYieldstrength / MPaMicro-hardness,HV0.1Elonga-tion / %C71000338 110 9032 SS304585 210 15242Fig. 1.Schematic illustration of friction stir butt welding.Fig. 2. Tool geometrical characteristics.1296Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Table 3. Friction stir welding process parametersRotational-to-linear speed ratio / (r ⋅mm –1)Linear speed / (mm ⋅min –1)Rotational speed / (r ⋅min –1)Sample No.20.00 40 800113.33 60 800 2 10.00 80 800 3 25.00 40 1000 4 16.65 60 1000 5 12.508010006The samples were transversely cut for metallographic studies. A marble solution was used for etching the micro-structure on the stainless steel side after sanding and polish-ing, whereas a nitric acid and distilled water solution was used for the copper–nickel alloy. Scanning electron micro-scopy (SEM) and X-ray diffraction (XRD) were used to evaluate the joint interface and examine the distribution and type of intermetallic compounds in the joint cross section. The mechanical properties of the joint were investigated us-ing a tensile test according to the ASTM E8-M03 standard. The tensile test was carried out at a crosshead speed of 1 mm/min. A Vickers microhardness testing machine with a load of 3 N and test time of 15 s was used to evaluate the hardness distribution of a joint cross section.3. Results and discussion3.1. Weld appearanceThe qualitative test of the welded samples indicated that samples No. 1–5 had defects. Longitudinal cracks on the copper–nickel side or tunneling defects on the stainless steel side were observed in all defected samples. Fig. 3 shows the effects of rotational and linear speed on the appearance of welding samples No. 1, 3, 4, and 6, as representative sam-ples containing cracks, tunneling defects, and defect-free welds. It is generally difficult to explain the causes of de-fects in the samples; however, the heat input may be an in-fluential factor. Many researchers have introduced various analytical, numerical, and empirical models in order to evaluate the relationship between rotational and linear tool speed and heat input and to examine their effects on the temperature distribution in the friction-stir welding proce-dure. With a proper estimation, the rotational-to-linear speed ratio can be considered a measure of welding heat input.In this study, samples No. 3 and 4 received the least and most heat input, respectively. The lower temperature of sample No. 3 caused insufficient material flow into the stir zone. After the tool was moved forward, the flow of material stopped before arriving at the advancing side. Therefore, there was inadequate material to fill the hole on the advanc-ing side (stainless steel). The tunnel hole led to the loss of joint strength in this sample, and the two parts were easily separated.Fig. 3. Surface appearance of welded samples: (a) No. 4; (b) No. 1; (c) No. 6: (d) No. 3.Fig. 3 presents the longitudinal cracks because of a tunneling defect in sample No. 3. Fig. 4 shows the effect of welding heat input on longitudinal crack length. It can be seen that by increasing the rotational-to-linear speed ratio (increasing welding heat input), the maximum temperature in the joint increased, which led to the higher temperature gradient in the welded samples. The significant difference in thermal conductivity of copper–nickel and stainless steel al-loys (thermal conductivity of copper–nickel is 2.8 times higher than that of stainless steel) [15–16] at a high temper-ature gradient produced longitudinal cracks as a result of thermal stress in the joint. According to the visual inspection of welded samples, a tunneling defect developed in the joint at a rotational-to-linear speed ratio of less than 10.00 r/mm.H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1297)On the other hand, at the rotational-to-linear speed ratio of 13.33 r/mm or higher, longitudinal cracks were formed at the joint interface.3.2. Macrostructure and microstructureThe evaluation of mechanical and metallurgical proper-ties was only carried out for sample No. 6 because it had no defects. The macrostructure of the joint and microstructure of different zones are shown in Figs. 5 and 6. The micro-structure of stainless steel included austenite and δ-ferrite with a grain size of (40 ± 5) µm (Fig. 5(b)). Although the quantity of ferrite phase was not significant, the presence of δ-ferrite could improve the formation of the sigma phase inalloys during friction-stir welding [17].Fig. 4. Effect of the rotational-to-linear speed ratio on cracklength.Fig. 5. Optical images of different zones of sample No. 4: (a) macrostructure of welded sample No. 4; (b) base metal of AISI 304; (c) base metal of C71000; (d) TMAZ in AISI 304 side; (e) SZ in AISI304 side; (f) TMAZ in C71000 side as marked by zone I in (a).1298 Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 6. Microstructure of stir zone of sample No. 4: (a) microstructure of zone II in Fig. 5(a); (b) SZ in C71000 side; (c) SEM image of zone I in (a); (d) SEM image of zone II in (a).The copper–nickel microstructure had a grain size of (50 ± 4) µm and an average particle size of (10 ± 3) µm in the grain boundaries. The results of energy dispersive X-ray spectroscopy (EDS) indicated that these particles were nick-el-rich oxides with iron and zinc (Fig. 7). The stir zone mostly consisted of copper–nickel alloy, which was likely due to the lower flow stress of copper–nickel alloy [18] and location of the main part of the tool on the copper–nickel side. Different behaviors of the two alloys in the etchant so-lution confirmed this finding.Fig. 7. Element mapping result of base metal C71000 alloy.As shown in Fig. 5(a), a steel layer was drawn from the advancing zone to the retreating zone (zone I). The joint cross section as a result of friction-stir procedure consisted of the stir zone (SZ), thermomechanically affected zoneH. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1299)(TMAZ), and heat-affected zone (HAZ). The mechanical behaviors of the welding materials, especially the welding zone hardness, were affected by high plastic deformation and high temperature during the friction-stir welding. The stir zone microstructure in the friction-stir weld had smaller and equiaxed grains in comparison with the base metal due to high plastic deformation and stir resulting from the tool pin.As presented in Fig. 5(f), grains in the thermomechani-cally affected zone on the copper–nickel side were elon-gated, which is exclusive to this zone [19]. The steel layer on the copper–nickel side contained recrystallized cop-per–nickel grains (Fig. 5(f)). On the other hand, the stir zone microstructure on the copper–nickel side contained equiaxed grains with a size of (15 ± 4) µm as a result of dynamic re-crystallization in this zone (Fig. 6(b)). The oxide particles observed in the copper–nickel base metal are shown in this figure. These particles were mainly at grain boundaries with a size of (5 ± 2) µm and prevented the growth of stir-zone grains.The EDS results showed that zinc and iron concentra-tions in the oxide particles increased (Figs. 7 and 8). The stir zone on the steel alloy side contained small recrystallized grains with a size of (5 ± 1) µm (Fig. 5(e)). Clearly, the grain size in the copper–nickel stir zone was greater than that of the steel-stir zone. The temperature and deformation rate in the friction-stir procedure had inverse effects on the grain size of the stir zone. In fact, an increase in the defor-mation rate led to a reduced grain size, and a rise in temper-ature increased the grain size in the stir zone [20].Fig. 8. Element mapping result of stir zone of C71000 side.The advancing side showed the highest temperature and deformation [21]. According to the stir zone microstructure results, the deformation effect was dominant on the steel side, and the grain size of stir zone reduced relative to the copper–nickel alloy side. On the contrary, elongated grains did not exist in the thermomechanically affected zone on the steel side (Fig. 5(d)). However, annealing twins were found across the base metal, whereas there were fewer twins in the thermomechanically affected zone of the steel. There were no twins in the stir zone on the steel side. An interesting point in the microstructure study was the occurrence of a specific layer-by-layer structure at the interface between copper–nickel and steel alloys near the tool shoulder (Fig. 6(a)). The SEM images of different zones in Fig. 6(a) are presented in Figs. 6(c) and 6(d).According to the line scan analysis presented in Fig. 9, the layer-by-layer structures consisted of copper-rich layers adjacent to iron-rich layers. Based on the comparison of the chemical composition of the copper-rich layer and cop-per–nickel base metal, this zone belonged to the cop-per–nickel base metal. However, the iron-rich layer did not match the chemical composition of steel base metal. The highest mass percentages of copper and chromium in the iron-rich layer were 9% and 30%, respectively. The iron-rich layer had a higher copper percentage, which in-creased to 31wt% in some layers.The high percentages of nickel and copper as austenite stabilizers could promote the formation of austenite phase. Generally, welding of austenite stainless steel can cause de-fects, including formation of the brittle phase, hot cracks, and carbide–chrome in grain boundaries. Copper, as an auste-nite-forming element, eliminates the δ-ferrite and sigma phases. Furthermore, the copper–nickel alloy limits the sigma phase by increasing the cooling rate from 600 to 800°C [22].1300Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 9. SEM image and line scans of chemical elements at the layer structure: (a, b) zone I in Fig. 6(a): (c, d) Fig. 5(f).The sigma phase is very hard and brittle. Its value in-creases by increasing the percentage of chromium, molyb-denum, and silicon, but decreases by increasing the nitrogen, nickel, and carbon contents. Prevention of sigma phase for-mation in stainless steel is difficult when the chrome per-centage is about 20wt%. When the chrome percentage is less than 20wt%, the sigma phase is not observable in auste-nite stainless steels. Due to the very low amount of chrome (up to 9wt%) in the layered structure, formation of sigma phase is not expected.Fig. 10 shows the XRD analysis of the iron-rich zone in the layered structure (point A in Fig. 6(c)); the austenite phase is the only existing phase in this zone. The high per-centage of nickel and copper prevented the formation of sigma phase as expected. The line scan analysis (Fig. 9) in-dicated that nickel concentrations reduced in layer bounda-ries but increased in the iron-rich layers due to nickel migra-tion from the interface to iron-rich layers.According to the EDS results (Fig. 11(a)) regarding point A in Fig. 6(c), the nickel and copper percentages were 24wt% and 21wt%, respectively, indicating the diffusion ofFig. 10. XRD pattern of iron rich layer structure.nickel and copper from the copper–nickel alloy at the inter-face of steel alloy due to the close proximity of this region to the tool shoulder and high temperature of the zone. Ac-cording to the EDS results (Fig. 11(b)), regardless of the in-creased percentage of copper and nickel in the grain boun-daries of the recrystallized zone on the steel side, the high percentage of chrome indicates the increased effect of this element by moving toward the stir zone of the stainless steel. Carbon present in the grain boundaries indicates chrome carbide formation at the joint interface near the tool shoulder.H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1301)The chemical compositions of these spots indicate that chrome and carbon move toward high-energy zones and form chrome carbide. Formation of carbide and a chrome-free zone around the grain boundary can severely degrade corro-sion resistance of the joint. Analysis of point C (Fig. 11(c)) indicates that this zone belongs to 304 stainless steel. The transient zone in the joint interface can affect the mechanical properties of the joint. Partial diffusion and formation of iron- and copper-rich layers, as shown in Fig. 9, are alsoobserved in zone I of Fig. 5(a).3.3. Hardness evaluationThe joint microhardness profile at the mid-thickness of the weld cross section is presented in Fig. 12. Hardness of the stir zone increases by moving from the steel base metal. According to the Hall-Petch equation, smaller grains have greater hardness; accordingly, hardness increases by de-creasing the grain size and increasing the particle boundary density. Hardness near the interface fluctuates considering the layer-by-layer structure. This structure produces impor-tant features, such as non-uniform hardness profiles and stress concentration zones. The stir zone on the cop-per–nickel side had a more uniform hardness profile and lower quantity. Hardness gradually decreased to the level of copper–nickel base metal by moving toward the cop-per–nickel base metal.3.4. Tensile properties and fractographyThe stress–strain curves for the base metals and joint are shown in Fig. 13. The yield strength and tensile strength of the joint are 103 MPa and 285 MPa, respectively, while elongation is 21%; these values are significantly lower than the corresponding values in the base metals. Tensile strength and elongation of joint were 84% and 65% of the corres-ponding values, respectively in the copper–nickel base metal. It should be noted that fracture occurred in the weld nugget and at the interface of steel and copper–nickel. The hardness profile shows sudden fluctuations, which cause stress con-centrations and joint strength degradation.Fig. 12. Microhardness profiles of cross-section of joint No. 6.Fig. 11. EDS analysis of points A (a), B (b), and C (c) in Fig. 6(c).1302 Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 13. Stress–strain curve of base metals and welded sample No. 6.The fractured cross section was investigated by SEM af-ter the tensile test. Fig. 14 shows the fractured section and SEM image. The SEM image of the fracture zone shows a brittle cleavage fracture, along with plastic deformation and small uniform holes on the surface. In the brittle cleavage fracture, the crack propagation corresponds to the successive and repeated breaking of atomic bonds along specific crys-tallographic planes. The fracture surface has a faceted tex-ture because of different orientations of the cleavage planes in the grains. In this type of fracture, no substantial plastic deformation occurs and the crack propagates very fast, nearly perpendicular to the direction of the applied stress. In the ductile fracture mode, spherical dimples correspond to microvoids initiating crack formation. Each dimple is half the size of the microvoid, which is formed and then sepa-rated during the fracture process. In the welded sample, brit-tle and ductile failures simultaneously occurred, which could be attributed to the transient zone (Fig. 12) and sud-den fluctuations in the hardness of the sample.Fig. 14. SEM image of fracture surface of the joint No. 6.4. ConclusionsThe present study investigated the friction-stir welding of C71000 and AISI304 stainless steel with a cylindrical pin tool and the following results were obtained.(1) Lack of proper material flow occurred as a result of low temperature at a rotational-to-linear speed ratio of 10 r/mm; therefore, there was not adequate material to fill the hole as the tool traveled forward on the advancing side (stainless steel). In case of rotational-to-linear speed ratio of greater than 20 r/mm, the high heat input produced a higher temperature gradient and resulted in the formation of longi-tudinal cracks as a result of thermal stress in the joint sec-tion.(2) The grain size on the copper–nickel side was larger than that of the stainless steel side. The stirring phenomena during friction-stir welding eliminated annealing twins in the stainless steel base metal and a uniform microstructure with small equiaxed grains formed in the stir zone. Tensile strength and elongation of joint were 84% and 65% of the corresponding values, respectively in the copper–nickel base metal. The fracture surface indicated brittle cleavage and plastic deformation behaviors.(3) Heat and plastic deformation caused element diffu-sion at copper- and iron-rich layers in the stir zone. Nickel and copper, as austenite stabilizers, led to the formation of austenite phase in the iron-rich layers. Chrome and carbon were transferred to grain boundaries, which were high-energy zones, and formed chrome carbide. The layer-by-layer structure and precipitation at the interface made the hardness profile non-uniform and formed possible stress concentra-tion zones.AcknowledgementThe author acknowledges the funding support of Babol Noshirvani University of Technology (No. BNUT/370167/97).References[1] M. Metikoš-Hukovi ć, R. Babi ć, I. Škugor, and Z. Gruba č,Copper-nickel alloys modified with thin surface films: Corro-sion behaviour in the presence of chloride ions, Corros. Sci., 53(2011), No. 1, p. 347.[2] M. Metikoš-Hukovi ć, R. Babi ć, I. Škugor Ron čevi ć, and Z.Gruba č, Corrosion resistance of copper–nickel alloy under fluid jet impingement, Desalination , 276(2011), No. 1-3, p. 228.[3] P. Carol, Corrosion and biofouling resistance evaluation of90-10 copper–nickel, Copper Development Association ,H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1303)2005, No. 63, p. 8.[4] S.G. Shiri, M. Nazarzadeh, M. Shariftabar, and M.S. Afarani,Gas tungsten arc welding of CP-copper to 304 stainless steel using different filler materials, Trans. Nonferrous Met. Soc.China, 22(2012), No. 12, p. 2937.[5] C.W. Yao, B.S. Xu, X.C. Zhang, J. Huang, J. Fu, and Y.X.Wu, Interface microstructure and mechanical properties of laser welding copper-steel dissimilar joint, Opt. Lasers Eng., 47(2009), No. 7-8, p. 807.[6] I. Magnabosco, P. Ferro, F. Bonollo, and L. Arnberg, An in-vestigation of fusion zone microstructures in electron beam welding of copper-stainless steel, Mater. Sci. Eng. A, 424(2006), No. 1-2, p. 163.[7] T.A. May and A.C. Spowage, Characterisation of dissimilarjoints in laser welding of steel–kovar, copper–steel and cop-per–aluminium, Mater. Sci. Eng. A, 374(2004), No. 1-2, p.224.[8] C. Roy, V.V. Pavanan, G. Vishnu, and P.R. Hari, M. Ariva-rasu, M. Manikandan, D. Ramkumar, and N. Arivazhagan, Characterization of metallurgical and mechanical properties of commercially pure copper and AISI 304 dissimilar weld-ments, Procedia Mater. Sci., 5(2014), p. 2503.[9] M. Velu and S. Bhat, Metallurgical and mechanical examina-tions of steel–copper joints arc welded using bronze and nickel-base superalloy filler materials, Mater. Des., 47(2013), p. 793.[10] Y. Imani, M.K. Besharati, and M. Guillot, Improving frictionstir welding between copper and 304L stainless steel, Adv.Mater. Res., 409(2012), p. 263.[11] A.J. Ramirez, D.M. Benati, and H.C. Fals, Effect of tool off-set on dissimilar Cu–AISI 316 stainless steel friction stir welding, [in] Proceeding of the Twenty-first International Offshore and Polar Engineering Conference, Maui, Hawaii, USA, 2011, p. 548.[12] A. Najafkhani, K. Zangeneh-Madar, and H. Abbaszadeh,Evaluation of microstructure and mechanical properties of friction stir welded copper/316L stainless steel dissimilarmetals, Int. J. ISSI, 7(2010), No. 2, p. 21.[13] M. Shamsujjoha, B.K. Jasthi, M. West, and C. Widener, Mi-crostructure and mechanical properties of FSW lap joint be-tween pure copper and 1018 mild steel using refractory metal pin tools, [in] Friction Stir Welding and Processing VII,TMS, San Antonio, Texas, 2013, p. 151.[14] M. Jafari, M. Abbasi, D. Poursina, A. Gheysarian, and B.Bagheri, Microstructures and mechanical properties of fric-tion stir welded dissimilar steel–copper joints, J. Mech. Sci.Technol., 31(2017), No. 3, p. 1135.[15] Copper Development Association Inc., Copper–NickelWelding and Fabrication, Copper Development Association Inc., McLean, Virginia [2013-02-01]. / applications/marine/cuni/fabrication/welding_and_fabrication.html[16] Smiths Metal Centres, 304/304L Stainless Steel Data Sheet,Smiths Metal Centres, Clerkenwell, London [2007-03-05]./datasheets.htm.[17] S.H.C. Park, Y.S. Sato, H. Kokawa, K. Okamoto, S. Hirano,and M. Inagaki, Rapid formation of the sigma phase in 304 stainless steel during friction stir welding, Scripta Mater.,49(2003), No. 12, p. 1175.[18] Y.V.R.K. Prasad, K.P. Rao, and S. Sasidhara, Hot WorkingGuide: A Compendium of Processing Maps, ASM Interna-tional, Materials Park, Ohio, 2015, p. 168.[19] Y. Sun and H. 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Microstructure and mechanical properties of ZrB 2–SiC nanocomposite ceramicQiang Liu,*Wenbo Han and Ping HuCenter for Composite Materials,Harbin Institute of Technology,Harbin 150001,ChinaReceived 28March 2009;accepted 30May 2009Available online 6June 2009A ZrB 2–SiC nanocomposite ceramic in which 20vol.%nanosized SiC powder was introduced into a ZrB 2matrix was fabricated by hot-pressing at 1900°C for 60min under a 30MPa uniaxed load.The composite microstructure showed intragranular nanostruc-tures that were peculiar to this material.Investigation of the mechanical properties revealed a flexural strength of 930±28MPa and a fracture toughness of 6.5±0.3MPa m 1/2.These improved mechanical properties were strongly dependent on the formation of the unusual intragranular nanostructures.Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Intragranular nanostructure;Mechanical properties;Microstructure;Fracture toughness;NanocompositeUltrahigh-temperature ceramics (UHTCs),suchas borides and carbides,were developed in the 1960s [1].Among UHTCs,zirconium diboride (ZrB 2)is a material of particular interest because of its excellent combination of high melting point,low theoretical den-sity,high electrical conductivity,good chemical inert-ness and superb wear resistance.These properties make it an attractive candidate for high-temperature applications such as refractory materials in foundries,electrical devices,nozzles and armor [2].Moreover,ZrB 2could be used for super-high-temperature struc-tural applications in aerospace [3,4].Its low mechanical properties,however,have long prevented this material from being used in a wide range of applications.Its sus-ceptibility to brittle fracture can lead to unexpected cat-astrophic failure,therefore its mechanical properties must be improved before the potential applications of ZrB 2can be fully realized.The introduction of a second phase of particles has been a successful strategy for improving the mechanical properties of monolithic diboride ceramics.With this aim,introduction of SiC particles [3–6]into ZrB 2yields a ZrB 2–SiC composite ceramic that is far stronger than monolithic ZrB 2.As a rule,however,improvement of mechanical properties is limited by the micro-sized par-ticles of the second phase.The mechanical properties of ceramics can be signif-icantly improved by introducing nanosized ceramic par-ticles into the ceramic-matrix grains or grain boundaries.The most significant achievements with this approach have been reported by Niihara and Nakahira [7–9],who first revealed that an introduction of 5vol.%of nanosized SiC particles into Al 2O 3increased the room-temperature strength of the composite from 350MPa to $1.0GPa (three-point flexure,30mm span).Similar improvements in strength have since been achieved in Al 2O 3–Si 3N 4,MgO–SiC and Si 3N 4–SiC composite systems.Materials constructed by these types of approaches are termed nanocomposite ceramics.At this point in time,however,there have been few attempts to create nanocomposite ceramics out of ZrB 2–SiC.Moreover,the effects of the composite micro-structure on the mechanical properties of ZrB 2–SiC nanocomposite ceramics have never been documented.Therefore,the aim of the present study was to investi-gate the microstructural features and effects on mechan-ical properties of a ZrB 2–SiC nanocomposite ceramic.The starting powders used in this study were:ZrB 2powder (Northwest Institute for Non-ferrous Metal Re-search,China),average particle size 2l m (>99%);and nanosized b -SiC powder (Kaier Nanotechnology Devel-opment Co.Ltd,China),average particle size 30nm (>98%).The nanosized SiC powder was first dispersed in ethanol,with 1h of ultrasonication.Then the powder mixture ZrB 2plus 20vol.%nanosized SiC particles were ball-milled using ZrO 2ball media and ethanol at1359-6462/$-see front matter Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.scriptamat.2009.05.041*Corresponding author.Tel./fax:+8645186402382;e-mail:dqz0402@Available online at Scripta Materialia 61(2009)690–692/locate/scriptamat180rpm for 12h.All ball-milling was performed in polyethylene bottles.After mixing,the resulting slurry was dried by rotary evaporation and then screened.The obtained powder mixtures were hot-pressed at 1900°C for 60min at a uniaxial pressure of 30MPa in Ar atmosphere.The microstructure of hot-pressed specimens was ob-served by using scanning electron microscopy (SEM,FEI Sirion,Holland)and transmission electron micros-copy (TEM,Hitachi H-9000,Japan)with an X-ray en-ergy dispersive spectroscopy (EDS,EDAX Inc.)analyzer attachment.Flexural strength (r )was tested in three-point bending on 3Â4Â36mm bars,using a 30-mm span and a crosshead speed of 0.5mm min À1.Each specimen was ground and polished with diamond slurries to a 1-l m finish.The edges of all the specimens were chamfered to minimize the effect of stress concen-tration resulting from machining flaws.Fracture tough-ness (K IC )was evaluated by a single-edge notched-beam test with a 16-mm span and a crosshead speed of 0.05mm min À1using 2Â4Â22mm test bars on the same jig used for the flexural strength.All flexural bars were fabricated with the tensile surface perpendicular to the hot-pressing direction.A minimum of five specimens was tested in each experimental condition.Figure 1shows the typical microstructural morphol-ogies of the ZrB 2–SiC nanocomposite ceramic under SEM (Fig.1a)and TEM (Fig.1b).As shown in Figure 1a,a number of submicron SiC particles (dark contrast)are located along the grain boundaries of the ZrB 2(gray contrast).Some smaller SiC particles also appear inside the ZrB 2grains (indicated by arrows);these are termed intragranular nanostructures.Higher magnification examination of the ZrB 2–SiC nanocomposite ceramic by TEM (Fig.1b)showed that the intragranular SiC particles (indicated by arrows)were approximately 100nm in size.The formation of the intragranular nanostructures was dependent on both the migration speed of ZrB 2ma-trix grain boundary and the migration speed of the SiC second phase [10,11].When the migration speed of the second phase was less than that of the matrix grain boundary,the nanosized SiC particles tended to be trapped within the ZrB 2grains during sintering.The fine ZrB 2particles would then coalesce around them,form-ing the intragranular nanostructures.Figure 2shows that the specimen fracture surface used for testing fracture toughness exhibited the typicalcharacteristics of a transgranular fracture.In monolithic ZrB 2ceramic,the predominant fracture mode would have been intergranular [12].There are two possible interpretations for this difference in fracture mode.The first is that the intergranular SiC particles in the ZrB 2–SiC nanocomposite ceramic were firmly bonded to the ZrB 2/ZrB 2interfaces.This rigid bonding could then have suppressed intergranular fracture [13].The other explanation is that there are differences in relaxation of the tensile residual stress around the SiC particles located between the intergranular and intra-granular.Because of the different thermal expansion coefficients between SiC and ZrB 2,a large internal stress will be generated during cooling after sintering.Assum-ing that a SiC particle is spherical,an internal tension will occur in a tangential direction to the ZrB 2matrix around the SiC particle.This will cause a crack to al-ways propagate towards the SiC particle.The internal tangential tension also would be relaxed by lattice and grain-boundary diffusion around the intragranular and intergranular particles,respectively.However,the tem-perature at which the grain-boundary diffusion is acti-vated would be lower than that required by lattice diffusion,thus the internal tangential tension around the intergranular SiC particles would be further relaxed during cooling.As a result,the internal tangential ten-sion around the intragranular SiC particles of the sin-tered body would always be greater than that around the intergranular particles.This would lead to a fracture surface that would always tend to be characteristic of a transgranular fracture.Thus,it is the intragranular nanostructures that predominantly induce the trans-granular fracture characteristic of the ZrB 2–SiC nano-composite ceramic.Examination of the mechanical properties of the ZrB 2–SiC nanocomposite ceramic revealed a fracture toughness that ranged from 6.4to 6.7MPa m 1/2.This represented an increase of approximately 83%over that of the monolithic ZrB 2(2.3–3.5MPa m 1/2)[2].In addi-tion,the flexural strength (920–945MPa)of this nano-composite ceramic was also significantly higher than that recently reported for the monolithic ZrB 2($565MPa)[4].The formation of the intragranular nanostructures appeared to play an important role in the improved mechanical properties of the ZrB 2–SiC nanocomposite ceramic,especially its increased fracture toughness and flexural strength.In order to investigate effects of the intragranular nanostructures on the mechanical properties oftheFigure 1.Typical microstructural morphologies of the ZrB 2–SiC nanocomposite ceramic:(a)SEM image of the sample and (b)TEM image of thesample.Figure 2.SEM image of the fracture surface of the ZrB 2–SiC nanocomposite ceramic.Q.Liu et al./Scripta Materialia 61(2009)690–692691ZrB 2–SiC nanocomposite ceramic,it is necessary to investigate a crack propagation behavior in this mate-rial.Figure 3shows TEM micrographs of crack propa-gation behavior in the ZrB 2–SiC nanocomposite ceramic.It was evident that the crack had never propa-gated in a straight line,but had been deflected,selecting the neighboring particles (Fig.3a).As stated previously,this deflection was caused by thermal internal stress in this material.It can be also seen in Figure 3a that a crack has penetrated through an intragranular particle (indicated by black arrow).The possible reason for this case is that the cracked particle may be an agglomera-tion composed of many fine SiC particles.Because the bond strength of this agglomeration is not high enough,it tends to fracture when a crack meets this kind of par-ticle.However,for other intragranular particles (<100nm),neither crack penetration through the intra-granular particles nor propagation along the particle/matrix interfaces was evident (Fig.3b).This phenome-non indicates that the intragranular particles bridged the crack,pointing to the existence of a particle-bridging mechanism.Based on the experimental observation above,a spe-cific explanation for this effect is as follows.When a pri-mary crack meets an intragranular nanosized SiC particle,it is normally impeded and thus bows (Fig.3a).The bowing crack bypasses the impenetrable particles and instead interacts with neighboring cracks.At this point,the bridging particles firmly pin the cracks and further prevent the crack from extending.As a re-sult,only by increasing the crack extension force can the crack further extend.In other words,it is by means of the particle-bridging mechanism that the strength and toughness of the ZrB 2nanocomposite ceramic are signif-icantly improved.Besides the explanation mentioned above,there is an-other one for the improvement in strength.After the for-mation of the intragranular nanostructures,there are many sub-interfaces within the ZrB 2matrix grains that belong to the interfaces between intragranular particles and matrix grains.As stated previously,moreover,be-cause of the difference in thermal expansion coefficients between the ZrB 2matrix and the SiC second phase,a large number of microcracks were formed around the intragranular particles,as shown in Figure 4.The for-mation of the sub-interfaces and microcracks can cause the matrix grains to be at a potential differentiation state,corresponding to the further grain refining.Thisthen improves the strength of this material according to the Hall–Petch equation [10].As discussed above,it is concluded that the formation of intragranular nanostructure is the fundamental rea-son for the significant increase in the mechanical proper-ties of this nanocomposite ceramic.In conclusion,a hot-pressed ZrB 2–SiC nanocompos-ite ceramic was fabricated by introducing nanosized SiC powder into a ZrB 2matrix.the intragranular nanostruc-tures were peculiar to this ceramic-based composite and induced a transgranular fracture characteristic.The mechanical properties of this nanocomposite ceramic,especially its flexural strength and fracture toughness,were much higher than those of monolithic ZrB 2.It is believed that the formation of intragranular nanostructures is a main reason for the improvements in mechanical properties of the ZrB 2–SiC nanocompos-ite ceramic.Intragranular particle bridging is believed to be the predominant toughening mechanism imparting the improved characteristics to this material.This work was supported by the NSFC(10725207),the Research Fund for the Doctoral Pro-gram of Higher Education (24403037)and National Natural Science Fund for Outstanding Youths (24402052).[1]E.V.Clougherty,R.L.Pober,L.Kaufman,Trans.Met.Soc.AIME 242(1968)1077.[2]F.Monteverde,S.Guicciardi,A.Bellosi,Mater.Sci.Eng.A 346(2003)310.[3]A.L.Chamberlain,W.G.Fahrenholtz,G.E.Hilmas,D.T.Ellerby,J.Am.Ceram.Soc.87(2004)170.[4]F.Monteverde,C.Melandri,S.Gicciardi,Mater.Chem.Phys.100(2006)513.[5]F.Monteverde,Appl.Phys.A 82(2006)329.[6]S.S.Hwang,A.L.Vasiliev,N.P.Padture,Mater.Sci.Eng.A 464(2007)216.[7]K.Niihara, A.Nakahira,in:P.Vincentini (Ed.),Advanced Structural Inorganic Composites,Elsevier Sci-ence Publishers,Trieste,Italy,1990,pp.637–664.[8]K.Niihara,A.Nakahira,Ann.Chim.16(1991)479.[9]K.Niihara,J.Ceram.Soc.Jpn.99(1991)974.[10]W.D.Kingery,H.K.Bowen,D.R.Uhlmann,Introduc-tion to Ceramics,Wiley,1976.[11]C.M.Wang,J.Mater.Sci.30(1995)3222.[12]S.Q.Guo,J.M.Yang,H.Tanaka,Y.Kagawa,Compos.Sci.Technol.68(2008)3033.[13]I.A.Ovid’ko,A.G.Sheinerman,Scripta Mater.60(2009)627.Figure 3.TEM micrographs of crack propagation behavior in the ZrB 2–SiC nanocomposite ceramic:crack propagation is from upper right to lowerleft.Figure 4.TEM micrograph of microcracks around an intragranular particle.692Q.Liu et al./Scripta Materialia 61(2009)690–692。

Characterization of microstructure,mechanical properties and corrosion resistance of dissimilar welded joint between 2205duplex stainless steel and 16MnRShaogang Wang *,Qihui Ma,Yan LiCollege of Material Science and Technology,Nanjing University of Aeronautics and Astronautics,Nanjing 210016,Chinaa r t i c l e i n f o Article history:Received 23March 2010Accepted 10July 2010Available online 16July 2010Keywords:A.Ferrous metals and alloys D.WeldingF.Microstructurea b s t r a c tThe joint of dissimilar metals between 2205duplex stainless steel and 16MnR low alloy high strength steel are welded by tungsten inert gas arc welding (GTAW)and shielded metal arc welding (SMAW)respectively.The microstructures of welded joints are investigated using scanning electron microscope,optical microscope and transmission electron microscopy respectively.The relationship between mechanical properties,corrosion resistance and microstructure of welded joints is evaluated.Results indicate that there are a decarburized layer and an unmixed zone close to the fusion line.It is also indi-cated that,austenite and acicular ferrite structures distribute uniformly in the weld metal,which is advantageous for better toughness and ductility of joints.Mechanical properties of joints welded by the two kinds of welding technology are satisfied.However,the corrosion resistance of the weldment produced by GTAW is superior to that by SMAW in chloride solution.Based on the present work,it is con-cluded that GTAW is the suitable welding procedure for joining dissimilar metals between 2205duplex stainless steel and 16MnR.Ó2010Elsevier Ltd.All rights reserved.1.IntroductionDuplex stainless steel (DSS)consists of approximately equal amounts of austenite and ferrite,which results in the favorable mechanical properties and corrosion resistance.The higher strength properties allow weight savings,which reduce fabrication costs and enable lighter support structures to be used.The higher corrosion resistance,in particular against stress corrosion cracking,makes them preferably applied in certain environments such as chemical tankers,pressure vessels,pipes to heat exchangers,paper machines and ocean engineering [1–3].With the growing applica-tion of new materials and higher requirements for materials,a great need occurs for component or structure of dissimilar metals.However,the joining of dissimilar metals is generally more chal-lenging than that of similar metals,which is usually due to several factors such as the differences in chemical compositions and ther-mal expansion coefficients,resulting in different residual stresses situation across the different regions of weldments as well as the migration of carbon element from the steel with higher carbon content to the steel with relatively lower carbon content.If the welding process is not well controlled,some weld defects such as dilutions and cracks will generate in the weld metal and leadto great decrease of properties of the welded joint.There are some researches about failure analysis or mechanical performance for dissimilar metals joints.Ul-Hamid et al.[4]have addressed that carbon diffusion in the dissimilar joint between carbon steel pipe and type 304stainless steel elbows resulted in cracking after a rel-atively short period of usage.Lee et al.[5]have also reported creep–fatigue damage of dissimilar weldment of modified 9Cr–1Mo steel (ASME Grade 91)and 316L stainless steel in a liquid me-tal reactor.In order to overcome the technical problems and take full advantage of the properties of different metals,it is necessary to pay more attention to the joining of dissimilar metals,so as to produce high quality welded joints between them.At present,some investigations have been conducted on weld-ing of duplex stainless steel,almost all common fusion welding techniques can be used to weld duplex stainless steel through selecting appropriate filler metals and parameters such as heat in-put [6,7].Explosive welding can be thought as a feasible method to produce composite plates.Kaçar and Acarer [8]have addressed that explosive welding process can be used successfully for clad-ding duplex stainless steel on the vessel steels without losing prop-erties such as corrosion resistance and mechanical properties.However,compared to the welding of similar metals,there is lim-ited information about microstructure/property relationships in dissimilar material welds between duplex stainless steel and low alloy high strength steel.Increasing application of these steels will0261-3069/$-see front matter Ó2010Elsevier Ltd.All rights reserved.doi:10.1016/j.matdes.2010.07.012*Corresponding author.Tel.:+8602552112901;fax:+8602552112626.E-mail address:sgwang@ (S.G.Wang).require a better understanding of the mechanics associated with welding of dissimilar metals.Since GTAW and SMAW are widely employed in engineering application,in the current work,a few at-tempts have been made to produce dissimilar material welded joint between DSS and low alloy high strength steel.At the same time,some results are presented as reference for the practical welding of these types of dissimilar metals.2.Experimental material and procedureThe base metals employed in this presentation are duplex stain-less steel2205and low alloy high strength steel16MnR.The chem-ical compositions of base metals andfiller metals are given in3.Results and discussion3.1.Microstructure of welded jointsThe preparation of microstructure samples of dissimilar metals joint is much difficult.Therefore,special operation procedure should be used.Both of the weld metal(WM)and2205base metal are etched by aqua-regia.However,the bonding region at the side of16MnR is etched by5%nital solution alone,and16MnR base me-tal should be prevented from being etched by aqua-regia.The interfacial microstructure of16MnR–WM is shown in Fig.2.It is a region with about30l m width near the fusion line.The existing of this region can be attributed to the thermal conductivity of theTable1Chemical compositions of base metals andfiller metals(wt.%).Elements C Mn P S Si Cr Ni Mo NBase metal SAF22050.0160.820.0240.0010.3622.48 5.46 3.120.16 16MnR0.15 1.380.0160.0140.32––––Filler metal ER22090.013 1.540.0180.0070.4922.928.61 3.180.17 E22090.0260.900.0250.0020.9022.1010.00 2.840.18 832S.G.Wang et al./Materials and Design32(2011)831–837Moreover,some short rod-like carbides,granular carbides and is-land-like carbides are observed at higher amplification electron microscope,as shown in Fig.4c and d respectively.However,the result shows that no carbides such as M23C6or martensite are ob-served in the unmixed zone.Therefore,it can be concluded that the development of such a morphology is attributed to decomposition of pearlite at16MnR side and formation of Fe3C at the WM side. The decomposition model is shown in Fig.5.The optical micrograph of weld metal is shown in Fig.6.From Fig.6,the morphology of acicular ferrite in austenite matrix has been observed,which is characterized by large amount of austen-ite.However,in terms of ferrite content in the joint,there is not much variation between the two weld metals in welded joints A and B,and the ferrite volume fraction is only17.3%and14.5%(ob-improving joint crack resistance and reducing the inhomogeneous distribution of weld structure during multi-pass welding.Generally,the formation of martensite,M23C6(chromium car-bide),Cr2N and r phase depends on the base materials joined and welding conditions according to Refs.[15,16].Therefore,X-ray diffraction analysis is carried out on the weld metal and the re-sults are shown in Fig.7.There are only a and c phases in both of the weld metals,and no precipitation of M23C6(chromium car-bide),Cr2N or r phase is found in the weld metal,which is advan-tageous to mechanical properties and corrosion resistance of the joint.3.2.Mechanical properties834S.G.Wang et al./Materials and Design32(2011)831–837si-cleavage fracture,as shown in Fig.9d.Microhardness profile across the joint interface is shown inFig.10.The microhardness distributions of two kinds of welded joints are almost the same.Obviously,the hardness value of weld metal is higher than that of the16MnR base metal and the 16MnR HAZ.With the distance increasing away from interface, the microhardness values vary to a certain extent.The highest hardness values of the two joint interface are approximately 224HV and220HV respectively.It is because the carbon element migrates from the16MnR side to weld metal during welding due to the difference of chemical compositions between16MnR and weld metal.Similar result is reported by Kaçar and Acarer[8],studied the explosively welded joint between DSS andCorrosion behaviororder to evaluate the corrosion resistance of weldis sealed with A/B glue,leaving about10mmÂ10mm area,3.5%NaCl solution is used as corrosion solution,the sche-matic diagram is shown in Fig.11.Electrochemical corrosion test results of2205DSS base metal and weld metal are shown in Fig.12and Table4respectively. These samples display more or less similar behaviors in terms of In general,the higher the value is,the better corrosion resistance of the material is.Therefore,in3.5%NaCl solution,corrosion resis-tance order of the samples is:2205DSS BM>joint A>joint B.The pitting corrosion resistance of the DSS BM is much better compared to the two weld metals,as can be seen from the polari-zation plot.The DSS BM sample does not display any corrosion in 3.5%NaCl solution and there is no pit in the sample examined after the potentiodynamic cyclic scanning.And the good pitting resis-tance behaviors of weld metal are attributed to the addition of Cr,Ni,elements[19].The alloying element Cr could improve the stability of passivefilms,and the Ni would decrease the overall dis-solution rates of Fe and Cr[20].Moreover,the heat input of joint A is different from that of joint B,which affects the weld microstruc-ture and results in the difference of formation condition of metal surface passivefilm.Generally,thefiner the grain is,the more eas-ily the compact passivefilm forms.As a result,the corrosive ions cannot readily diffuse through the passivefilm and the metal pre-sents better corrosion resistance,so the joint A has better corrosion resistance compared to joint B.When welded joint is etched in chloride solution,defects gener-ated in the welding process(such as welding spatter or inclusion) possibly make it lose its ability to protect the surface passivefilm. As a result,a chromium-depleted zone appears around weld metal, which makes the surface activated,and the joint presents an ac-tive–passive behavior.The initiation sites for the pits are located at the ferrite–austenite grain boundaries and once formed they rapidly propagate from ferrite to austenite,as described in Ref.[21].It can be seen from Fig.13that ferrite grains are etched,leav-ing lots of grooves at the ferrite–austenite grain boundaries,and the remaining white strips are austenite.This selective localized corrosion is attributed to difference of the electrochemical poten-tial,caused by the ratio of biphase in weld metal.It is concluded that the austenite grains are by far more resistant to the chloride environment than that of the ferrite grains.4.ConclusionsThe investigation of welding between2205DSS and16MnR by GTAW and SMAW respectively reach the following conclusions:Fig.10.Hardness curves of16MnR–WM interface. Fig.12.Polarization curves of DSS BM and weld metals.Attribute to decomposition of pearlite at16MnR side and for-mation of Fe3C at the WM side,a decarburization layer and an un-mixed zone are observed at the interface of16MnR/WM.The microstructure of weld metal consists of austenite and acic-ular ferrite,and both of the two kinds of joints are characterized by a high content of austenite,which is beneficial to mechanical prop-erties and corrosion resistance.Sigma phase or M23C6intermetallic compounds are not observed in current case through analysis of XRD.The impact toughness of the weld metal is similar to that of 16MnR,but it is much higher than that of16MnR HAZ.The weld metal and16MnR HAZ welded by GTAW present a ductile mode of fracture,while the pattern of16MnR HAZ welded by SMAW be-longs to quasi-cleavage fracture.The fracture of the joint welded by GTAW occurs in the zone of 16MnR base metal,while the fracture position of the joint welded by SMAW is in16MnR HAZ.However,the average tensile strength of welded joints is582.4MPa,564.6MPa respectively.Both of them can meet the tensile strength requirements of engineering structure.The joint produced by SMAW has higher susceptibility to pitting corrosion in chloride solution than that of weldment produced by GTAW.Based on the present work,it is summarized that GTAW with filler metal ER2209is the suitable welding procedure for dissimilar metals joining between2205duplex stainless steel and16MnR in the practical application.References[1]Sieurin H,Sandström R.Austenite reformation in the heat-affected zone ofduplex stainless steel2205.Mater Sci Eng A2006;418:250–6.[2]Olsson J,Snis M.Duplex–A new generation of stainless steels for desalinationplants.Desalination2007;205:104–13.[3]Kurt B.The interface morphology of diffusion bonded dissimilar stainless steeland medium carbon steel couples.J Mater Process Technol2007;190:138–41.[4]Ul-Hamid A,Tawancy HM,Abbas NM.Failure of weld joints between carbonsteel pipe and304stainless steel elbows.Eng Fail Anal2005;12:181–91. 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