Effects of erbium on microstructure and mechanical properties of as-cast Mg-7Zn-3Al alloy

Ferroelectric and piezoelectric properties of tungsten substituted SrBi 2Ta 2O 9ferroelectric ceramicsIndrani Coondoo *,S.K.Agarwal a ,A.K.Jha ba Superconductivity and Cryogenics Division,National Physical Laboratory,Dr K.S.Krishnan Road,New Delhi 110012,India bDepartment of Applied Physics,Delhi College of Engineering,Bawana Road,Delhi 110042,India1.IntroductionDefects in crystals significantly influence physical and various other properties of materials [1].For instance,as it is well known,doping by other elements leads to significant changes in the electrical properties of silicon.Historically,‘‘defect engineering’’has been developed in the field of semiconducting materials such as compound semiconductors as well as in diamond,Si and Ge [2–4].Subsequently,the concept of defect engineering has been applied to other functional materials,and the significant improve-ment in material properties have been achieved in high transition-temperature superconductors [5],amorphous SiO 2[6],photonic crystals [7]and also in the field of ferroelectrics,such as BaTiO 3,Pb(Ti,Zr)O 3(PZT),etc.[8,9].Various structural and electrical properties of bismuth layer-structured ferroelectrics (BLSF)are also strongly affected on deviation from stoichiometric composi-tions and defects have been recognized as a crucially important factor [10–13].It has been found that in BLSF small changes in chemical composition result in significantly altered dielectric and ferroelectric properties including dielectric constant and remanent polarization.In SrBi 2Ta 2O 9(SBT)and SrBi 2Nb 2O 9(SBN),orthor-hombic structural distortions with non-centrosymmetric spacegroup A 21am cause spontaneous ferroelectric polarization (P s )along a axis [14,15].SBT,a member of the BLSF family,has occupied an important position among the Pb-free ferroelectric memory materials [16–18].Tungsten (W 6+)has recently been investigated as a dopant for bismuth titanates and lanthanum doped bismuth titanates,in which the remanent polarization was observed to enhance when a small amount of Ti 4+was substituted by W 6+[19,20].With the objective to improve structural,dielectric and ferroelectric proper-ties,the hexavalent tungsten (W 6+)was chosen as a donor cation for partial replacement of the pentavalent tantalum (Ta 5+)SBT.In this report,the effect of tungsten substitution in SBT (SBTW),on the microstructural,ferroelectric and piezoelectric properties is reported.The results including the improvement in polarization properties have been discussed.2.ExperimentalSamples of compositions SrBi 2(W x Ta 1Àx )2O 9(SBWT),with x =0.0,0.025,0.050,0.075,0.10and 0.20were synthesized by solid-state reaction method taking SrCO 3,Bi 2O 3,Ta 2O 5and WO 3(all from Aldrich)in their stoichiometric proportions.The powder mixtures were thoroughly ground and passed through sieve of appropriate size and then calcined at 9008C in air for 2h.The calcined mixtures were ground and admixed with about 1–1.5wt%polyvinyl alcohol (Aldrich)as a binder and then pressed at $300MPa into disk shaped pellets.The pellets were sintered at 12008C for 2h in air.Materials Research Bulletin 44(2009)1288–1292A R T I C L E I N F O Article history:Received 3October 2008Received in revised form 5December 2008Accepted 6January 2009Available online 15January 2009Keywords:A.CeramicsC.X-ray diffractionD.FerroelectricityA B S T R A C TTungsten substituted samples of compositions SrBi 2(W x Ta 1Àx )2O 9(x =0.0,0.025,0.050,0.075,0.10and 0.20)were synthesized by solid-state reaction method and studied for their microstructural,electrical conductivity,ferroelectric and piezoelectric properties.The X-ray diffractograms confirm the formation of single phase layered perovskite structure in the samples with x up to 0.05.The temperaturedependence of dc conductivity vis-a`-vis tungsten content shows a decrease in conductivity,which is attributed to the suppression of oxygen vacancies.The ferroelectric and piezoelectric studies of the W-substituted SBT ceramics show that the remanent polarization and d 33values increases with increasing concentration of tungsten up to x 0.05.Such compositions with low conductivity and high P r values should be excellent materials for highly stable ferroelectric memory devices.ß2009Elsevier Ltd.All rights reserved.*Corresponding author.Present address:Liquid Crystal Group,National Physical Laboratory,Dr K.S.Krishnan Road,New Delhi 110012,India.Tel.:+919810361727;fax:+911125170387.E-mail address:indrani_coondoo@ (I.Coondoo).Contents lists available at ScienceDirectMaterials Research Bulletinj o ur n a l h o m e p a g e :w w w.e l se v i e r.c om /l oc a t e /m a t r e sb u0025-5408/$–see front matter ß2009Elsevier Ltd.All rights reserved.doi:10.1016/j.materresbull.2009.01.001X-ray diffractograms of the sintered samples were recorded using a Bruker diffractometer in the range 108 2u 708with CuK a radiation.The sintered pellets were polished to a thickness of 1mm and coated with silver paste on both sides for use as electrodes and cured at 5508C for half an hour.Electrical conductivity was performed using Keithley’s 6517A Electrometer.The polarization–electric field (P –E )hysteresis measurements were done at room temperature using an automatic P –E loop tracer based on Sawyer–Tower circuit.Piezoelectric charge co-efficient d 33was measured using a Berlincourt d 33meter after poling the samples in silicone–oil bath at 2008C for half an hour under a dc electric field of 60–70kV/cm.3.Results and discussion3.1.Structural and micro-structural studiesThe phase formation and crystal structure of the ceramics were examined by X-ray diffraction (XRD),which is shown in Fig.1.The XRD patterns of the samples show the characteristic peaks of SBT.The peaks have been indexed with the help of a computer program–POWDIN [21]and the refined lattice parameters are given in Table 1.It is observed that a single phase layered perovskite structure is maintained in the range 0.0 x 0.05.Owing to the same co-ordination number i.e.6and the smallerionic radius of W (0.60A˚)in comparison to Ta (0.64A ˚),there is a high possibility of tungsten occupying the tantalum site.The observance of unidentified peak of very low intensity in the compositions with x >0.05indicates the solubility limit of W concentration in SBT.The unidentified peak is possibly due to tungsten not occupying the Ta sites in the structure as the intensity of this peak is observed to increase with tungsten content.Composition and sintering temperature influences the micro-structure such as grain growth and densification of the specimen,which in turn control other properties of the material [11,13].The effects of W substitution on the microstructure have been examined by SEM and the obtained micrographs are shown in Fig.2.It shows the microstructure of the fractured surface of the studied samples.It is clearly observed that W substitution has pronounced effect on the average grain size and homogeneity of the grains.Randomly oriented and anisotropic plate-like grains are observed in all the samples.It is also observed that the average grain size increases gradually with increasing W content.The average grain size in the sample with x =0.0is $2–3m m while that in the sample with x =0.20the size increases to $5–7m m.3.2.Electrical studiesThe electrical conductivity of ceramic materials encompasses a wide range of values.In insulators,the defects w.r.t.the perfect crystalline structure act as charge carriers and the consideration of charge transport leads necessarily to the consideration of point defects and their migration [22].Many mechanisms were put forward to explain the conductivity mechanism in ceramics.Most of them are approximately divided into three groups:electronic conduction,oxygen vacancies ionic conduction,and ionic and p-type mixed conduction [22].Intrinsic conductivity results from the movement of the component ions,whereas conduction resulting from the impurity ions present in the lattice is known as extrinsic conductivity.At low temperature region (ferroelectric phase),the conduction is dominated by the extrinsic conduction,whereas the conduction at the high-temperature paraelectric phase ($300–7008C)is dominated by the intrinsic ionic conduction [23,25].Fig.3shows the temperature dependence of dc conductivity (s dc )for the undoped and doped SBT samples.The curves show that the conductivity increases with temperature.This is indicative of negative temperature coefficient of resistance (NTCR)behavior,a characteristic of dielectrics [22].It is observed in Fig.3that throughout the temperature range,the dc conductivity of the doped samples are nearly two to three orders lower than that of the undoped sample.Two predominant conduction mechanisms indicated by slope changes in the two different temperature regions are observed in Fig.3.Such changes in the slope in the vicinity of the ferro-paraelectric transition region have been observed in other ferroelectric materials as well [23,24].In addition,it is also observed (Table 2)that the activation energy calculated using the Arrhenius equation [22]in the paraelectric phase increase from $0.80eV for the undoped sample to $2eV for the doped samples.The X-ray photoemission spectroscopic study has confirmed that when Bi 2O 3evaporates during high-temperature processing,vacancy complexes are formed in the (Bi 2O 2)2+layers [26].As a result,defective (Bi 2O 2)2+layers are inherently present in SBT.The undoped SBT shows n-type conductivity,since when oxygen vacancies are created,it leaves behind two trapped electrons [27]:O o !12O 2"þV o þ2e 0(1)where O o is an oxygen ion on an oxygen site,V o is a oxygen vacant site and e 0represents electron.The conductivity in the perovskites can be described as an ordered diffusion of oxygen vacancies [28].Their motion is manifested by enhanced ionic conductivity associated with an activation energy value of $1eV [26].These oxygen vacancies can be suppressed by addition of donors,since the donor oxide contains more oxygen per cation than the host oxide it replaces [29].It has been reported that conductivity in Bi 4Ti 3O 12(BIT)can be significantly decreased,up to three orders of magnitude with the addition of donors,such as Nb 5+and Ta 5+at the Ti 4+sites [23,30].A few other studies on layered perovskites have also reported a decrease inconductivityFig.1.XRD patterns of SrBi 2(W x Ta 1Àx )2O 9samples sintered at 12008C.Table 1Lattice parameters of SrBi 2(W x Ta 1Àx )2O 9samples.Concentration of W a (A ˚)b (A ˚)c (A ˚)0.0 5.5212 5.513924.92230.025 5.5214 5.520225.10790.05 5.5217 5.519925.05850.075 5.5191 5.504525.05670.10 5.5142 5.506125.0850.205.51335.493925.0861I.Coondoo et al./Materials Research Bulletin 44(2009)1288–12921289with addition of donors [23,24,31].In the present study,the Ta 5+-site substitution by W 6+in SBT can be formulated using a defect chemistry expression as WO 3þV o!Ta 2O 512W Ta þ3O o (2)It shows that the oxygen vacancies are reduced upon the substitution of donor W 6+ions for Ta 5+ions.Hence,it is reasonable to believe that the conductivity in SBT is suppressed by donor addition.As per the above discussion,the high s dc observed in the undoped SBT (Fig.3)can be attributed to the motion of oxygen vacancies.As already discussed,the doped samples show reduced conductivity because the transport phenomena involving oxygen vacancies are greatly reduced.The high E a value of $1.75–2eVcorresponding to the high-temperature region in the doped ceramics is consistent with the fact that in the donor-doped materials,the ionic conduction reduces [32].The activation energy E a in the low temperature ferroelectric region (Table 2)corre-sponds to extrinsic conduction.At lower temperatures the extrinsic conductivity results from the migration of impurity ions in the lattice.Some of these impurities may also be associated with lattice defects.Pure SBT has large number of Schottky defects (oxygen vacancies)in addition to impurity ions whereas in the doped samples,due to charge neutrality,there is relatively less content of oxygen vacancies.Thus,in the doped samples the conductivity in the low temperature region is largely due to the impurity ions only.This explains the high activation energy in pure SBT in the low temperature region compared to doped samples (Table 2).In the high-temperature region,the value of E a in the doped samples is observed to increase with W concentration up to x =0.05but beyond that,it decreases (Table 2).The decrease in the activation energy for samples with x >0.05suggests an increase in the concentration of mobile charge carriers [33].This observation can be ascribed to the existence of multiple valence states of tungsten.Since tungsten is a transitional metal element,the valence state of W ions in a solid solution most likely varies from W 6+to W 4+depending on the surrounding chemical environment [34].When W 4+are substituted for the Ta 5+sites,oxygen vacancies would be created,i.e.one oxygen vacancy would be created for every two tetravalent W ions entering the crystal structure,whichFig.3.Variation of dc conductivity with temperature in SrBi 2(W x Ta 1Àx )2O 9samples.Fig.2.SEM micrographs of fractured surfaces of SrBi 2(W x Ta 1Àx )2O 9samples with (a)x =0.0,(b)x =0.025,(c)x =0.050,(d)x =0.075,(e)x =0.10and (f)x =0.20Table 2Activation energy (E a )in the high-temperature paraelectric region and low temperature ferroelectric region;Curie temperature (T c )in SrBi 2(W x Ta 1Àx )2O 9samples.Concentration of W E a (high temp.)(eV)E a (low temp.)(eV)T c (8C)0.00.790.893110.025 1.920.593080.05 1.960.543250.075 1.940.543380.10 1.860.573680.201.740.54390I.Coondoo et al./Materials Research Bulletin 44(2009)1288–12921290explains the increase in the concentration of mobile charge carriers which ultimately results in an decrease in the E a beyond x>0.05. Hence it is reasonable to conclude that W ions in the SBWT exists as a varying valency state,i.e.at lower doping concentration they exist in hexavalent state(W6+)and at a higher doping concentra-tion,they tend to exist in lower valency states[8].The P–E loops of SrBi2(Ta1Àx W x)2O9are shown in Fig.4.It is observed that W-doping results in formation of well-defined hysteresis loops.Fig.5shows the compositional dependence of remanent polarization(2P r)and the coercivefield(2E c)of SrBi2(Ta1Àx W x)2O9samples.Both the parameters depend on W content of the samples.It is observed that2P rfirst increases with x and then decreases while2E cfirst decreases with x and then increases(Fig.5).The optimum tungsten content for maximum2P r ($25m C/cm2)is observed to be x=0.075.It is known that ferroelectric properties are affected by compositional modification,microstructural variation and lattice defects like oxygen vacancies[10,35,36].In hard ferroelectrics, with lower valent substituents,the associated oxide vacancies are likely to assemble in the vicinity of domain walls[37,38].These domains are locked by the defects and their polarization switching is difficult,leading to an increase in E c and decrease in P r[38]. On the other hand,in soft ferroelectrics,with higher valent substituents,the defects are cation vacancies whose generation in the structure generally increases P r.Similar observations have been made in many reports[38–41].Watanabe et al.[42]reported a remarkable improvement in ferroelectric properties in the Bi4Ti3O12ceramic by adding higher valent cation,V5+at the Ti4+ site.It has also been reported that cation vacancies generated by donor doping make domain motion easier and enhance the ferroelectric properties[43].Further,it is known that domain walls are relatively free in large grains and are inhibited in their movement as the grain size decreases[44].In the larger grains, domain motion is easier which results in larger P r.Also for the SBT-based system,it is known that with increase in the grain size the remanent polarization also increases[45,46].Based on the obtained results and above discussion,it can be understood that in the undoped SBT,the oxygen vacancies assemble at sites near domain boundaries leading to a strong domain pinning.Hence,as observed,well-saturated P–E loop for pure SBT is not obtained.But in the doped samples,the suppression of the oxygen vacancies reduces the pinning effect on the domain walls,leading to enhanced remanent polarization and lower coercivefield.Also,the increase in grain size in tungsten added SBT,as observed in SEM micrographs(Fig.2)contribute to the increase in polarization values.In the present study,the grain size is observed to increase with increasing W concentration.However, the2P r values do not monotonously increase and neither the E c decreases continuously with increasing W concentration(Fig.5). The variation of P r and E c beyond x>0.05,seems possibly affected by the presence of secondary phases(observed in XRD diffracto-grams),which hampers the switching process of polarization [47–50].Also,beyond x>0.05the increase in the number of charge carriers in the form of oxygen vacancies leads to pinning of domain walls and thus a reduction in the values of P r and increase in E c is observed.Fig.6shows the variation of piezoelectric charge coefficient d33 with x in the SrBi2(Ta1Àx W x)2O9.The d33values increases with increase in W content up to x=0.05.A decrease in d33values is observed in the samples with x!0.075.The piezoelectric coefficient,d33,increases from13pC/N in the sample with x=0.0to23pC/N in the sample with x=0.05.It is known that the major drawback of SBT is its relatively higher conductivity,which hinders proper poling[51].High resistivity is therefore important for maintenance of poling efficiency at high-temperature[52,53].The W-doped SBT samples show an electrical conductivity value up to three orders of magnitude lower than that of undoped sample(Fig.3).The positional variation of2P r and2E c in SrBi2(W x Ta1Àx)2O9samples.Fig.6.Variation of d33in SrBi2(W x Ta1Àx)2O9samples.Fig. 4.P–E hysteresis loops in SrBi2(W x Ta1Àx)2O9samples recorded at roomtemperature.I.Coondoo et al./Materials Research Bulletin44(2009)1288–12921291decrease in conductivity upon donor doping improve the poling efficiency resulting in the observed higher d33values.Moreover, since the grain size increases with W content in SBT,it is reasonable to believe that the increase in grain size will also contribute to the increase in d33values[54].The decrease in the value of d33for samples with x!0.075is possibly due to the presence of secondary phases as observed in diffractograms[1,51,55]and the increase in oxygen vacancies for samples with x>0.05.4.ConclusionsX-ray diffractograms of the samples reveal that the single phase layered perovskite structure is maintained in the samples with tungsten content x0.05.SEM micrographs reveal that the average grain size increases with increase in W concentration. The temperature dependence of the electrical conductivity shows that tungsten doping results in the decrease of conductivity by up to three order of magnitude compared to W free SBT.All the tungsten-doped ceramics have higher2P r than that of the undoped sample.The maximum2P r($25m C/cm2)is obtained in the composition with x=0.075.The reduced conductivity allows high-temperature poling of the doped samples.Such compositions with low loss and high P r values should be excellent materials for highly stable ferroelectric memory devices.The d33value is observed to increase with increasing W content up to x0.05.The value of d33 in the composition with x=0.05is$23pC/N as compared to$13 pC/N in the undoped sample.AcknowledgmentsThe authors sincerely thank Prof.P.B.Sharma,Dean,Delhi College of Engineering,India for his generous support and providing ample research infrastructure to carry out the research work.The authors are thankful to Dr.S.K.Singhal,Scientist, National Physical Laboratory,India for his fruitful discussion and suggestions.References[1]Y.Noguchi,M.Miyayama,K.Oikawa,T.Kamiyama,M.Osada,M.Kakihana,Jpn.J.Appl.Phys.41(2002)7062.[2]A.Bonaparta,P.Giannozzi,Phys.Rev.Lett.84(2000)3923.[3]S.Connell,E.Siderashaddad,K.Bharuthram,C.Smallman,J.Sellschop,M.Bos-senger,Nucl.Instrum.Methods B85(1994)508.[4]T.Derry,R.Spits,J.Sellschop,Mater.Sci.Bull.11(1992)249.[5]K.Salama,D.F.Lee,Supercond.Sci.Technol.7(1994)177.[6]H.Hosono,Y.Ikuta,T.Kinoshita,M.Hirano,Phys.Rev.Lett.87(2001)175501.[7]S.Noda,A.Chutinan,M.Imada,Nature407(1999)608.[8]S.Shannigrahi,K.Yao,Appl.Phys.Lett.86(2005)092901.[9]G.H.Heartling,nd,J.Am.Ceram.Soc.54(1971)1.[10]H.Watanabe,T.Mihara,H.Yoshimori,C.A.Paz De Araujo,Jpn.J.Appl.Phys.34(1995)5240.[11]T.Atsuki,N.Soyama,T.Yonezawa,K.Ogi,Jpn.J.Appl.Phys.34(1995)5096.[12]T.Noguchi,T.Hase,Y.Miyasaka,Jpn.J.Appl.Phys.35(1996)4900.[13]M.Noda,Y.Matsumuro,H.Sugiyama,M.Okuyama,Jpn.J.Appl.Phys.38(1999)2275.[14]R.E.Newnham,R.W.Wolfe,R.S.Horsey,F.A.D.Colon,M.I.Kay,Mater.Res.Bull.8(1973)1183.[15]A.D.Rae,J.G.Thompson,R.L.Withers,Acta Crystallogr.Sect.B:Struct.Sci.48(1992)418.[16]H.M.Tsai,P.Lin,T.Y.Tseng,J.Appl.Phys.85(1999)1095.[17]Y.Shimakawa,Y.Kubo,Y.Nakagawa,T.Kamiyama,H.Asano,F.Izumi,Appl.Phys.Lett.74(1999)1904.[18]Y.Noguchi,M.Miyayama,T.Kudo,Phys.Rev.B63(2001)214102.[19]J.K.Kim,T.K.Song,S.S.Kim,J.Kim,Mater.Lett.57(2002)964.[20]W.T.Lin,T.W.Chiu,H.H.Yu,J.L.Lin,S.Lin,J.Vac.Sci.Technol.A21(2003)787.[21]Wu E.,POWD,An interactive powder diffraction data interpretation and indexingprogram Ver2.1,School of Physical Science,Flinders University of South Australia, Bedford Park,S.A.JO42AU.[22]R.C.Buchanan,Ceramic Materials for Electronics:Processing,Properties andApplications,Marcel Dekker Inc.,New York,1998.[23]H.S.Shulman,M.Testorf,D.Damjanovic,N.Setter,J.Am.Ceram.Soc.79(1996)3124.[24]M.M.Kumar,Z.G.Ye,J.Appl.Phys.90(2001)934.[25]Y.Wu,G.Z.Cao,J.Mater.Res.15(2000)1583.[26]B.H.Park,S.J.Hyun,S.D.Bu,T.W.Noh,J.Lee,H.D.Kim,T.H.Kim,W.Jo,Appl.Phys.Lett.74(1999)1907.[27]C.A.Palanduz,D.M.Smyth,J.Eur.Ceram.Soc.19(1999)731.[28]C.R.A.Catlow,Superionic Solids&Solid Electrolytes,Academic Press,New York,1989.[29]M.V.Raymond,D.M.Symth,J.Phys.Chem.Solids57(1996)1507.[30]S.S.Lopatin,T.G.Lupriko,T.L.Vasiltsova,N.I.Basenko,J.M.Berlizev,Inorg.Mater.24(1988)1328.[31]M.Villegas,A.C.Caballero,C.Moure,P.Duran,J.F.Fernandez,J.Eur.Ceram.Soc.19(1999)1183.[32]Y.Wu,G.Z.Cao,J.Mater.Sci.Lett.19(2000)267.[33]B.H.Venkataraman,K.B.R.Varma,J.Phys.Chem.Solids66(2005)1640.[34]C.D.Wagner,W.M.Riggs,L.E.Davis,F.J.Moulder,Handbook of X-ray Photoelec-tron Spectroscopy,Perkin Elmer Corp.,Chapman&Hall,1990.[35]Y.Noguchi,I.Miwa,Y.Goshima,M.Miyayama,Jpn.J.Appl.Phys.39(2000)1259.[36]M.Yamaguchi,T.Nagamoto,O.Omoto,Thin Solid Films300(1997)299.[37]W.Wang,J.Zhu,X.Y.Mao,X.B.Chen,Mater.Res.Bull.42(2007)274.[38]T.Friessnegg,S.Aggarwal,R.Ramesh,B.Nielsen,E.H.Poindexter,D.J.Keeble,Appl.Phys.Lett.77(2000)127.[39]Y.Noguchi,M.Miyayama,Appl.Phys.Lett.78(2001)1903.[40]Y.Noguchi,I.Miwa,Y.Goshima,M.Miyayama,Jpn.J.Appl.Phys.39(2000)L1259.[41]B.H.Park,B.S.Kang,S.D.Bu,T.W.Noh,L.Lee,W.Joe,Nature(London)401(1999)682.[42]T.Watanabe,H.Funakubo,M.Osada,Y.Noguchi,M.Miyayama,Appl.Phys.Lett.80(2002)100.[43]S.Takahashi,M.Takahashi,Jpn.J.Appl.Phys.11(1972)31.[44]R.R.Das,P.Bhattacharya,W.Perez,R.S.Katiyar,Ceram.Int.30(2004)1175.[45]S.B.Desu,P.C.Joshi,X.Zhang,S.O.Ryu,Appl.Phys.Lett.71(1997)1041.[46]M.Nagata,D.P.Vijay,X.Zhang,S.B.Desu,Phys.Stat.Sol.(a)157(1996)75.[47]J.J.Shyu,C.C.Lee,J.Eur.Ceram.Soc.23(2003)1167.[48]I.Coondoo,A.K.Jha,S.K.Agarwal,Ferroelectrics326(2007)35.[49]T.Sakai,T.Watanabe,M.Osada,M.Kakihana,Y.Noguchi,M.Miyayama,H.Funakubo,Jpn.J.Appl.Phys.42(2003)2850.[50]C.H.Lu,C.Y.Wen,Mater.Lett.38(1999)278.[51]R.Jain,V.Gupta,A.Mansingh,K.Sreenivas,Mater.Sci.Eng.B112(2004)54.[52]I.S.Yi,M.Miyayama,Jpn.J.Appl.Phys.36(1997)L1321.[53]A.J.Moulson,J.M.Herbert,Electroceramics:Materials,Properties,Applications,Chapman&Hall,London,1990.[54]H.T.Martirena,J.C.Burfoot,J.Phys.C:Solid State Phys.7(1974)3162.[55]R.Jain,A.K.S.Chauhan,V.Gupta,K.Sreenivas,J.Appl.Phys.97(2005)124101.I.Coondoo et al./Materials Research Bulletin44(2009)1288–1292 1292。

Mg掺杂BaTiO3介电陶瓷中氧空位缺陷的EPR监控第３５卷⼀第１期⼀吉⼀林⼀化⼀⼯⼀学⼀院⼀学⼀报Ｖｏｌ.３５Ｎｏ.１⼀２０１８年１⽉ＪＯＵＲＮＡＬＯＦＪＩＬＩＮＩＮＳＴＩＴＵＴＥＯＦＣＨＥＭＩＣＡＬＴＥＣＨＮＯＬＯＧＹＪａｎ.⼀２０１８收稿⽇期:２０１７￣１０￣２３基⾦项⽬:国家⾃然科学基⾦项⽬(２１２７１０８４)?吉林省科技发展计划项⽬(２０１２１８２５)?长⽩⼭学者特聘教授⽀持计划(２０１５０４７)作者简介:郑永顺(１９９３￣)?男?吉林梅河⼝⼈?吉林化⼯学院研究⽣?主要从事⽆机介电陶瓷材料⽅⾯的研究.?通信作者:路⼤勇?Ｅ￣ｍａｉｌ:ｄｙｌｕ＠ｊｌｉｃｔ.ｅｄｕ.ｃｎ⼀⼀⽂章编号:１００７￣２８５３(２０１８)０１￣００４０￣０４Ｍｇ掺杂ＢａＴｉＯ３介电陶瓷中氧空位缺陷的ＥＰＲ监控郑永顺?路⼤勇?(吉林化⼯学院材料科学与⼯程研究中⼼?吉林吉林１３２０２２)摘要:在烧结温度Ｔｓ＝１４００?下?采⽤固相反应法制备(Ｂａ１－ｘＭｇｘ)ＴｉＯ３(ｘ＝０.０１５)和Ｂａ(Ｔｉ１－ｘＭｇｘ)Ｏ３(ｘ＝０.０１５)陶瓷.以电⼦顺磁共振(ＥＰＲ)技术作为关键技术?研究了陶瓷的点缺陷.结果表明:ＢＭＴ为六⽅和伪⽴⽅钛酸钡结构的混合相?ＢＴＭ为六⽅和四⽅钙钛矿结构的混合相.在低于－１００?的菱⽅相中?探测到与氧空位相关ｇ＝１.９５６的ＥＰＲ信号?且Ｂａ空位⼆Ｔｉ空位和Ｏ空位共存.关键词:镁掺杂钛酸钡?Ｘ射线衍射?电⼦顺磁共振?点缺陷中图分类号:Ｏ６１４.２３⽂献标志码:ＡＤＯＩ:１０.１６０３９/ｊ.ｃｎｋｉ.ｃｎ２２－１２４９.２０１８.０１.０１０⼀⼀在介电场中?ＢａＴｉＯ３基陶瓷是使⽤最⼴泛的电介质.镁离⼦(Ｍｇ２＋)具有固定的价态?通常与稀⼟离⼦作为共掺杂剂掺杂在ＢａＴｉＯ３中作为Ｘ７Ｒ型应⽤于多层陶瓷电容器(ＭＬＣＣ)中[１￣６].在１１５０?的较低的烧结温度(Ｔｓ)下?认为Ｍｇ２＋倾向于靠近晶粒的表⾯.对于Ｍｇ掺杂的ＢａＴｉＯ３陶瓷?Ｍｇ２＋被认为是替代ＢａＴｉＯ３中的Ｔｉ位的受体(Ｍｇ?Ｔｉ)[７￣９]?因为６￣ＣＮＭｇ２＋(０.７２?)与Ｔｉ４＋(０.６０５?)的离⼦尺⼨相近[１０].这⾥采⽤了Ｋｒ?ｇｅｒ和Ｖｉｎｋ提出的缺陷符号[１１].离⼦半径作为配位数(ＣＮ)的函数在表１中给出[１０].然⽽Ｍｇ２＋在ＢａＴｉＯ３￣ＭｇＴｉＯ３体系中代替Ｂａ位和Ｔｉ位的两性⾏为很少被考虑.根据报道?当Ｍｇ２＋代替Ｔｉ４＋位点时?每个Ｍｇ２＋离⼦(Ｍｇ?Ｔｉ)被认为伴随着⼀个氧空位(Ｖ七Ｏ)[７].Ｖ七Ｏ的存在与Ｍｇ掺杂的ＢａＴｉＯ３中的半导体⾏为有关?这归因于由受体掺杂的Ｍｇ２＋离⼦产⽣的氧空位的迁移率[７]具体请参照参考⽂献[７].在研究者的报道[７]中没有在Ｍｇ掺杂的ＢａＴｉＯ３中观察氧空位存在的直接证据.本实验?在烧结温度Ｔｓ＝１４００?下制备了ＢＭＴ￣ＢＴＭ陶瓷.采⽤电⼦顺磁共振(ＥＰＲ)技术对ＢＭＴ￣ＢＴＭ中的Ｖ七Ｏ进⾏检测?并提供了低温菱⽅相中存在Ｖ七Ｏ的证据.在Ｔｓ＝１４００?时?发现了Ｔｉ４＋与Ｔｉ３＋的还原以及在六⽅相中与Ｔｉ３＋(３ｄ１)有关的两种ＥＰＲ信号.表１⼀离⼦半径作为配位数(ＣＮ)的函数离⼦配位数(ＣＮ)半径(?)Ｂａ２＋１２１.６１Ｔｉ４＋６０.６０５Ｔｉ３＋６０.６７Ｍｇ２＋１２１.２３Ｍｇ２＋６０.７２１⼀实验部分１.１⼀试剂与仪器以ＢａＣＯ３(分析纯?上海帝阳)⼆ＴｉＯ２(９９.９％?上海帝阳)和ＭｇＯ(基准试剂?国药)粉末为初始原料.采⽤了国内丹东皓圆仪器有限公司的ＤＸ￣２７００型Ｘ射线衍射仪测定.以ＣｕＫα为射线源?扫描范围为２０??２θ?８５??步宽为０.０２??采样时间为３ｓ?管电压⼆管电流分别为４０ｋＶ和３０ｍＡ?采⽤步进扫描⽅式进⾏扫描.利⽤美国Ａｃｃｅｌｒｙｓ公司的ＭＳＭｏｄｅｌｉｎｇ软件的Ｒｅｆｌｅｘ模块并结合Ｓｍａ４Ｗｉｎｅ软件?对ＸＲＤ谱进⾏处理?除去ＣｕＫα２散射和衍射背底的贡献?进⾏晶体结构计算.采⽤法国ＪＹ公司的５３２ｎｍ激光ＬａｂＲＡＭＸｐｌｏＲＡ型拉曼光谱仪在室温下对样品进⾏测量?采⽤德国ＢｒｕｋｅｒＣｏｒｐｏｒａｔｉｏｎＡ３００￣１０/１２型电⼦顺磁共振仪对ＢＭＴ和ＢＴＭ样品进⾏ＥＰＲ谱测试?选⽤Ｘ波段频率为９.８４ＧＨｚ?在－１５０?温度点进⾏测定?测试范围为５００￣６５００Ｇ.ｇ值由ｈν０＝ｇβＨ关系计算.１.２⼀实验过程采⽤固相冷压陶瓷技术?按照名义分⼦式(Ｂａ１－ｘＭｇｘ)ＴｉＯ３(ＢＭＴ)⼆Ｂａ(Ｔｉ１－ｘＭｇｘ)Ｏ３(ｘ＝０.０１５)(ＢＴＭ)进⾏原料配⽐?将混合的原料在玛瑙研钵中研磨１ｈ后以３５０?/ｈ升⾄１１００?预烧５ｈ进⾏脱碳.⽤聚⼄烯醇⽔溶液(ＰＶＡ)粘合剂?将混合物在２００ＭＰａ压⼒下加压２ｍｉｎ压制成⽚.最后在１４００?下空⽓中烧结１２ｈ制备碳酸钡陶瓷.２⼀结果与讨论２.１⼀ＸＲＤ谱图分析图１显⽰的是ＢＭＴ￣ＢＴＭ陶瓷粉末在室温下测量的ＸＲＤ谱图与模拟的⽴⽅⼆四⽅⼆六⽅钛酸钡ＸＲＤ谱图的对⽐图.２θ/?图１⼀ＢＭＴ￣ＢＴＭ室温的室温ＸＲＤ谱图与模拟⽴⽅⼆四⽅⼆六⽅钛酸钡ＸＲＤ谱图的对⽐图观察图谱发现?ＢＭＴ在~４５?处的峰(图１ｂ)逐渐演变为四⽅晶系ＢａＴｉＯ３(空间群:Ｐ４ｍｍ)(ＪＣＰＤＳ卡号６￣５２６)的(００２)/(２００)峰与六⽅晶系ＢａＴｉＯ３(ＪＣＰＤＳ卡号３４￣１２９)(空间群:Ｐ６３/ｍｍｃ)的(２０４)峰(图１(ａ)所⽰)?即四⽅相和六⽅相共存在ＢＭＴ中?ＢＴＭ在~４５?处的峰(图１ｃ)逐渐演变为⽴⽅体ＢａＴｉＯ３(空间群:Ｐｍ３ｍ)ＪＣＰＤＳ卡号３１￣１７４)的(２００)与六⽅晶系ＢａＴｉＯ３的(２０４)峰(图１ａ所⽰)?即⽴⽅相和六⽅相共存在ＢＴＭ中.２.２⼀ＲＳ谱图分析室温下ＢＭＴ￣ＢＴＭ的拉曼光谱如图３(ａ)所⽰.四⽅ＢａＴｉＯ３具有四种常见的光学模式:Ａ１(ＴＯ２)?Ｂ１＋Ｅ(ＴＯ＋ＬＯ)?Ａ１(ＴＯ３)和Ａ１(ＬＯ３)＋Ｅ(ＬＯ３)?特征峰分别为~２６０?~３０５?~５２０?和~７２０ｃｍ－１.ＢＭＴ￣ＢＴＭ的拉曼散射实验阐明了三个事实:(１)~３０５ｃｍ－１的尖峰通常被认为是四⽅相的证明?并且在⽴⽅ＢａＴｉＯ３相中消失[１１￣１３].因此?３１０ｃｍ－１峰(图３(ａ)存在?说明ＢＴＭ中的⽴⽅相是伪⽴⽅.(２)对于ＢＭＴ与ＢＴＭ在１９５ｃｍ－１处观察到特征峰?这归因于伪⽴⽅相或四⽅晶相中存在正交变形[１４?１５].(３)Ｂａ(Ｔｉ１￣ｘＦｅｘ)Ｏ３－δ(ＢＴＦ)陶瓷具有完全六⽅相的结构(ａ＝５.７２４??ｃ＝１３.９９０??Ｖ０＝３９７.０４?３)(图３所⽰)?主要表现出三个强带?１５２?２０８和６４０ｃｍ－１(图３ｂ)?分别对应于Ｅ１ｇ?Ｅ１ｇ＋Ｅ２ｇ和Ａ１ｇ声⼦[１６].对于ＢＭＴ?这三个频带较弱?重叠在四⽅声⼦光谱(图３(ａ))?显⽰六⽅晶相和四⽅晶相共存?⽽对于ＢＴＭ?这三个频带较明显(图３(ａ))?显⽰六⽅晶相和伪⽴⽅晶相共存.Ｒａｍａｎｓｈｉｆｔ/ｃｍ－１图２⼀(ａ)在Ｔｓ＝１４００?烧结的ＢＭＴ￣ＢＴＭ和(ｂ)六⽅Ｂａ(Ｔｉ０.９５Ｆｅ０.０５)Ｏ３(ＢＴＦ)陶瓷的室温拉曼光谱插图显⽰了ＢＴＦ的ＸＲＤ图谱１４⼀⼀第１期郑永顺?等:Ｍｇ掺杂ＢａＴｉＯ３介电陶瓷中氧空位缺陷的ＥＰＲ监控⼀⼀⼀２.３⼀ＥＰＲ谱图分析图３为(ａ)ＢＭＴ与(ｂ)ＢＴＭ在－１５０?下测得的变温ＥＰＲ谱图.从图中可以看出在低于－１００?的菱⽅相中?探测到与氧空位相关ｇ＝１.９５６的ＥＰＲ信号[１７]?并探测到与Ｔｉ空位相关ｇ＝２.００４的ＥＰＲ信号[１８]?与Ｂａ空位相关ｇ＝１.９７４的ＥＰＲ信号[２０].ｇ＝１.９３４和１.９４２处的两个Ｔｉ３＋相关信号与ＢＭＴ⼆ＢＴＭ中的六⽅相有关?因为Ｔｉ３＋相关信号只能在低温下观察到?这是由于通过低温有效地延长了⾃旋￣晶格弛豫时间(τ)[１９].所以推断出在ＢＭＴ与ＢＴＭ中Ｂａ空位⼆Ｔｉ空位和Ｏ空位共存.Ｈ/ＧＨ/Ｇ图３⼀(ａ)ＢＭＴ与(ｂ)ＢＴＭ在－１５０?下测得的变温ＥＰＲ谱图３⼀结⼀⼀论在Ｔｓ＝１４００?的烧结温度下?采⽤固相反应法制备(Ｂａ１－ｘＭｇｘ)ＴｉＯ３和Ｂａ(Ｔｉ１－ｘＭｇｘ)Ｏ３(ｘ＝０.０１５)陶瓷.通过ＸＲＤ测试结果表明ＢＭＴ为六⽅和伪⽴⽅钛酸钡结构的混合相?ＢＴＭ为六⽅和四⽅钙钛矿结构的混合相?通过电⼦顺磁共振(ＥＰＲ)技术测试的结果表明ＢＭＴ与ＢＴＭ陶瓷在低于－１００?的菱⽅相中?探测到与氧空位相关ｇ＝１.９５６的ＥＰＲ信号?且Ｂａ空位⼆Ｔｉ空位和Ｏ空位共存.参考⽂献:[１]⼀Ｈ.Ｋｉｓｈｉ.?Ｎ.Ｋｏｈｚｕ.?Ｊ.Ｓｕｇｉｎｏ?ｅｔａｌ.Ｔｈｅｅｆｆｅｃｔｏｆｒａｒｅ￣ｅａｒｔｈ(Ｌａ?Ｓｍ?Ｄｙ?ＨｏａｎｄＥｒ)ａｎｄＭｇｏｎｔｈｅｍｉｃｒｏ￣ｓｔｒｕｃｔｕｒｅｉｎＢａＴｉＯ３[Ｊ].Ｊ.Ｅｕｒ.Ｃｅｒａｍ.Ｓｏｃ.?１９９９?１９(６￣７):１０４３￣１０４６.[２]⼀Ｊ.Ｈ.Ｈｗａｎｇ.?Ｓ.Ｋ.Ｃｈｏｌ.?ａｎｄＹ.Ｈ.Ｈａｎ.Ｄｉｅｌｅｃｔｒｉｃｐｒｏｐ￣ｅｒｔｉｅｓｏｆＢａＴｉＯ３ｃｏｄｏｐｅｄｗｉｔｈＥｒ２Ｏ３ａｎｄＭｇＯ[Ｊ].Ｊｐｎ.Ｊ.Ａｐｐｌ.Ｐｈｙｓ.?２００１?４０(８):４９５２￣４９５５.[３]⼀Ｈ.Ｋｉｓｈｉ.?Ｙ.Ｍｉｚｕｎｏ.?Ｈ.Ｃｈａｚｏｎｏ.Ｂａｓｅ￣ｍｅｔａｌｅｌｅｃｔｒｏｄｅ￣ｍｕｌｔｉｌａｙｅｒｃｅｒａｍｉｃｃａｐａｃｉｔｏｒｓ:ｐａｓｔ?ｐｒｅｓｅｎｔａｎｄｆｕｔｕｒｅｐｅｒｓｐｅｃｔｉｖｅｓ[Ｊ].Ｊｐｎ.Ｊ.Ａｐｐｌ.Ｐｈｙｓ.?２００３?４２(１):１￣１５.[４]⼀Ｓ.Ｗａｎｇ.?Ｓ.Ｚｈａｎｇ.?Ｘ.Ｚｈｏｕ.?ｅｔａｌ.Ｅｆｆｅｃｔｏｆｓｉｎｔｅｒｉｎｇａｔｍｏｓｐｈｅｒｅｓｏｎｔｈｅｍｉｃｒｏｓｔｒｕｃｔｕｒｅａｎｄｄｉｅｌｅｃｔｒｉｃｐｒｏｐ￣ｅｒｔｉｅｓｏｆＹｂ/Ｍｇｃｏ￣ｄｏｐｅｄＢａＴｉＯ３ｃｅｒａｍｉｃｓ[Ｊ].Ｍａｔｅｒ.Ｌｅｔｔ.?２００５?５９(１９):２４５７￣２４６０.[５]⼀Ｃ.￣Ｙ.Ｃｈａｎｇ.?Ｗ.￣Ｎ.Ｗａｎｇ.?Ｃ.￣Ｙ.Ｈｕａｎｇ.ＥｆｆｅｃｔｏｆＭｇＯａｎｄＹ２Ｏ３ｄｏｐｉｎｇｏｎｔｈｅｆｏｒｍａｔｉｏｎｏｆｃｏｒｅ￣ｓｈｅｌｌｓｔｒｕｃｔｕｒｅｉｎＢａＴｉＯ３ｃｅｒａｍｉｃｓ[Ｊ].Ｊ.Ａｍ.Ｃｅｒａｍ.Ｓｏｃ.?２０１３?９６(８):２５７０￣２５７６.[６]⼀Ｃ.￣Ｈ.Ｋｉｍ.?Ｋ.￣Ｊ.Ｐａｒｋ.?Ｙ.￣Ｊ.Ｙｏｏｎ.?ｅｔａｌ.ＲｏｌｅｏｆＹｔｔｉｕｍａｎｄｍａｇｎｅｓｉｕｍｉｎｔｈｅｆｏｒｍａｔｉｏｎｏｆｃｏｒｅ￣ｓｈｅｌｌｓｔｒｕｃｔｕｒｅｏｆＢａＴｉＯ３ｇｒａｉｎｓｉｎＭＬＣＣ[Ｊ].Ｊ.Ｅｕｒ.Ｃｅｒａｍ.Ｓｏｃ.?２００８?２８(６):１２１３￣１２１９.[７]⼀Ｓ.Ｈ.Ｃｈａ.?Ｙ.Ｈ.Ｈａｎ.ＥｆｆｅｃｔｓｏｆＭｎｄｏｐｉｎｇｏｎｄｉｅｌｅｃｔｒｉｃｐｒｏｐｅｒｔｉｅｓｏｆＭｇ￣ｄｏｐｅｄＢａＴｉＯ３[Ｊ].Ｊ.Ａｐｐｌ.Ｐｈｙｓ.?２００６?１００(１０):１０４１０２.[８]⼀Ｍ.Ｄｏｎｇ.?Ｈ.Ｍｉａｏ.?Ｇ.Ｔａｎ.?ｅｔａｌ.ＥｆｆｅｃｔｓｏｆＭｇ￣ｄｏｐｉｎｇｏｎｔｈｅｍｉｃｒｏｓｔｒｕｃｔｕｒｅａｎｄｐｒｏｐｅｒｔｉｅｓｏｆＢａＴｉＯ３ｃｅｒａｍｉｃｓｐｒｅｐａｒｅｄｂｙｈｙｄｒｏｔｈｅｒｍａｌｍｅｔｈｏｄ[Ｊ].Ｊ.Ｅｌｅｃ￣ｔｒｏｃｅｒａｍ.?２００８?２１(１):５７３￣５７６.[９]⼀Ｓ.￣Ｈ.Ｙｏｏｎ.?Ｃ.Ａ.Ｒａｎｄａｌｌ.?Ｋ.￣Ｈ.Ｈｕｒ.Ｅｆｆｅｃｔｓｏｆａｃｃｅｐｔｏｒｃｏｎｃｅｎｔｒａｔｉｏｎｏｎｔｈｅｂｕｌｋｅｌｅｃｔｒｉｃａｌｃｏｎｄｕｃｔｉｏｎｉｎａｃｃｅｐｔｏｒ(Ｍｇ)￣ｄｏｐｅｄＢａＴｉＯ３[Ｊ].Ｊ.Ａｐｐｌ.Ｐｈｙｓ.?２０１０?１０７(１０):１０３７２１.[１０]Ｒ.Ｄ.Ｓｈａｎｎｏｎ..Ｒｅｖｉｓｅｄｅｆｆｅｃｔｉｖｅｉｏｎｉｃｒａｄｉｉａｎｄｓｙｓ￣２４⼀⼀吉⼀林⼀化⼀⼯⼀学⼀院⼀学⼀报⼀⼀２０１８年⼀⼀ｔｅｍａｔｉｃｓｔｕｄｉｅｓｏｆｉｎｔｅｒａｔｏｍｉｃｄｉｓｔａｎｃｅｓｉｎｈａｌｉｄｅｓａｎｄｃｈａｌｃｏｇｅｎｉｄｅｓ[Ｊ].ＡｃｔａＣｒｙｓｔａｌｌｏｇｒ.ＳｅｃｔＡ.?１９７６?３２(５):５１￣７６７.[１１]Ｄ.￣Ｙ.Ｌｕ.?Ｘ.￣Ｙ.Ｓｕｎ.?Ｍ.Ｔｏｄａ.Ａｎｏｖｅｌｈｉｇｈ￣ｋＹ５Ｖｂａｒｉｕｍｔｉｔａｎａｔｅｃｅｒａｍｉｃｓｃｏ￣ｄｏｐｅｄｗｉｔｈｌａｎｔｈａｎｕｍａｎｄｃｅｒｉｕｍ[Ｊ].Ｊ.Ｐｈｙｓ.Ｃｈｅｍ.Ｓｏｌｉｄｓ.?２００７?６８(４):６５０￣６６４.[１２]Ｍ.Ｐ.Ｆｏｎｔａｎａ.?Ｍ.Ｌａｍｂｅｒｔ.Ｌｉｎｅａｒｄｉｓｏｒｄｅｒａｎｄｔｅｍｐｅｒ￣ａｔｕｒｅｄｅｐｅｎｄｅｎｃｅｏｆＲａｍａｎｓｃａｔｔｅｒｉｎｇｉｎＢａＴｉＯ３[Ｊ].ＳｏｌｉｄＳｔａｔｅＣｏｍｍｏｎ.?１９７２?１０(１):１￣４.[１３]Ａ.Ｍ.Ｑｕｉｔｔｅｔ.?Ｍ.Ｌａｍｂｅｒｔ.ＴｅｍｐｅｒａｔｕｒｅｄｅｐｅｎｄｅｎｃｅｏｆｔｈｅＲａｍａｎｃｒｏｓｓｓｅｃｔｉｏｎａｎｄｌｉｇｈｔａｂｓｏｒｐｔｉｏｎｉｎｃｕｂｉｃＢａＴｉＯ３[Ｊ].ＳｏｌｉｄＳｔａｔｅＣｏｍｍｏｎ.?１９７３?１２(１０):１０５３￣１０５５.[１４]Ｔ.Ｎａｇａｉ.?Ｋ.Ｉｉｊｉｍａ.?Ｈ.Ｊ.Ｈｗａｎｇ.?ｅｔａｌ.ＥｆｆｅｃｔｏｆＭｇＯｄｏｐｉｎｇｏｎｔｈｅｐｈａｓｅｔｒａｎｓｆｏｒｍａｔｉｏｎｓｏｆＢａＴｉＯ３[Ｊ].Ｊ.Ａｍ.Ｃｅｒａｍ.Ｓｏｃ.?２０００?８３(１):１０７￣１１２.[１５]Ｃ.Ｈ.Ｐｅｒｒｙ.?Ｄ.Ｂ.Ｈａｌｌ.ＴｅｍｐｅｒａｔｕｒｅｄｅｐｅｎｄｅｎｃｅｏｆｔｈｅＲａｍａｎｓｐｅｃｔｒｕｍｉｎＢａＴｉＯ３[Ｊ].Ｐｈｙｓ.Ｒｅｖ.Ｌｅｔｔ.?１９６５?１５(１７):７００￣７０２.[１６]Ｈ.Ｙａｍａｇｕｃｈｉ.?Ｈ.Ｕｗｅ.?Ｔ.Ｓａｋｕｄｏ.?ｅｔａｌ.Ｒａｍａｎ￣ｓｃａｔ￣ｔｅｒｉｎｇｓｔｕｄｙｏｆｔｈｅｓｏｆｔｐｈｏｎｏｎｍｏｄｅｓｉｎｈｅｘａｇｏｎａｌｂａｒ￣ｉｕｍｔｉｔａｎａｔｅ[Ｊ].Ｊ.Ｐｈｙｓ.Ｓｏｃ.Ｊｐｎ.?１９８７?５６(２):５８９￣５９５.[１７]彭研焱?路⼤勇.Ｔｂ在ＢａＴｉＯ３中的⾃补偿模式探索及介电性质[Ｊ].吉林化⼯学院学报?２０１６?３３(１):６１￣６４.[１８]刘婷婷?路⼤勇.Ｌｕ在ＢａＴｉＯ３中⾃的位占据研究[Ｊ].吉林化⼯学院学报?２０１６?３３(１):６５￣６８.[１９]Ｐ.Ｓ.Ｄｏｂａｌ.?Ａ.Ｄｉｘｉｔ.?Ｒ.Ｓ.Ｋａｔｉｙａｒ.?ｅｔａｌ.Ｍｉｃｒｏ￣ＲａｍａｎｓｔｕｄｙｏｆＢａ１－ｘＳｒｘＴｉＯ３ｃｅｒａｍｉｃｓ[Ｊ].Ｊ.ＲａｍａｎＳｐｅｃｔｒｏｓｃ.?２００１?３２(２):１４７￣１４９.ＥＰＲＭｏｎｉｔｏｒｉｎｇｏｆＯｘｙｇｅｎ￣ｖａｃａｎｃｙＤｅｆｅｃｔｓｉｎＭｇ￣ｄｏｐｅｄＢａＴｉｏ３ＤｉｅｌｅｃｔｒｉｃＣｅｒａｍｉｃｓＺＨＥＮＧＹｏｎｇ￣ｓｈｕｎ?ＬＵＤａ￣ｙｏｎｇ?(ＲｅｓｅａｒｃｈＣｅｎｔｅｒｆｏｒＭａｔｅｒｉａｌｓＳｃｉｅｎｃｅａｎｄＥｎｇｉｎｅｅｒｉｎｇ?ＪｉｌｉｎＩｎｓｔｉｔｕｔｅｏｆＣｈｅｍｉｃａｌＴｅｃｈｎｏｌｏｇｙ?ＪｉｌｉｎＣｉｔｙ１３２０２２?Ｃｈｉｎａ)Ａｂｓｔｒａｃｔ:(Ｂａ１－ｘＭｇｘ)ＴｉＯ３ａｎｄＢａ(Ｔｉ１－ｘＭｇｘ)Ｏ３－ｘ(ｘ＝０.０１５)ｃｅｒａｍｉｃｓｗｅｒｅｐｒｅｐａｒｅｄｂｙｔｈｅｓｏｌｉｄ￣ｓｔａｔｅｒｅａｃｔｉｏｎｍｅｔｈｏｄａｔａｓｉｎｔｅｒｉｎｇｔｅｍｐｅｒａｔｕｒｅｏｆＴｓ＝１４００?.Ｅｌｅｃｔｒｏｎｐａｒａｍａｇｎｅｔｉｃｒｅｓｏｎａｎｃｅ(ＥＰＲ)ｗａｓｅｍｐｌｏｙｅｄａｓａｋｅｙｔｅｃｈｎｉｑｕｅｔｏｓｔｕｄｙｔｈｅｐｏｉｎｔｄｅｆｅｃｔｓｏｆｃｅｒａｍｉｃｓ.ＴｈｅｒｅｓｕｌｔｓｓｈｏｗｔｈａｔＢＭＴｃｅｒａｍｉｃｅｘｈｉｂｉｔｅｄｔｈｅｍｉｘｅｄｐｈａｓｅｓｏｆｈｅｘａｇｏｎａｌａｎｄｐｓｅｕｄｏｃｕｂｉｃｐｅｒｏｖｓｋｉｔｅｓｔｒｕｃｔｕｒｅｓ?ＢＴＭｃｅｒａｍｉｃｅｘｈｉｂｉｔｅｄｔｈｅｍｉｘｅｄｐｈａｓｅｓｏｆｈｅｘａｇｏｎａｌａｎｄｔｅｔｒａｇｏｎａｌｐｅｒｏｖｓｋｉｔｅｓｔｒｕｃｔｕｒｅｓ.ＡｎＥＰＲｓｉｇｎａｌａｔｇ＝１.９５６?ｗｈｉｃｈｉｓａｓｓｏｃｉａｔｅｄｗｉｔｈｏｘｙｇｅｎ￣ｖａｃａｎｃｙｄｅｆｅｃｔｓ?ｗａｓｄｅｔｅｃｔｅｄｉｎｔｈｅｒｈｏｍｂｏｈｅｄｒａｌｐｈａｓｅｂｅｌｏｗＴ＝－１００?.ＴｈｒｅｅｔｙｐｅｓｏｆｐｏｉｎｔｄｅｆｅｃｔｓｏｆＢａ?Ｔｉ?Ｏｖａｃａｎｃｉｅｓｃｏｕｌｄｃｏｅｘｉｓｔｉｎｃｅｒａｍｉｃｓ.Ｋｅｙｗｏｒｄｓ:Ｍｇ￣ｄｏｐｅｄＢａＴｉＯ３?Ｘ￣ｒａｙｄｉｆｆｒａｃｔｉｏｎ?ｅｌｅｃｔｒｏｎｐａｒａｍａｇｎｅｔｉｃｒｅｓｏｎａｎｃｅ?ｐｏｉｎｔｄｅｆｅｃｔ３４⼀⼀第１期郑永顺?等:Ｍｇ掺杂ＢａＴｉＯ３介电陶瓷中氧空位缺陷的ＥＰＲ监控⼀⼀⼀。

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].。

稀土元素Gd对Al-Si-Mg铸造合金微观组织和力学性能的影响刘文祎;徐聪;刘茂文;肖文龙;马朝利【摘要】为了系统地研究稀土Gd对铸造Al-Si-Mg(A357)合金组织和性能的影响,采用OM,SEM,EPMA,XRD,DSC,T EM及拉伸实验等方法对不同Gd含量A357合金进行研究.结果表明:Gd的添加可以细化A357合金的晶粒并减小二次枝晶间距.此外,Gd可以有效地细化合金中的共晶硅,但是对片状共晶硅的形貌影响不大.晶粒和共晶硅的细化及二次枝晶间距的减小使添加Gd后的A357合金的力学性能有了显著的提高.其中,A357-0.5Gd(质量分数/％)合金热处理态抗拉强度为355M Pa,相对于未添加Gd元素的A357合金提高了37M Pa.当Gd质量分数为1.0％时,尽管组织得到进一步细化,但是大量粗大Al2 Si2 Gd第二相的形成导致了合金力学性能的下降.同时对Gd的细化机制进行探究,结合T EM分析结果可以推断,Gd变质处理后共晶硅上的孪晶密度并不足以引起共晶硅形貌的转变,使得Gd变质效果较弱.而Gd对共晶硅的细化作用可能与Gd增加成分过冷以及形成纳米相阻碍共晶硅生长有关.【期刊名称】《材料工程》【年(卷),期】2019(047)006【总页数】7页(P129-135)【关键词】铝硅合金;稀土Gd;细化;显微组织;力学性能【作者】刘文祎;徐聪;刘茂文;肖文龙;马朝利【作者单位】北京航空航天大学材料科学与工程学院空天先进材料与服役教育部重点实验室,北京100191;北京航空航天大学材料科学与工程学院空天先进材料与服役教育部重点实验室,北京100191;北京航空航天大学材料科学与工程学院空天先进材料与服役教育部重点实验室,北京100191;北京航空航天大学材料科学与工程学院空天先进材料与服役教育部重点实验室,北京100191;北京航空航天大学材料科学与工程学院空天先进材料与服役教育部重点实验室,北京100191【正文语种】中文【中图分类】TG146.2铸造铝硅合金由于具有优良的铸造性能，较高的比强度与韧性以及良好的抗疲劳性能和耐蚀性能，被广泛应用于航空、航天和军事等领域。

Material Sciences 材料科学, 2019, 9(9), 898-903Published Online September 2019 in Hans. /journal/mshttps:///10.12677/ms.2019.99111Effect of Aging Treatment onMicrostructure and MechanicalProperties of AlSi7Mg0.4 AlloysNa Zhao*, Suiqun Zhu, Yi Cao#SkyTeam Motor (Technology) Ltd., ShanghaiReceived: Sep. 4th, 2019; accepted: Sep. 23rd, 2019; published: Sep. 30th, 2019AbstractAlSi7Mg0.4 alloy is an aluminum alloy with strengthened heat treatment. In the paper, AlSi7Mg0.4 alloys with different aging treatment were investigated. The structure was characterized by X-ray diffraction (XRD) and metallographic microscope. Dynamic change characteristics of the sample with the temperature were tested through differential scanning calorimeter (DSC) method.Through tensile and fatigue tests, the results showed that the yield strength was decreased with aging temperature decrease, and the break elongation was increased with aging temperature de-crease. And the fatigue performance was better at lower aging temperature.KeywordsAl-Si Casting Alloy, AlSi7Mg, Aging Treatment, Texture, Mechanical Property时效处理对AlSi7Mg0.4合金组织结构和力学性能的影响赵娜*，朱随群，曹懿#天合汽车科技(上海)有限公司，上海收稿日期：2019年9月4日；录用日期：2019年9月23日；发布日期：2019年9月30日*第一作者。

精 密 成 形 工 程第16卷 第5期 48JOURNAL OF NETSHAPE FORMING ENGINEERING 2024年5月收稿日期：2024-01-29 Received ：2024-01-29基金项目：国家重点研发计划（2022YFB4602300）Fund ：National Key R&D Program of China (2022YFB4602300) 引文格式：段宇航, 王磊磊, 郝璐静, 等. 热处理时效时间对激光增材制造Al-Mg-Sc-Zr 合金拉伸性能的影响研究[J]. 精密成形工程, 2024, 16(5): 48-54.DUAN Yuhang, WANG Leilei, HAO Lujing, et al. Effect of Heat Treatment Aging Time on Tensile Properties of Laser Additive Manufactured Al-Mg-Sc-Zr Alloy[J]. Journal of Netshape Forming Engineering, 2024, 16(5): 48-54. *通信作者（Corresponding author ）热处理时效时间对激光增材制造Al-Mg-Sc-Zr合金拉伸性能的影响研究段宇航，王磊磊*，郝璐静，原帅超，赵艳秋，占小红（南京航空航天大学 材料科学与技术学院，南京 211106）摘要：目的 研究热处理时效时间对激光增材制造Al-Mg-Sc-Zr 合金微观组织与拉伸性能的影响，揭示微观组织与力学性能的内在关联机制。

方法 采用控制单一变量的试验方法进行时效热处理，设定保温温度为325 ℃，冷却方式为空冷，在不同保温时间（2、4、6、8 h ）下进行组织与性能共通性及差异性分析。

结果 经325 ℃时效热处理4 h 后，在激光增材制造Al-Mg-Sc-Zr 高强铝合金中形成了Al 3Sc 、Al 3(Sc,Zr)析出相，抗拉强度达到最大值486 MPa ，相较于未热处理，提升了21.8%，随着保温时间的进一步延长，析出相的高温停留时间变长，组织形核长大，Al 3Sc 、Al 3(Sc,Zr)强化相尺寸明显增大，最大尺寸可达0.6 μm 。

Journal of Materials Processing Technology 211(2011)750–758Contents lists available at ScienceDirectJournal of Materials ProcessingTechnologyj o u r n a l h o m e p a g e :w w w.e l s e v i e r.c o m /l o c a t e /j m a t p r o t ecPerformance of a cutting tool made of steel matrix surface nano-composite produced by in situ laser melt injection technologyO.Verezub a ,Z.Kálazi b ,A.Sytcheva c ,L.Kuzsella d ,G.Buza b ,N.V.Verezub e ,A.Fedorov f ,G.Kaptay g ,h ,∗aUni.Miskolc,Dep.Production Eng.,Egyetemvaros,3515Miskolc,HungarybBAY-ATI (RI for Materials Science and Engineering),Fehervari ut 130,Budapest,Hungary cBAY-NANO (RI for Nanotechnology),Dep.Nano-metrology,3515Miskolc,Egyetemvaros,E/7,Hungary dUni.Miskolc,Dep.Polimer Eng.,Egyetemvaros,3515Miskolc,Hungary eNational Technical University (“Kharkov Polytechnical Institute”),Frunze st.21,Dep.Integrated Technologies in Machine Building,Kharkov,Ukraine fARC -The Australian Reinforcing Company,518Ballarat Road,Sunshine,3020VIC,Melbourne,Australia gBAY-NANO (RI for Nanotechnology),Dep.Nano-composites and Uni.Miskolc,Egyetemvaros,E/7,3515Miskolc,Hungary hUniversity of Miskolc,Dep.Nanotechnology,Egyetemvaros,E/7,3515Miskolc,Hungarya r t i c l e i n f o Article history:Received 15May 2010Received in revised form 4December 2010Accepted 7December 2010Available online 15December 2010Keywords:Cutting toolMetal matrix nano-composite Laser processing Tool-lifea b s t r a c tSteel-matrix (105WCr6steel)surface nano-composites with (Ti,W)C micron-sized and (Fe,W)6C nano-sized carbide precipitates were produced by in situ laser melt injection technology with subsequent heat treatment.The microhardness of a 1mm thick nano-composite layer was found to be higher than that of the initial matrix.The machinability of the surface nano-composite by a cubic boron nitride (CBN)wheel was found lower,but still reasonable compared to the initial matrix.Cutting tools produced from our new nano-composite by the CBN wheel were found to have higher wear resistance,longer tool life and provided lower cutting forces against a C45steel workpiece compared to the initial matrix of the nano-composite.©2010Elsevier B.V.All rights reserved.1.IntroductionMaterial removal and machining processes play a key role in generating value-added activities to materials and machine parts since their introduction about 3centuries ago (Shaw,2005).Nev-ertheless,there is a constant interest in this field due to the high variety of newly developed materials (Biswas,2006)and superhard coatings (Veprek and Veprek-Heijman,2008)that can be used as cutting tools.The optimum combination of hard particles and duc-tile metallic matrices can lead to higher wear resistance.As this principle is widely recognized,particles reinforced metal matrix composites (MMCs)have been developed for cutting tools in a Al-matrix composites (Uday et al.,2009)and in steel matrix com-posites (Li et al.,2010).Among many requirements to cutting tools,cost is always at the top of the list.That is why this research was concentrated on the∗Corresponding author at:BAY-NANO (RI for Nanotechnology),Dep.Nano-composites and Uni.Miskolc,Egyetemvaros,E/7,3515Miskolc,Hungary.Tel.:+36304150002;fax:+3646362916.E-mail addresses:olga ver79@mail.ru (O.Verezub),kalazi@bzaka.hu (Z.Kálazi),kubaisy@mail.ru (A.Sytcheva),femkuzsy@uni-miskolc.hu(L.Kuzsella),buza@bzaka.hu (G.Buza),nikverezub@mail.ru (N.V.Verezub),FedorovA@.au (A.Fedorov),kaptay@ (G.Kaptay).cheapest possible steel matrix.To keep the cost of the cutting tool low,only the surface layer of the cutting tool will be improved,i.e.steel -matrix surface composites will be considered in this paper.It has been established by Iglesias et al.(2007)that the wear resistance of the composite increases with decreasing the size of the hard,reinforcing particles.That is why a special technology has been developed by Verezub et al.(2009)to ensure that the reinforc-ing particles are as small as ing smaller hard particles in the steel matrix of the cutting tool is expected to improve the performance of the cutting tool.There is a wide variety of reinforcing particles for the steel matrix.The most ‘popular’particles are TiC particles (Ala-Kleme et al.,2007)and WC particles embedded in the matrix of M2high-speed tool steels (Riabkina-Fishman et al.,2001)or in Ni–Cr matrix (St-Georges,2007).These particles are hard and thermodynami-cally stable.While TiC has superiority in both hardness (Kalpakjian,1995)and thermodynamic stability (Barin,1993)over WC,both of these particles are metallic in nature,and thus are well wettable by the steel matrix ensuring strong adhesion between the matrix and the reinforcing particles.This is essential to prevent reinforc-ing particles to turn out of the matrix without being worn,what is known to be one of the wear mechanisms of MMCs (Kaptay et al.,1997).In this respect WC is superior to TiC,as WC is perfectly wet-ted,while TiC is only moderately wetted by liquid steel as reviewed0924-0136/$–see front matter ©2010Elsevier B.V.All rights reserved.doi:10.1016/j.jmatprotec.2010.12.009O.Verezub et al./Journal of Materials Processing Technology211(2011)750–758751Table1Composition(in wt%)of the HVG steel(similar to105WCr6steel)(balance:Fe).C Mn Si S P Cu Cr Ni Al Ti Mo Nb W V O0.99 1.000.380.0140.0210.14 1.030.190.0280.0050.10.01 1.340.030.0056by Eustathopoulos et al.(1999)and shown theoretically by Kaptay (2005).Among the many possible technologies to produce steel matrix surface nano-composites,the laser melt injection(LMI)technology has been selected by the authors.This technology was developed 3decades ago by Ayers and Tucker(1980)to produce a surface composite layer.In this technique large(around100␮m)carbide particles are blown by a gas stream into a moving laser melted pool of a substrate metal.The method is superior to all coating technologies in providing perfect adhesion between the compos-ite and the substrate and also in providing large thickness(around 1mm)allowing to re-ground the cutting ter the LMI tech-nology has been proven to be efficient to produce WC particles reinforced steel matrix composites by Liu et al.(2008),using par-ticularly X40CrMoV5–1steel surface layer by Dobrzanski et al. (2005)and duplex stainless steels matrix by Do Nascimento et al. (2008).The combination of WC+Co particles was used by Bitay and Roósz(2006).TiC particles were added into liquid steel by Fábián et al.(2003).The drawback of LMI technology is that only large carbide par-ticles with a sufficiently large kinetic energy can break the high surface tension liquid metal/gas interface,as was proven for liq-uid steel by Farias and Irons(1985)and for liquid aluminum by Vreeling et al.(2000).This is especially true for low-density TiC particles(compared to the density of liquid steel)that are not‘per-fectly’wetted by liquid steel as shown by Verezub et al.(2005). In fact,incorporation problems for the TiC/liquid steel couple was mentioned in the veryfirst paper by Ayers and Tucker(1980)and was later confirmed by Králik et al.(2003).Thus,the steel rein-forcing matrix,the TiC particles,the LMI technique and the desire to produce surface nano-composite seem to be contradictory.To solve this problem,an in situ LMI technology was developed by Verezub et al.(2009)to produce steel matrix carbide reinforced surface nano-composites.The in situ production route of steel-matrix TiC reinforced composites has been known since the work by Terry and Chinyamakobvu(1991).This method has been developed further by using reaction casting by Feng et al.(2005)and high-energy electron beam irradiation by Lee et al.(2006).The method was extended to produce Fe/(TiW)C composite powder by Correa et al. (2007).Good tribological behaviour of TiC–ferrous matrix com-posites was shown by Kattamis and Suganuma(1990).The Fe/TiC composites were found to have excellent wear properties by Galgali et al.(1999),confirmed also for elevated temperatures by Degnan et al.(2001).The samefinding was extended by Dogan et al.(2001) for cast chromium steels reinforced by TiC particles.Nevertheless, the in situ production of Fe/TiC composites and the LMI technology was combined for thefirst time by Verezub et al.(2009).The goal of the present paper is to evaluate the machinability(upon producing a cutting tool from it)and also the performance as a cutting tool of a steel matrix(TiW)C reinforced surface nano-composite produced on a cheap steel matrix by the in situ LMI process.2.Materials and methods2.1.MaterialsLow-alloyed tool steel plates of grade HVG(Russian GOST5950-73,1973being the analogue of steel105WCr6)have been selected as a base material for the current research.Detailed chemical com-position of the HVG substrate is given in Table1.The initial size of the substrates was8mm×8mm×4mm.Additionally,tungsten carbide and metallic titanium powders of chemical purity,both with a particle size of45–70␮m were used.These two powders were mixed at a1:1molar ratio.This molar ratio was chosen as the most stable carbides in the Ti–W–C system are the TiC and WC car-bides,and in this way the exchange reaction Ti+WC=W+TiC can be ensured between them.Of-course,the C-content of the original steel will also play some role(see below).2.2.Production of the nano-compositesSchematic diagram of LMI-equipment used in the current study is shown in Fig.1.The upper8mm×8mm plane of the HVG sub-strate was coated by a thin layer of graphite to increase laser beam absorption efficiency during the LMI process.The other side of the steel substrate was brazed onto a large,water-cooled Cu-plate to ensure fast cooling of the substrate.The top surface of the substrate was melted by a2.5kW CO2continuous wave laser(type Trumph TLC105),with a laser spot of2mm in diameter.The laser spot was moving along the sample with a scanning speed of400mm/min. The(WC+Ti)powder mixture was blown into the melted pool at an angle of45◦using argon as carrier gas.The following three pow-der feeding rates were used during our experiments:1.3g/min, 2.3g/min and3.8g/min.Several laser tracks were drawn parallel to each other with a50%overlapping.After the LMI process,the rapidly cooled samples were heat treated under the following conditions:austenitizing at a tempera-ture of1000–1050◦C during10–15s in a high frequency induction furnace,followed by rapid cooling and tempering at a temperature of350◦C for1h.The second round of tempering was performed during1h at560◦C.For reasons of more correct comparison,the initial HVG samples were heat treated under usual conditions (hardening at840◦C and tempering at170◦C).The samples were grinded,polished,etched and analyzed uti-lizing special techniques.The microstructure of the substrates was observed using an AMRAY1810i SEM(Scanning Electron Microscopy with micro resolution),equipped with EDS(Energy Dispersive X-ray Spectroscopy).The identification of nano-sized particles was performed by a high resolution SEM(HitachiS-4800,Fig.1.The schematic diagram of the laser melt injection(LMI)equipment(1–laser, 2–powder nozzle,3–steel substrate to be treated,4–copper cooling plate,5–working table,6–cooling water input and output).752O.Verezub et al./Journal of Materials Processing Technology211(2011)750–758Japan).Quantitative analysis of the samples was performed by ImageJ software.In different parts of the paper the following short sample names are used:i.“LMI”is the sample produced here by the LMI procedure includ-ing heat treatment.ii.“HVG”is the original HVG sample(see Table1)heat treated as described above.iii.“HSS”is a commercially available high speed steel sample with 6%W+5%Mo.2.3.Microhardness measurement of the nano-compositeThe microhardness profiles were measured using TUKON2100B equipment(Wilson Instr.)using load of500g and time of pressing of10s.The samples were polished and etched before the micro-hardness measurements.Microhardness was scanned along two lines:(i)perpendicular to the surface,as function of depth,and(ii) parallel along the surface,at the depth of0.40mm.2.4.Machinability of the nano-composite by cubic BN wheelA cubic boron nitride(CBN)grindingflaring cup wheel of type L010(125/100)–100%–B1–58(Russian standard)which cor-responds to grinding wheel B120C100vitrified bond(Stephenson and Agapiou,2006)was used to remove small quantities of the sur-face composite material to produce the required shape and surface quality for the cutting tool insert.Additionally,the grinding ratio of the CBN wheel was studied by removing the same thickness of 1mm from each substrate(HVG,LMI,HSS).The grinding ratio G,is defined as the ratio of worn mass of the grinding wheel(mg)to the mass of the removed material(g).The CBN wheel was studied by SEM+EDS after the grinding experiments.2.5.Performance of a cutting tool made of LMI nano-composite materialSteel C45(0.45%C+0.6%Mn+0.25%Si)was used as a workpiece for the cutting experiments.Machining of the steel C45workpiece was performed by the HVG,LMI and HSS cutting tools.All the exper-iments were run with the following cutting conditions:cutting speed V=20–60m/min,feed f=0.05–0.3mm/rev and depth of cut d=0.25–1.75mm.The cutting force components were measured by a piezoelectric dynamometer(Kistler).SEM and EDS analysis of the cutting tool and the removed chips were applied after the cutting experiments.3.Results and discussion3.1.Structure and composition of the nano-compositeFig.2shows SEM pictures of cross section of the characteristic LMI sample.Fig.2a shows that the depth of the surface composite layer is approximately1mm.Due to multiple scanning by the laser beam,the depth of the melted layer shows a certain pattern in Fig.2a,with minima in the depth separated by a distance of about 1mm(what is half of the2mm laser spot diameter due to50% of overlapping).As one can see from Fig.2,the microstructure of the melted layer seems to be macroscopically homogeneous.This is due to the high velocity of Marangoni convection of the laser melted pool during the LMI process.Fig.3shows enlarged SEM pictures of the LMI sample with two types of precipitates.The several micron sized(Ti,W)C carbide pre-cipitates(Fig.3a)formed during fast cooling at the latest stage ofthe Fig.2.SEM pictures of the cross sections of the steel substrate after the LMI treat-ment(a)and the general view of the microstructure within the laser treated zone (b)(powder feeding rate is2.3g/min).LMI process by in situ precipitation from the molten steel matrix. The core of these precipitates is rich in Ti,while the outer region of the precipitates is rich in W.This is so due to higher thermodynamic stability of TiC compared to WC.The second type of precipitates is below100nm in diameter and is formed only during the subse-quent heat treatment.These nano-particles are(Fe,W)6C carbides (Fig.3a and b),precipitating from the supersaturated solid steel matrix(for more details see Verezub et al.,2009).The volume%of micron sized(Ti,W)C particles are shown in Fig.4as function of the powder feeding rate.The theoretical maximum,shown in Fig.4was calculated from the technologi-cal parameters and from the cross section of the melted zone(see Fig.2a).One can see that in the as-received LMI samples the amount of incorporated(Ti,W)C particles is somewhat lower compared to the theoretical maximum.The incorporation ratio decreases from 89%(for1.3g/min)to76%(for3.8g/m)with increasing the pow-der feeding rate.It is probably due to the gradual increase in the effective viscosity of the suspension with increasing its solid con-tent,what makes further incorporation and dissolution of(Ti+WC) particles more difficult.During heat treatment of the LMI samples, the volume%of micron-sized(Ti,W)C particles is decreased further by about20%.This is due to the partial dissolution of the W-rich outer region of the micron-sized(Ti,W)C precipitates.The amount of nano-sized(Fe,W)6C particles is found around25±5vol%,being independent of the powder feeding rate.These nano-sized particles form during the heat treatment,from the over-saturated matrix and partially from the dissolved outer regions of the micron-sized precipitates.As follows from materials balance,the majority of the content of these(Fe,W)6C nano-particles originate from the mate-rial of the matrix.Further investigation is needed to clarify howO.Verezub et al./Journal of Materials Processing Technology 211(2011)750–758753Fig.3.SEM micrographs of the cross section of the LMI sample in two different magnifications (powder feeding rate is 2.3g/min).the conditions of heat treatment influence the micro-and nano-structure of the composite and the amount and size distribution of (Fe,W)6C particles.It should be mentioned that at the highest powder feeding rate of 3.8g/min the LMI samples appeared to be cracked.This is probably due to the too high volume %of the carbide phase in the matrix.The two other samples (produced at the powder feeding rates of 1.3and 2.3g/min)are free of cracks.The latter is more promising as the higher amount of carbide phase leads to improved mechanical properties of the composite,if the formation of cracks is avoided.3.2.Microhardness of the LMI nano-composite sampleThe depth profile of microhardness of the LMI nano-composite sample with powder feeding rate of 2.3g/min is shown in Fig.5.All measurements are made after the heat treatment described in the experimental part.The depth profile can be divided into three regions:010*******12345powder feeding rate, g/min(T i ,W )C , v o l %Fig.4.The volume %of the micron-sized (Ti,W)C particles as function of powder feeding rate after the LMI process (before and after the heat treatment procedure).020040060080010001200140000.51 1.52M i c r o h a r d n e s s , HVDistance from surface, mmFig.5.Depth profile of microhardness of the LMI nano-composites (powder feeding rate is 2.3g/min).i.The upper surface layer of about 500␮m thickness has a highest microhardness of about 1200HV.ii.The initial substrate (below 1mm from the top surface)has a lowest microhardness of about 1000HV.iii.There is a transition zone between 500and 1000␮m measuredfrom the top surface,within which the microhardness gradually changes between the above mentioned limits.In evaluation of these results let us remind that carbon can diffuse from the non-melted part of the substrate into the melted LMI part of the substrate during the heat treatment.The increased microhardness of the upper surface layer of LMI nano-composite sample is obviously due to the precipitated micron-sized (Ti,W)C and nano-sized (Fe,W)6C hard carbide par-ticles.The existence of the intermediate zone could be due to the interplay between solidification rate (solidification goes from the bottom of the melted zone upwards)and the feeding and mixing rates of the added powder mixture (powder mixture is added to the top and the incorporated particles together with the dissolved atoms move downwards mainly by the Marangoni convection).In Fig.6a the measured microhardness is shown parallel along the sample surface,at the depth of about 0.4mm for both the as received LMI sample and the heat treated LMI sample.One can see that the microhardness of the LMI samples increase due to the heat treatment,what is probably due to the formation of (Fe,W)6C02004006008001000120014000.511.522.53M i c r o h a r d n e s s , H VDistance, mm00.20.40.60.8100.511.52 2.53d e p t h , m mdistance, mmFig.6.Microhardness scanned parallel along the surface,at the depth of about 0.4mm for both as received LMI sample and the heat treated LMI sample (a)and the depth of the melted pool as function of the same path (b)(see the pattern in Fig.2a).754O.Verezub et al./Journal of Materials Processing Technology 211(2011)750–7580.0340f , m m /p a s sv f ,/m i nFig.7.The grinding ratio of the LMI nano-composite sample (powder feeding rate is 2.3g/min,V =25m/s).nanoparticles.It is also obvious that the heat treatment flatters out the large fluctuations in the microhardness of the as received LMI sample.The minima in the microhardness fluctuations (Fig.6a)approximately coincide with the minima in the depth of the melted zone (see Fig.6b and the pattern in Fig.2a).This can be explained by Fig.5,measured at the largest depth of the melted zone.The smaller is the depth of the melted zone,the higher becomes the relative depth of the same absolute depth of 0.4mm,and thus,in accordance with Fig.5,the smaller is the microhardness.As follows from Figs.5–6,the microhardness of the produced nano-composite layer is around 12GPa.For this value the optimum grinding wheel and the optimum workpiece to be machined should be selected such that the ratio of microhardnesses of the machining and that of the to be machined materials should be at least 3.As a machining tool,CBN (cubic boron-nitride)has been selected with its microhardness of about 50GPa (Kalpakjian,1995)being about 4.2times stronger compared to the hardness of our LMI sample.On the other hand,the C45workpiece has been selected with its microhardness of about 2.7GPa,being about 4.4times less strong compared to our LMI sample.Thus,the microhardness of our LMI nano-composite sample is positioned almost in the middle (in a logarithmic scale)of the interval between the microhardness val-ues of the machining CBN tool and that of the to be machined C45workpiece.3.3.Machinability of the LMI nano-composite by cubic BN wheel During the LMI treatment the surface of the substrate melts,and thus it becomes quite uneven after solidification (the subsequent heat treatment does not provide any significant improvement).As a result,the as-received LMI nano-composite sample cannot be used as a cutting tool.Therefore,the as-received LMI nano-composite sample was grinded by a CBN wheel to obtain the shape and surface quality required for cutting tools.Fig.7shows the grinding ratio of the LMI nano-composite sam-ple as function of the feed rate of the workpiece (v f ,m/min)and a feed (f ,mm/pass).The combination of a feed of f =0.04m/pass and a feed rate of v f =3m/min leads to a maximum grinding ratio of about G =45–50mg/g.Based on the results shown in Fig.7,the optimal grinding conditions are selected as:f =0.01–0.02mm/pass,v f =1–2m/min and V =25m/s.Under these conditions the grinding ratio can be kept at a reasonable level of G =8–15mg/g.In compari-son,under the same conditions the grinding ratio for the HVG steel was found to be 6.8mg/g,while the grinding ratio for HSS is known to be about 5–6mg/g (Lisanov,1978).The increased grinding ratio of our LMI sample is obviously due to the hard (Ti,W)C and (Fe,W)6C particles in the surface of newly developed material.Fig.8.The EDS spectra of CBN wheels after grinding HVG (a)and LMI (b)samples (powder feeding rate is 2.3g/min).In Fig.8the energy dispersive X-ray spectra of two CBN wheels are compared after identical grinding runs of the HVG and LMI samples.In addition to the C-and Fe-peaks after grinding the HVG sample,large W and Ti peaks are observed after grinding the LMI nano-composite sample.This can be explained by stabilisation of the C-content of the initial steel substrate by added Ti and by the attraction between (Ti,W)and (B,N)atoms,respectively,being due to the existence of stable titanium boride,titanium nitride and tungsten boride compounds as reported by Barin (1993).Thus,dur-ing the grinding process part of the Ti-and W-content of the LMI nano-composite substrate adheres to the CBN surface.The pres-ence of solid titanium and tungsten carbides causes loosening of the CBN grains and their fallout,leading to intensive wear of the wheel,also shown by Klimenko et al.(1996).Forming the rake and flank surfaces during grinding of the LMI nano-composite samples resembles the grinding of high-speed steel cutting tools as shown by Mamalis et al.(2002).When the high quality alloyed layer is achieved,cutting edge without visible chip-ping is obtained.At the same time the edge roughness,as well as the radius of the cutting edge are higher for the HVG substrate com-pared to the LMI nano-composite substrate (Table 2).Increase of the feed rate and that of the feed lead to further increase in roughness of the tool’s cutting edge.Table 2Roughness of tool’s cutting edge and surfaces after grinding by CBN wheels (param-eters:V =25m/s,v f =2m/min,f =0.01mm/pass).Tool material Cutting edgeroughness R a ,␮m Roughness of rake and flank surfaces R a ,␮m HSS 1.2–1.30.15–0.18LMI 1.3–1.50.17–0.20HVG1.4–1.60.21–0.24O.Verezub et al./Journal of Materials Processing Technology 211(2011)750–75875500,10,20,30,40,5050100150200machining time, minV B , m mFig.9.The influence of machining time on flank wear for different cutting tool materials (V =25m/min,f =0.1mm/rev,d =0.5mm).Curves correspond to the HVG steel,LMI nano-composite produced with different powder feeding rates (figures on curves correspond to the unit of g/min),and HSS.Overall it can be concluded that CBN wheels can be used with optimum grinding parameters of f =0.01–0.02mm/pass,v f =1–2m/min and V =25m/s to convert the as-received LMI nano-composite into the cutting tool.The required shape and roughness of the cutting tool can be obtained with a reasonable grinding ratio of about G =8–15mg/g.3.4.Tool life of the cutting tool made of our nano-composite During testing of a new LMI nano-composite cutting tool on C45workpiece,crater wear was found to be negligible compared to flank wear.These two types of wear are the most common measured forms of tool wear.Thus,the tool life of this newLMI2040608010012014016018001234powder feeding rate, g/mint o o l l i f e , m i nFig.10.Tool life as function of the powder feeding rate during the LMI process (V =25m/min,f =0.1mm/rev,d =0.5mm)(the point at zero feeding rate refers to a different heat history of a sample,that is why this point is connected to other points by a thin line).nano-composite cutting tool is determined from the measured flank wear.Fig.9shows flank wear measurements for HVG steel used as a base material,LMI nano-composite produced with different pow-der feeding rates and HSS.The critical flank wear of 0.45mm was chosen based on values recommended for replacing or re-grounding alloyed tool materials (Kalpakjian,1995).The machining time during which the actual flank wear achieves the critical value is called tool life (T ,min).Tool life as function of the powder feeding rate is shown in Fig.10.It shows that an optimum value of the powder feeding rate exists for the maximum tool life.When Fig.10is rationalized in combination with Fig.4,it can be seenthat the volume %of carbide particles in the compositeFig.11.SEM images of the cutting tool made of the LMI sample after its service (a–c)and EDS spectrum (d)of the worn surface.756O.Verezub et al./Journal of Materials Processing Technology211(2011)750–758Fig.12.Removed chip from steel C45by the LMI cutting tool(a)and its EDS spec-trum(b).gradually increases with the increase of powder feeding rate and,as a consequence,tool life also increases.However,as was mentioned above,cracks were formed in the substrate,made by the powder feeding rate of3.8g/min.As a result,a cutting tool made of this substrate has a lower tool life.One can suppose that there is an optimum feeding rate in the interval between2.3and3.8g/min, when the volume%of carbide particles is somewhat larger than for the2.3g/min feeding rate,but still without crack formation.The SEM images of the LMI nano-composite cutting tool faces after machining of C45steel are shown in Fig.11.Fig.11a shows the overlapping of the LMI tracks and the traces of theflank wear(mea-sured as0.45mm).In Fig.11b–c carbide particles being similar to those shown in Figs.2–3are shown.The difference is that the steel matrix is worn away in between the hard carbide particles after machining compared to the initial state of the LMI nano-composite substrate(compare Fig.11b–c to Figs.2–3).Therefore,it is evident that theflank wear is the result of abrasive wear of the LMI cutting tool.The EDS spectrum(Fig.11d)of the worn surface shows Fe,Ti, W as basic components.In Fig.12the SEM picture and EDS spectrum of the removed chip from the C45workpiece is shown,after its machining by the LMI nano-composite cutting tool.It can be seen that the removed chip is continuous,and the main elements of the nano-composite (Ti and W)are missing from its X-ray spectrum.Thus,there was no adhesion of Ti and/or W to the C45steel workpiece during its machining by the LMI nano-composite cutting tool.In order to position our cutting tool made of the LMI substrate on a tool-life scale,tool-life tests have been conducted.The effect of cutting speed V on tool-life T has been assessed using Taylor’s tool life equation(Eq.(1))(Taylor,1907)and the results are shown Table3Experimental tool life(T,min)of different cutting tool materials against C45work-piece(f=0.1mm/rev and d=0.5mm).V,m/min T,min(experimental)HVG LMI HSS20385500535 2550160190 30157088 3542550 40–1726 45–715 50–47 60–24in Table3.C1=V·T n(1) Eq.(1)is widely used and recognized in the industry.It relates tool life to the cutting speed through empirical tool life constants n and C1.Table4shows the range of values n and C1for different cutting tool materials obtained from the data in Table3.The data(Table4)indicate that the LMI process of inserting (TiW)C particles into HVG substrate improved tool life of the base material by300–400%when cutting speed V was25–35m/min. However,this new material was felt short to surpass tool life of HSS cutting material by just20%in the same cutting speed range.No sig-nificant difference between tool life of HSS and LMI was observed during machining at20m/min.Performance of cutting tools made of LMI nano-composite is similar to the performance of HSS and is limited by wear resistance at cutting speeds above40–45m/min (Stephenson and Agapiou,2006).Cutting tools made of HVG steel can be used at speeds up to30m/min.3.5.Cutting force componentsDuring machining of the C45workpieces by the cutting tool made of HVG and LMI substrates,the two main force components F z (N)and F x(N)have been measured as function of the depth of cut d (mm)and feed f(mm/rev).The effects of depth of cut and feed on the measured F z and F x force components for the two different cutting materials(HVG and LMI nano-composite)are shown in Figs.13–14. The cutting speed increase within limits of V c=20–60m/min does not sufficiently influence the value of the cutting force(Fedorov, 2005)and therefore has not been tested in this paper.The effect of depth of cut d on the measured force components for two different cutting tool materials is shown in Fig.13a and b. Forces F z and F x increase with the increase in depth of cut because the increase in depth of cut leads to increase in the area of cut and length of the cutting edge in contact.The influence of feed f on the forces F z and F x is shown in Fig.14a and b.The increases in feed lead to increase in cut thickness,which,in turn,increases the area of cut and as a consequence,the force components.Figs.13–14show that machining with LMI nano-composite cut-ting tool material decreases cutting forces F z and F x in comparison with HVG cutting tool material.However,significant force reduc-tions can only be observed when depth of cut is greater than1mm and feed is greater than0.2mm/rev.The force components can be described as function of parameters d and f by the followingTable4Values of n and C1for different tool materials(f=0.1mm/rev,d=0.5mm).Cutting tool material n C1HSS0.2280.89 LMI0.1967.18 HVG0.1241.58。

Gd含量对Al-Zn-Mg-Cu-Zr合金微观组织与力学性能的影响梅飞强;王少华;房灿峰;孟令刚;贾非;郝海;张兴国【摘要】采用铸锭冶金工艺制备6种Gd含量不同的A1-Zn-Mg-Cu-Zr-xGd合金.采用金相观察、力学性能测试、扫描电镜、电子探针及透射电镜等分析手段,研究质量分数x,分别为0%、0.10%、0.15%、0.20%、0.25%和0.30%的Gd对基体合金铸态及时效态显微组织和力学性能的影响.结果表明：Gd含量对A1-Zn-Mg-Cu-Zr合金的微观组织和力学性能的影响显著,当Gd含量低于0.25%时,随Gd含量的增加细化效果、强度以及伸长率都增加；当Gd含量为0.25%时,铸态组织中基体晶粒最小,达到32μm左右；此时T6态合金组织的强度和伸长率达到最高,抗拉强为624.54 MPa,屈服强度为595.00 MPa,伸长率为13.3%,且固溶组织具有良好的抗再结晶作用；而当Gd含量超过0.25%时,合金的微观的组织与力学性能变差.【期刊名称】《中国有色金属学报》【年(卷),期】2012(000)009【总页数】9页(P2439-2447)【关键词】A1-Zn-Mg-Cu-Zr合金;Gd;显微组织;力学性能【作者】梅飞强;王少华;房灿峰;孟令刚;贾非;郝海;张兴国【作者单位】大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024;大连理工大学材料科学与工程学院,大连 116024【正文语种】中文【中图分类】TG146.2+1超高强铝合金是20世纪60年代以航空航天材料为背景发展起来的一种高强度铝合金材料[1−3]。

一般将屈服强度为500 MPa以上的铝合金称为超高强铝合金[4]。

128精密成形工程 2023年7月[9] HUANG Meng, XU Chao, FAN Guo-hua, et al. Role ofLayered Structure in Ductility Improvement of LayeredTi-Al Metal Composite[J]. Acta Materialia, 2018, 153: 235-249.[10] 王锟. 基于内聚力-GTN混合模型的钛-铝层状复合板损伤研究[D]. 洛阳: 河南科技大学, 2018: 43-46.WANG Kun. Study on Damage of Ti-Al Laminated Composite Plate Based on CZM-GTN Hybrid Model[D].Luoyang: Henan University of Science and Technology,2018: 43-46.[11] 郭照灿. 铜/铝双金属复层材料应变光学检测及力学性能研究[D]. 郑州: 郑州轻工业大学, 2022: 31-42.GUO Zhao-can. Optical Strain Detection and Mechani-cal Properties of Cu/Al Bimetallic Clad Materials[D].Zhengzhou: Zhengzhou University of Light Industry, 2022: 31-42.[12] 李莎, 贾燚, 刘欣阳, 等. 层状镁/铝复合板轧制工艺研究进展[J]. 精密成形工程, 2021, 13(6): 1-11.LI Sha, JIA Yi, LIU Xin-yang, et al. Research Progresson Rolling Process of Laminated Mg/Al Clad Plate[J].Journal of Netshape Forming Engineering, 2021, 13(6):1-11.[13] 高勃兴, 邹德坤, 谢红飙, 等. 铝/钢轧制复合有限元二次开发模拟与实验研究[J]. 精密成形工程, 2021, 13(6): 56-63.GAO Bo-xing, ZOU De-kun, XIE Hong-biao, et al.Simulation and Experimental Study on Finite Element Secondary Development of Aluminum/Steel Rolling Composite[J]. Journal of Netshape Forming Engineering, 2021, 13(6): 56-63.[14] LIU H S, ZHANG B, ZHANG G P. Enhanced Tough-ness and Fatigue Strength of Cold Roll Bonded Cu/Cu Laminated Composites with Mechanical Contrast[J].Scripta Materialia, 2011, 65(10): 891-894.[15] HUANG M, CHEN J S, WU H, et al. Strengthening andToughening of Layered Ti-Al Metal Composites by Controlling Local Strain Contribution[J]. IOP Confer-ence Series: Materials Science and Engineering, 2017, 219: 012028.[16] LI Z, LIN Y C, ZHANG L, et al. In-situ Investigation onTensile Properties of a Novel Ti/Al Composite Sheet[J].International Journal of Mechanical Sciences, 2022, 231: 107592.[17] HUANG C X, WANG Y F, MA X L, et al. InterfaceAffected Zone for Optimal Strength and Ductility inHeterogeneous Laminate[J]. Materials Today, 2018, 21(7): 713-719.[18] CHEN Wen-huan, HE Wei-jun, CHEN Ze-jun, et al.Extraordinary Room Temperature Tensile Ductility ofLaminated Ti/Al Composite: Roles of Anisotropy andStrain Rate Sensitivity[J]. International Journal of Plas-ticity, 2020, 133: 102806.[19] XING Bing-hui, HUANG Tao, XU Liu-jie, et al. Effectof Heat Treatment Process on the Microstructure of theInterface of Ti/Al Laminated Composite[J]. CompositeInterfaces, 2022, 29(7): 749-764.[20] 金属材料拉伸试验第1部分: 室温试验方法: GB/T228. 1—2010[S]. 2011.Metallic Materials-Tensile Testing-Part 1: Method ofTest at Room Temperature: GB/T 228.1—2010[S]. 2011.[21] 杨方方, 皇涛, 陈拂晓, 等. 异质金属层状复合板分层应力-应变关系研究[J]. 塑性工程学报, 2018, 25(1):187-191.YANG Fang-fang, HUANG Tao, CHEN Fu-xiao, et al.Research on Delamination Stress-Strain Relationship ofHeterostructure Laminated Composite Sheets[J]. Journalof Plasticity Engineering, 2018, 25(1): 187-191.[22] HUANG Tao, PEI Yan-bo, CHEN Fu-xiao, et al. ANovel Layered Finite Element Model for Predicting theDamage Behavior of Metal Laminated Composite[J].Composite Structures, 2023, 311: 116786.[23] TVERGAARD V, HUTCHINSON J W. The Relationbetween Crack Growth Resistance and Fracture ProcessParameters in Elastic-Plastic Solids[J]. Journal of theMechanics and Physics of Solids, 1992, 40(6): 1377-1397.[24] GURSON A L. Continuum Theory of Ductile Ruptureby Void Nucleation and Growth: Part I-Yield Criteriaand Flow Rules for Porous Ductile Media[J]. Journal ofEngineering Materials and Technology, 1977, 99(1): 2-15.[25] TVERGAARD V. Influence of Voids on Shear BandInstabilities under Plane Strain Conditions[J]. Interna-tional Journal of Fracture, 1981, 17(4): 389-407.[26] HUANG M, FAN G H, GENG L, et al. Revealing Ex-traordinary Tensile Plasticity in Layered Ti-Al MetalComposite[J]. Scientific Reports, 2016, 6(1): 1-10.责任编辑：蒋红晨第15卷 第7期 精 密 成 形 工 程2023年7月JOURNAL OF NETSHAPE FORMING ENGINEERING129收稿日期：2023‒03‒22 Received ：2023-03-22基金项目：国家自然科学基金（52001188）Fund ：National Natural Science Foundation of China (52001188)作者简介：孙捷（1988—），男，博士，讲师，主要研究方向为镁合金塑性变形机理Biography ：SUN Jie (1988-), Male, Doctor, Research focus: plastic deformation mechanism of magnesium alloy. 引文格式：孙捷, 曲京儒, 阎玉芹, 等. 稀土元素对轧制Mg–Zn–Zr 合金板材微观组织和力学性能的影响[J]. 精密成形工程, 2023, 15(7): 129-135.SUN Jie, QU Jing-ru, YAN Yu-qin, et al. Effect of Rare Earth (RE) on Microstructure and Mechanical Property of Rolled 稀土元素对轧制Mg–Zn–Zr 合金板材微观组织和力学性能的影响孙捷，曲京儒，阎玉芹，赵彦华（山东建筑大学 机电工程学院，济南 250101）摘要：目的 为了使Mg–Zn–Zr 合金在热加工过后具有良好的力学性能及变形各向同性，在Mg–2Zn–0.5Zr 合金中添加不同含量的稀土元素，研究稀土元素对Mg–2Zn–0.5Zr 合金轧制后微观组织和力学性能的影响规律，以解决变形镁合金织构强、变形各向异性强的问题。