Fracture behavior of Metallic multilayers

INTERFACE SCIENCEEffect of interface structures on the fracture behavior of two-dimensional carbon/carbon composites by isothermal chemical vapor infiltrationYong-Gang He •Ke-Zhi Li •He-Jun Li •Jian-Feng Wei •Qian-Gang Fu •Dong-Sheng ZhangReceived:22July 2009/Accepted:25November 2009/Published online:10December 2009ÓSpringer Science+Business Media,LLC 2009Abstract The interface structures and fracture behavior of the two-dimensional carbon/carbon composites by isother-mal vapor infiltration have been investigated.The results show that the graphene layers exhibit long-range order in high/textured pyrocarbon matrix and are curved in about 5-nm interface region of the fiber/high-textured.Some globular nanoparticles are formed on the fiber surface and the high-textured layer about 10nm exists in the interface of the fiber/low-textured.The graphene layers stacks are scrolled and folded in the medium-textured and they are waved together in the interface of the fiber/medium-tex-tured.The pseudo-plastic fracture behavior of the two-dimensional carbon/carbon composites is resulted from the dominant high-textured matrix and a moderate interfacial bonding force.A strong adhesion of the fiber/low-textured and the thicker fiber increased by surrounding low-textured layer result in the increasing flexural strength.The single medium-textured and a very strong bonding force of the fiber/medium-textured lead to the brittle fracture behavior.IntroductionCarbon/carbon (C/C)composites have attracted extensive attentions due to their excellent mechanical properties at ultrahigh temperature,such as high specific strength,specific modulus,and steady strength above 2473K [1,2].Therefore,they are suitable for aircraft and aerospace applications.The C/C composites are often fabricated by isothermal chemical vapor infiltration (ICVI)of porous carbon fiber preforms.Benzinger and coworkers [3–5]have investigated the correlation between ICVI process parameters and pyrocarbon matrix microstructures.Mean-while,several models have been proposed to correlate the degree of texture in a qualitative way with the composition of the gas phase.The textures of the bulk pyrocarbon matrix could be controlled well by choosing appropriate ICVI process parameters.However,the texture of the very first layer of pyrocarbon about a few nanometers thick is considerably difficult to control.Furthermore,the structure and characteristics of the interface determine the adhesion between fiber and matrix.The mechanical properties of C/C composites tremendously depend on load transfer at fiber/matrix (F/M)interface.A weak interface may impair the integrity of composites,whereas a strong bond may induce brittle fracture behavior [6–8].Therefore,study of interface characteristics is essential for the development of C/C composites with desirable and predictable behavior.The two-dimensional carbon/carbon (2D C/C)compos-ites are widely used as brake disks for airplane in aero-nautic fields.It is well known that the microstructure of pyrocarbon matrix and the property of interface greatly affect their mechanical and tribological properties.A number of researchers have extensively studied the microstructure of pyrocarbon matrix and failure mecha-nism of C/C composites by ICVI [9–11].However,the earlier study mainly focused on the structure and mechanical properties of the carbon fiber felts by ICVI.Although a few studies have been reported on the influence of the matrix texture on the fracture behavior of 2D C/C composites [12,13],the structural properties of the F/MY.-G.He ÁK.-Z.Li (&)ÁH.-J.Li ÁJ.-F.Wei ÁQ.-G.Fu ÁD.-S.ZhangCarbon/Carbon Composites Technology Research Center,Northwestern Polytechnical University,Xi’an,Shaanxi 710072,People’s Republic of China e-mail:likezhi@ Y.-G.Hee-mail:heyg@J Mater Sci (2010)45:1432–1437DOI 10.1007/s10853-009-4089-0interface and their effect on the fracture behavior of2D C/C composites by ICVI have not appeared yet in the literature.In this study,the micro-and nanostructures of the F/M interface of2D C/C composites by ICVI were analyzed systematically by high-resolution transmission electron microscopy.In particular,the effect of interface micro-structures on the facture behavior of2D C/C composite was discussed.ExperimentalPreparation of the2D C/C compositesThe investigated samples are2D needle-punchedfiber felt with afiber volume fraction of25%and a special archi-tecture of0°/90°/0°/90°.The PAN(polyacrylonitrile)-based carbonfibers forming the felts have a typical mean diameter of7l m.Thefiber density is1.72g/cm3.All the preforms were heat-treated at1773–1973K before densi-fication to remove the impurity on thefiber surface.The infiltration was carried out by means of ICVI process using methane as precursor gas and nitrogen as the diluted gas at negative pressure.The experimental operation and infil-tration procedure were described in detail in Ref.[14].The infiltration conditions were adjusted to obtain bulk samples with different pyrocarbon matrix.Samples A and B were densified at1373K with the methane partial pressure of10 and20kPa,respectively.Sample C was infiltrated at 1323K with the methane partial pressure30kPa.The same gas residence time of0.1s was used for preparing the three samples.The bulk densities of the three samples were increased to1.75g/cm3by different densification time. Characterization methodsThe infiltrated2D C/C composites were heated in high-temperature graphitization furnace at2373K for2h at an argon atmosphere.Three rectangular bars of55mm9 10mm94mm along thefiber layer direction were cut from the inner region of the samples A–C,respectively. Three-point bending tests were carried out in accordance with Q/GB95-92on a CMT5304-30KN universal testing machine.The optical observation was performed on polarized light microscopy(PLM,Leica DMLP).The extinction angle(Ae)was measured according to the pro-cedure described by Bourrat et al.[15].Three kinds of the pyrocarbon can be distinguished:low-textured(LT, Ae\12°),medium-textured(MT,12°B Ae\18°),and high-textured(HT,Ae C18°).The morphologies of frac-ture surface and the textures of pyrocarbon matrix were examined by scanning electron microscopy(SEM,JSM-6700F).The microstructures of the F/M interface were analyzed by(high-resolution)transmission electron microscopy(TEM/HRTEM,JEM-3010)combining with selected area electron diffraction(SAED).The orientation angle(OA)is obtained by SAED patterns.According to Reznik and H}u ttinger[16],the pyrocarbon textures are classified into LT,80°B OA\180°,MT,50°B OA \80°,and HT,OA\50°.Results and discussionTextures and microstructures of the interfaceThe polarized light micrographs of the samples A-C are shown in Fig.1.The textures vary significantly with the changing of infiltration process conditions.The matrix of sample A infiltrated at1373K with the methane partial pressure of10kPa exhibits pure HT.The matrix of sample B densified at identical temperature with a higher methane partial pressure20kPa possesses double textures around thefiber.The inner dark layer about1l m thick is LT layer and the outer layer is HT.The carbonfibers are surrounded by single MT in sample C.The PLM provided the information about the texture on a coarse scale.For analyzing the microstructures close to the F/M interface in more detail,TEM and HRTEM with SAED were used in this study.Figure2shows the TEM micrographs of the F/M interface and the SAED patterns of the matrix in the as-received composites.In Fig.2a,the OA of23°±2°corresponds to HT pyrocarbon matrix and the inserted SAED pattern illustrates the high degree of parallel alignment of the(002)basal planes in HT matrix. Moreover,it can be seen that the HT layer contains numerous microcracks,which are approximately parallel to one another and oriented parallel to thefiber surface. Owing to the weak vander Waals forces between layers, microcracks are easily induced in HT matrix during the course of mechanically thinned and ion milling.However, no microcracks are displayed in the interface of thefiber–HT matrix.It demonstrates that there is a relatively good adhesion between thefiber and HT matrix.In Fig.2b,c, the inserted SAED patterns and the corresponding OA angles indicate two quite different textures,i.e.,LT matrix, OA=105°±2°and MT matrix,OA=65°±2°.In contrast to HT matrix,there are not any cracks within LT and MT matrix in the vicinity of thefiber surface.It can be seen that the LT and MT are composed of particles with different sizes and shapes.These particles are cross-linked to form mosaic-like texture.Furthermore,few microcracks can be detected in the interface of thefiber–LT matrix. Reznik and Gerthsen[17]also observed that the cracks are generated not directly at thefiber–LT matrix interface butat some distance in the LT matrix.These observations are indicative of a strong adhesion between the fiber and LT matrix.Figure 2c indicates that there is a microcrack between the fiber and MT matrix,which is produced by heat treatment.However,the microcrack does not spread throughout the fiber surface.Figure 3displays the HRTEM micrographs of the F/M interface region of 2D C/C composites.It should be noted that there are obviously different interface microstructuresbetween the fibers and HT,LT,and MT,respectively.Figure 3a reveals that the highly ordered graphene layers exhibit long-range order and parallel to the fiber surface in the HT matrix.The 002lattice fringes are curved to some extent in about 5-nm interface region of the fiber–HT matrix.These curved lattice fringes may be produced by stress or they replicate the unevenness of the fiber surface.Therefore,debonding cannot probably take place between the fiber and HT matrix in the course of mechanically thinned and ion milling.It is clear in Fig.3b that some globular particles about a few nm are first formed on the fiber surface.The following is the HT layer of up to 10nmFig.1Polarized light micrographs of a sample A;b sample B;c sampleCFig.2TEM micrographs of the fiber/matrix interface in the 2D C/C composites a sample A;b sample B;c sample C.Inserts correspond to SAED patterns of the pyrocarbon matrixand the next layer is LT matrix.The globular particles possibly nucleate directly at active sites on the fiber surface and grow up gradually in the following deposition.The polyhedral particles forming LT matrix exhibit oriented graphite-like shells and a relatively disordered central part.In regard to the forming of HT layer on a scale of a few nanometers,Gerthsen et al.[18]also observed similarhighly oriented layers in the interface of fiber–LT matrix and proposed stress-induced ordering mechanism.The detailed mechanism of the generation of such layer is not clear yet at present.Nevertheless,the ordered graphene layers of the shells of LT particles should be helpful to the formation of the thin HT layer.It can be concluded that a considerably strong adhesion between the fiber and LT matrix is achieved.The contact area microstructure of the interface region between the fiber and MT matrix is shown in Fig.3c.It can be seen that in the vicinity of the PAN carbon fiber surface contains extended domains of stacked graphene layers which are often oriented parallel to the fiber surface.Meanwhile,the graphene layers stacks are scrolled and folded in the MT matrix.Moreover,some graphene layers in fiber and MT matrix waved together in the bonding area.They can improve the bonding force of the fiber with MT matrix like a hook.Therefore,the interfacial bonding force of the fiber–MT matrix is very strong.Fracture behavior of the 2D C/C compositesThe typical nominal stress–strain curves of the investigated samples A–C are shown in Fig.4.It can be seen from the curves that samples A and B exhibit pronounced pseudo-plastic fracture behavior,while sample C experiences a brittle fracture behavior.Sample C possesses the highest flexural strength.Samples A and B have the lowest and medium flexural strength,respectively.SEM micrographs of fracture surface of the samples after three-point bending tests are demonstrated in Fig.5.A larger extent of fiber-pullout in samples A and B agrees with the pseudo-plastic fracture behavior.Less obvious fiber-pullout proves the brittle fracture behavior of sample C.In addition,Fig.6Fig.3HRTEM micrographs of the fiber/matrix interface a the fiber–HT interface in sample A;b the fiber–LT interface in sample B;c the fiber–MT interface in sampleCFig.4Typical nominal stress–strain curves of the 2D C/C compositesshows the SEM micrographs of fracture surface of the different pyrocarbon matrices in samples A–C.In Fig.6a,b,a stepwise morphology and relative rough fracture sur-face are observed in HT layers.The HT layer exhibits an intensive fragmentation of kinked sublayers.From Fig.6b,we can see that the fiber and its vicinal LT pyrocarbon layers suffer from fracture in the same plane.As a result,they are pulled out together from HT layers.Figure 6c demonstrates that the fracture surface of MT layers is a very smooth plane which is vertical to the fiber surface.The fracture behavior of 2D C/C composites is deter-mined by the matrix texture and the F/M interface in the composites.The experimental results show that the cracking occurs mainly within HT carbon layers and less frequently in MT and LT carbon layers.The rough fracture surface of HT matrix proves that a multiple crack deflec-tion takes place in HT layers leading to energy dissipation,which makes contribution to the toughness enhancement.In sample A with pure HT matrix,the crack propagation is blocked by the curved graphene layers on the fiber surface.The debonding generates not directly at the F/M interface but in the highly ordered graphene layers.The pseudo-plastic fracture behavior of sample A can be ascribed to the pure HT layer matrix and a comparatively moderate interfacial bonding force.The globular particles growing at active sites on the fiber surface and thin HTlayerFig.5SEM micrographs of fracture surface of the samples after three-point bending tests a sample A;b sample B;c sampleC Fig.6SEM micrographs of fracture surface of the different pyro-carbon matrices a HT;b HT and LT;c MToriginating from LT matrix make thefiber and LT matrix adhere closely.The microcracks from the outer layer HT matrix take place defection at the LT–HT interface.They propagate along the surface of LT matrix,and thus the sliding is brought out between the LT and HT matrix. Therefore,the toughness fracture behavior of sample B is caused by the dominant HT layer and the interfacial sliding between different texture matrices.Both a strong adhesion of thefiber–LT matrix and the thickerfiber increased by surrounding LT pyrocarbon result in the increasing of flexural strength.Owing to the mutual winding of graphene layers in the F/M interface and MT matrix,the defection and propagation of the cracks become very difficult. Consequently,sample C experiences a brittle fracture behavior and possesses the highestflexural strength. ConclusionsThe interface microstructures of thefiber–HT,–LT,and –MT matrices are obviously different.The highly ordered graphene layers exhibit parallel to thefiber surface in HT matrix.The microcracks are not produced in the interface of thefiber–HT matrix,but in HT matrix.Some globular particles on a scale a few nanometers exist at the interface of thefiber–LT matrix.An HT layer of about10nm is formed before LT deposition.The graphene layers infibers and MT matrix are waved together in the bonding area, which can improve the bonding force of thefiber–MT matrix.The pseudo-plastic fracture behavior of the2D C/C composites resulted from the dominant HT matrix and a comparatively moderate interfacial bonding force.A strong adhesion of thefiber–LT matrix and the thickerfiber increased by surrounding LT layers result in the increasing flexural strength.The single MT matrix and a very strong bonding force of thefiber–MT matrix lead to a brittle fracture behavior.Acknowledgments The authors thank thefinancial support from the National Natural Science Foundation of China(No.90716024)and the‘‘111’’Project(No.B08040).References1.Fitzer E,Manocha LM(1998)Carbon reinforcements and car-bon/carbon composites.Springer,Berlin2.Li HJ,Li AJ,Bai RC,Li KZ(2005)Carbon43:29373.Benzinger W,H}u ttinger KJ(1999)Carbon37:9314.Hu ZJ,Zhang WG,H}u ttinger KJ,Reznik B,Gerthsen D(2003)Carbon41:7495.Zhang WG,H}u ttinger KJ(2003)Carbon41:23256.Tang SF,Deng JY,Liu WC,Yang K(2006)Carbon44:28777.Li XT,Li KZ,Li HJ,Wei J,Wang C(2007)Carbon45:16628.Tezcan J,Ozcan S,Gurung B,Filip P(2008)J Mater Sci43:1612.doi:10.1007/s10853-007-2333-zvenac J,Langlais F,Fe´ron O,Naslain R(2001)Compos SciTechnol61:33910.Lee JY(2005)J Mater Sci40:3573.doi:10.1007/s10853-005-2884-911.Guellali M,Oberacker R,Hoffmann MJ(2005)Carbon43:195412.Xu GZ,Li HJ,Bai RC,Wei J,Zhai YQ(2008)Mater Sci Eng A478:31913.Chen TF,Reznik B,Gerthsen D,Zhang WG,H}u ttinger KJ(2005)Carbon43:308814.Li HL,Xu GZ,Li KZ,Wang C,Li W,Li ML(2009)J Mater SciTechnol25:10915.Bourrat X,Trouvat B,Limousin G,Vignoles G,Doux F(2000)JMater Res15:9216.Reznik B,H}u ttinger KJ(2002)Carbon40:62117.Reznik B,Gerthsen D(2003)Carbon41:5718.Gerthsen D,Bach D,Pauw VD,Kalho¨fer S,Reznik B,Send W(2006)Int J Mater Res97:1052。

Fracture Mechanisms in Bulk Metallic Glassy MaterialsZ.F.Zhang,1,2,*G.He,1J.Eckert,1and L.Schultz11IFW Dresden,Institute of Metallic Materials,01069Dresden,Germany2Shenyang National Laboratory for Materials Science,Institute of Metal Research,Chinese Academy of Sciences,110016Shenyang,People’s Republic of China(Received7March2003;published24July2003)W efind that the failure of bulk metallic glassy(BMG)materials follows three modes,i.e.,shearfracture with a fracture plane significantly deviating from45 to the loading direction,normal tensilefracture with a fracture plane perpendicular to the loading direction,or distensile fracture in a break orsplitting mode with a fracture plane parallel to the loading direction.The actually occurring type offailure strongly depends on the applied loading mode and the microstructure of the material.Extensiveevidence indicates that the Tresca fracture criterion is invalid,and for thefirst time,three fracturecriteria are developed for isotropic materials with high strength,such as advanced BMGs or the newlydeveloped bulk nanostructural materials.DOI:10.1103/PhysRevLett.91.045505P ACS numbers:62.20.Mk,46.50.+aBulk metallic glass(BMG)materials possess remark-able physical,chemical,and mechanical properties and have potential applications in many areas,which have created extensive interest among scientists and engineers [1,2].Normally,metallic glasses often deform by the formation of localized shear bands and display little plas-ticity before fracture at room temperature[3–6].How-ever,abundant observations[7–20]show that BMGs always fracture along a plane deviating from the maxi-mum shear stress plane,which is not in agreement with the Tresca criterion[19].On the other hand,recently developed BMG composites containing nanocrystals, hard particles,strongfibers,or ductile phases can exhibit high strength and good ductility[21–27].However, multifarious fracture behavior is frequently observed [24–27]and can be attributed to the special nature of the BMG materials.The present research will elucidate all the possible fracture phenomena of BMGs induced by uniaxial loading.Based on our experiments,we have developed comprehensive criteria to explain the various existing fracture mechanisms,which can be widely used for isotropic materials with high strength.The Tresca criterion[19]considers that fracture always occurs along the maximum shear stress plane at a critical shear stress 0.Therefore,all metallic glassy specimens should preferentially fracture along the maximum shear stress plane(45 )no matter whether under tension or compression.However,extensive investigations from the1970s to present show that the fracture angles between the stress axis and the shear plane of metallic glasses always deviate from45 ,as listed in Table I[7–20].Our recent investigations also reveal such a shear fracture behavior for different BMGs under tension and compres-sion.The typical shear fracture behavior of BMGs is illustrated in Figs.1(a)–1(c).For Zr59Cu20Al10Ni8Ti3, Zr54:5Cu20Al10Ni8Ti7:5,Zr55Cu30Al10Ni5,and Zr52:5Cu17:9Al10Ni14:6Ti5metallic glasses,the compres-sive fracture angle C between the compressive axis and the fracture plane is in the range of40 –43 . For some BMG composites containing Ta particles or dendritic precipitates[ Zr0:55Cu0:30Al0:10Ni0:05 95Ta5or Ti56Cu16:8Ni14:4Sn4:8Nb8],their C is only about32 or27 ,respectively,revealing an even more significant deviation from45 .Obviously,Ta particles or dendritic precipitates greatly change the shear fracture plane of the BMG composites.Moreover,we also found that the ten-sile fracture angles T of Zr59Cu20Al10Ni8Ti3and Zr52:5Cu17:9Al10Ni14:6Ti5metallic glasses are equal to 54 and56 ,respectively.This agrees well with previous findings,i.e., T approximately deviates5 –20 from45 for different BMG materials.Based on the data available (Table I)and the present results,it can be concluded that the shear fracture of metallic glasses does not obey the Tresca criterion[19],namely, T is larger than45 ,but C is smaller than45 :0< C<45 < T<90 :(1) Accordingly,there must exist a shear fracture mechanism different from the Tresca criterion.Since all the BMG materials possess very high strength(1.5–2.5GPa),the normal stress applied on the shear plane is also very high ( 1GPa)and must play an important role in the frac-ture process.Under tension or compression,the stress state( ; )of a specimen on any plane can be illus-trated as in Fig.1(d).If the normal stress( C or T)on the shear plane is taken into account,similar to the Mohr-Coulomb model[19],we can express the critical fracture stress( C or T)for both shear compressive and tensile failure asC 0 C C for compression ;(2a)T 0ÿ T T for tension :(2b) Here C and T are two constants for compression and tension and will be discussed later.In a ÿ coordinate,Eqs.(2a)and(2b)can be plotted in Fig.1(e),where T and C are equal to the slopes of the critical tensile and compressive fracture lines AB and AE,respectively.Meanwhile,another constant 0can be regarded as the critical normal fracture stress under the condition without shear stress and T 0= 0.Under tension or compression,the stress state ; on any shear plane will distribute on the Mohr circles a or b in Fig.1(e).When the Mohr circles touch the critical fracture lines AB or AE,the tangency points C ; or D ; correspond to the critical shear fracture stress state of the specimen.Therefore,the diameters of the two Mohr circles represent tensile or compressive shear frac-ture strength( T F or C F)and can be deduced asT F2 01 T 2pT;(3a)C F2 01 C 2pÿ C:(3b)From Fig.1(e),tensile and compressive shear fracture angles T and C can be expressed asT arctan1 T 2qT >45 ;(4a)C arctan1 C 2qÿ C <45 :(4b) Therefore,the shear fracture strengths( T F or C F)and shear fracture angles( T and C)are functions of 0, C, and T( 0= 0).From Eqs.(4a)and(4b)or Fig.1(e),it is apparent that T is larger than45 ,as given by Eq.(1). This explains why the shear fracture of most BMG ma-terials does not occur at the maximum shear stress plane[7–20].Fractography observations showed that the typical compressive shear fracture surface exhibits only a vein-like structure with a rather uniform arrangement,show-ing a pure shear fracture mode[Fig.2(a)].This vein structure has been widely observed before and was attrib-uted to melting of metallic glasses[12–16,20].However, on the tensile fracture surfaces,two typical features,i.e., some round cores and radiating veins,are observed for thefirst time,as shown in Fig.2(b).W e consider that (i)the veins originate from the cores and propagateTABLE pressive and tensile fracture angles C and T for different BMG materials.Investigators Composites CHe et al.[10]Zr52:5Ni14:6Al10Cu17:9Ti540 –45Lowhaphandu et al.[14]Zr62Ti10Ni10Cu14:5Be3:541:6 2:1Donovan[19]Pd40Ni40P2041:9 1:2Wright et al.[20]Zr40Ti14Ni10Cu12Be2442Present results(I)Zr59Cu20Al10Ni8Ti343 Compressive Zr55Cu30Al10Ni540 –43 Fracture Zr54:5Cu20Al10Ni8Ti7:542Zr52:5Ni14:6Al10Cu17:9Ti542Zr55Cu30Al10Ni5 1ÿ0:95Ta531 –33Ti56Cu16:8Ni14:4Sn4:8Nb827Ti56Cu16:8Ni14:4Sn4:8Nb8Split( 0 )Zr55Cu30Al10Ni5a Break or splitTi50Cu20Ni23Sn7Break or split Investigators Composites TAlpas et al.[7]Ni78Si10B1255Bengus et al.[8]Fe70Ni10B2060Davis and Y eow[9]Ni49Fe29P14B6Si253He et al.[10]Zr52:5Ni14:6Al10Cu17:9Ti555 –65Inoue et al.[11]Zr65Ni10Al7:5Cu7:5Pd1050Inoue et al.[12]Cu60Zr30Ti1054 Tensile Liu et al.[13]Zr52:5Ni14:6Al10Cu17:9Ti553 –60 Fracture Lowhaphandu et al.[14]Zr62Ti10Ni10Cu14:5Be3:557 3:7 Megusar et al.[15]Pd80Si2050Mukai et al.[16]Pd40Ni40P2056Noskova et al.[17]Co70Si15B10Fe560Takayama[18]Pd77:5Cu6Si16:550 –51Present results(II)Zr52:5Ni14:6Al10Cu17:9Ti556Zr59Cu20Al10Ni8Ti354Zr52:5Ni14:6Al10Cu17:9Ti5a 90Zr59Cu20Al10Ni8Ti3a 90a The three metallic glassy materials were annealed at420 C for5,10,and20min for partial or full crystallization.towards the outside in a radial mode;(ii)tensile fracture does not occur in a pure shear mode and is distinctly different from the compressive fracture feature.These observations further provide powerful evidence for the significant difference in the fracture features of BMGs induced by compression and tension.Therefore,both the macroscopic stress state and the micrometer-scale frac-ture feature of BMGs support the deviation of T and C from the maximum shear stress plane due to a restraining effect under normal compressive stress or a promotion effect under normal tensile stress.Besides the shear fracture,there are some other frac-ture modes for BMG materials.For example,annealed Zr59Cu20Al10Ni8Ti3and Zr52:5Cu17:9Al10Ni14:6Ti5,which contain a high volume fraction of nanocrystals(V f> 50%),always show nearly zero tensile fracture strength and their fracture surfaces are approximately perpendicu-lar to the loading axis[Fig.3(a)].Another typical failure mode is that the specimens often break into many small pieces rather than fracture in a shear mode,when the partially or fully crystallized metallic glass(for example, Zr55Cu30Al10Ni5)is subjected to compressive loading.In some cases,BMG composites containing dendritic phases (Ti50Cu23Ni20Sn7or Ti56Cu16:8Ni14:4Sn4:8Ta8)split into several parts[Fig.3(b)]or fracture along a plane nearly parallel to the compressive axis[Fig.3(c)],which is different from the wing cracks proposed by Ashby et al.[28].A similar fracture phenomenon,such as buckling and consequent delamination,also occurs in other BMG composites[24–27].This indicates that BMG composites with nanocrystals,hard particles,strongfibers,or den-dritic phases do not always fracture in a shear mode. Here,we define the two observed failure modes as normal tensile fracture(TF)with a fracture plane perpendicular to the loading direction and distensile fracture(DF)in a break or splitting mode with a fracture plane parallel to the loading direction,respectively.Normal TF of the crystallized metallic glasses always occurs at very low tensile fracture strength and can be schematically illustrated as in Fig.3(d).The tensile stress Mohr circle a is very small and the tensile strength T is approximately equal to 0.Therefore,the critical tensile fracture line AB is nearly perpendicular to thenormal FIG.2.Typical shear fracture surface features of metallicglasses under(a)compression and(b)tension.FIG.3.(a)Normal tensile fracture of the crystallizedZr59Cu20Al10Ni8Ti3metallic glass containing brittle Zr2Cunanocrystals.(b)Fracture of Ti50Cu20Ni23Sn7metallicglassy composite containing dendritic precipitates under com-pression.(c)Split of Ti56Cu16:8Ni14:4Sn4:8Ta8metallic glassycomposite containing dendritic precipitates under compression.(d)V ariation of the critical compressive and tensile fractureconditions due to the change in the microstructure of the BMGmaterials.FIG.1.Typical shear fracture modes of BMG materialsunder tension and pressive shear frac-ture of(a)Zr55Cu30Al10Ni5metallic glass and(b)Ti56Cu16:8Ni14:4Sn4:8Nb8composite containing dendritic pre-cipitates;(c)tensile shear fracture of Zr52:5Ni14:6Al10Cu17:9Ti5metallic glass;(d)illustration of shear and normal stresses onany plane of a specimen under compression or tension;(e)schematic illustration of the critical tensile and compressivefracture lines(AB and AE)and the stress distribution on the twoMohr circles under tension and compression.stress axis and T is quite close to90 .From Eq.(4a), when T!1,normal tensile fracture will occur at an angle T close to90 .The great increase of T 0= 0!1can be attributed to the extremely low value of 0due to the crystallization treatment,which results in the formation of nanocrystals,such as brittle intermetal-lic compounds[10,11].The DF mechanism is difficult to understand and was never explained before.As shown in Fig.3(d)and Table I, since the compressive specimens break or split into two or several parts,here we define a critical distensile fracture stress D as shown in Fig.3(d).When D is smaller than the critical shear fracture stress C F,i.e.,D< C F2 01 C 2pÿ C;(5)the stress Mohr circle b in Fig.3(d)willfirst touch the distensile fracture line DE.In this case,the specimen will fracture in a distensile modefirst rather than in a shear mode,due to local cracking rather than macroscopic shear cracking of the composite.On the contrary,when D is higher than C F,the specimen will still fracture a shear mode.Therefore,the fracture plane of BMGs can make any angle from0 to90 with respect to the stress axis.The actual failure of a BMG material is a competitive process between shear,distensile,and normal tensile fracture, depending on the loading mode and the inhomogeneity of the microstructure.Normal tensile and distensile frac-ture can be regarded as two special cases of shear frac-ture.The multifarious fracture mechanisms can be attributed to either an inhomogeneous distribution of the strengthening phases in the composites or the weak bonding of the interfaces between the matrix and the strengthening phases.Since the strengthening phases can also change the shear fracture angles C and T of the BMG materials,the physical meaning of the two constants C and T must reflect the degree of homoge-neity in the microstructure of BMG materials.The occur-rence of distensile fracture or not depends on the values of C and T and is a result of the interactions between the strengthening phases and the matrix of the composites [21–27],which are substantially important for further optimizing the properties of the BMG composites through the control of the microstructure.With the development of advanced materials,new fracture features are of interest in the mechanics and physics of the materials.BMGs and their composites,as typical super-high-strength materials,display abnormal fracture behavior in comparison with traditional crystal-line materials.Thefinding of the new fracture modes of BMG materials and their fracture criteria will contribute to the optimum design and fabrication of new high-performance materials.For example,recently,we suc-cessfully fabricated novel BMG materials with high strength and good ductility through the control of composition and microstructure[27].W e suggest that the present fracture criteria can be widely employed for isotropic materials with high strength,such as advanced BMGs or the newly developed bulk nanostructural materials.Z.Z.and G.H.acknowledge the Alexander von Humboldt Foundation forfinancial support.This work was partially supported by the Special Funding for the National100Excellent Ph.D.Thesis(Z.Z.)provided by Chinese Academy ofSciences.*Corresponding author.Present 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1.the Iron Age 铁器时代2.covalent bonding 共价键，共价结合posites 复合材料4.crystal lattice 晶体点阵，晶格position and structure 成分和结构6.tensile strength抗拉强度，抗张强度7.ferrous metals 黑色金属8.gray cast iron 灰口铸铁9.austennitic stainless 奥氏体不锈钢10.weldability and hardenability 可焊性和可淬性11.refractory metals 难溶金属11.carbide and nitride碳化物和氮化物12.stiffness 刚度13.corrosion 腐蚀14.the Bronze Age 铜器时代15.metallic bonding 金属键，金属结合16.polymers 高分子材料17.ceramics and glasses 陶瓷和玻璃18.elementaty cell 晶胞19.direction indices晶向指数20.synthesis and processing 合成和加工21.yeild strength 屈服强度22.nonferrous metals 有色金属23.white cast iron白口铸铁24.martensitic stainless steels 马氏体不锈钢25.castability and formability 铸造性能与模锻性能26.titanium and nickel钛和镍27.precious metals 贵金属28.oxide and sulfide氧化物和硫化物29.die cast alloy压铸合金30.elasticity 弹性，弹力31.brittleness脆性32.fatigue strength 疲劳强度33.corrosion腐蚀34.annealing 退火35.high compressive strength 高压缩强度材料工程materials engineering 金属及其化合物metals and their alloys 面心立方晶格face-centered cubic lattice 材料塑性the plasticity of materials 普碳钢plain-carbon steels 陶瓷ceramics 合金元素alloying elements 表面处理surface treatment 金属物理性能the physical property of metals 材料科学materials science 金属材料metallic materials 体心立方晶格body-centered cubic lattice 材料的强度the strength of materials 有色金属nonferrous metals 合金钢alloy steels 铝及铝合金aluminums and aluminum alloys 加工硬化work hardening 热处理heat treated 金属力学性能mechanical propertyAbsorbed energy吸收功transition temperature转变温度modulus of elasticity弹性模量conductivity导电性thermal expansion热膨胀heat capacity 热容mold铸型rolling轧制forming 模压thermosetting ploymers热固性材料thermoplastic ploymers 热塑性材料stress versus strain应力应变pig iron生铁wrought iron熟铁steel malking 炼钢smelting熔炼blast furnace鼓风炉castability可锻性machinability机加工性nonmachinable不可机加工的hardenability可淬硬性nonmagnetic非磁铁alloyed steels合金钢anneal退火stree-corrsion cracking应力腐蚀断裂high-strength low-alloy steel高强度低合金钢cast iron alloys铸铁合金heat-treatable 可热处理的solubility溶解度thermo-mechanical 热加工性plain-carbon steel普碳钢electrolytic iron电解铁Introduction to materials材料概论coordination nunber配位数polycrystals多晶体anisotropy各向异性hexagonal close-packed structure 密排六方结构impact strength冲击强度tensile strength拉伸强度yield point屈服点utimate strength极限强度breaking strength破坏强度fracture toughness断裂韧度thoughness 韧性elastic limit弹性极限creep strength蠕变强度creep蠕变fatigue life 疲劳寿命corrosion resistance抗腐蚀性wear-resistance 耐磨性wear rate磨损率oxidation resistance抗氧化性imperfection缺陷austenitic马氏体martensitic 马氏体pearlite珠光体ferritic铁素体iron carbide渗碳体stainless steel不锈钢fracture 断裂compouds化合物specific strength比强度allotropic同素异形体reractory metals耐火材料anodize阳极电镀forging锻造casting铸造hardness硬度gray cast iron灰口铸铁magnetin磁性rust铁锈ingots铸锭malleable cast iron 可锻铸铁brittle materials脆性材料white cast iron 白口铸铁gears齿轮shafts轴weldable可焊接的unweldable不可焊接的weldablility可焊接性tool steels工具钢metallic bonding金属键covalent bonding共价键ionic bonding离子键hydrogen bonding 氢键crystal lattice晶格crystalline晶体amorphous非晶体packing factor致密度crystallographic indices结晶指数slip planes滑移面close-packed planes密排面elasticity弹性Elongation rate延伸率stress-rupture properties应力开裂reduction in area断面收缩句子：1,It is generally the behavior of materials is which limits the performance of machines and equipment.材料的性能通常限制着机器和设备的性能2,It is useful to consider the extent of metallic behaverior in the currently known range of chemical elements.在目前已知的化学范围内考察其金属性的程度是很有用的3,The packing factor is determined as the ratio of the volume of all elementary particles per elementary cell to the total volume of the elementary cell.填充因子的大小取决于每个晶胞中所有基本微粒的体积之和与整个晶胞体积之间的比率4,In most materials more than one phase is present, with each phase having its unique atomic arrangement and properties. Control of the type, size, distribution, and amount of these phases within the main body of the material provides an additional way to control properties of a material.在大多数材料中，往往存在着不止一种相，每一种相都有其各自的原子排列和特性。

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A B C D E F G H I G K L M N O P Q R S T U V W X Y Zapparent porosity显气孔率annual capacity 年产能alkali resistance 耐碱性airtight packing 密封包装alkali bursting 碱金属损毁antioxidant agent 抗氧化剂blast furnace runner 高炉铁钩binder phase 结合相brittle layer 脆性层bulk density体积密度bulk density （of granular material）颗粒体积密度bonding strength粘结强度blast furnace高炉bottom cooling 炉底冷却bath zone 熔池bottom joint of converter 炉底打结cooled jet 冷却水coal injection 喷煤closed porosity闭气孔率closed pores闭气孔chemical bond化学结合ceramic bond陶瓷结合cold compressive strength，cold crushing strength常温耐压强度crude steel 粗铁continuous operation 连续作业ceramic cup 陶瓷杯coal-tar bond 煤焦油结合cast house 铸铁场castable浇注料ceramic burner 陶瓷燃烧器checker chamber 蓄热室combustion chamber 燃烧室converter 转炉creep resistance 抗蠕变性能continuous casting 连铸combined blowing 复吹cold isostatic press 冷等静压continental Europe 欧洲大陆crucible 坩埚cooling system 冷却系统dry strength干燥强度direct bonding直接结合desiliconization脱硅dephosphorization 脱磷desulphurization 脱硫dry gunning mix 干式喷补料dome 炉顶expansion joint 膨胀缝even flow 均匀流入extractor system 排烟系统erosion resistance 耐磨损electric arc furnace 电弧炉expansion allowance 膨胀缝electrode ring 电极圈fireclay core 耐火粘土fused silica 熔融石英fused mullite电熔莫来石fireclay brick 粘土砖fracture behavior断裂性能flexural strength 抗折强度fused magnesia 电熔镁砂flake graphite 鳞片石墨free flow 自流式grain bulk density颗粒体积密度gas distribution chamber 气室gunning mix 喷补hydraulic bond水化结合hearth diameter 炉缸直径hot blast 热风炉homogeneity of the material 材料的均匀heat capacity 热容hot strength 热态强度inject gun 喷枪intensive cooling 强冷却iron tap hole 出铁口ill-defined setting 无法确定的固化long nozzle 长水口little known鲜为人知mechanical property 力学性能mix 炮泥main trough 主沟mineralogical composition 矿相组成method of manufacture 生产工艺metallic additives 金属添加物nitride bond 氮化物结合open pores开口气孔open flame coal 非结焦性煤open hearth furnace 平炉oxygen partial 氧分压pig iron ladle 铁水包preventive measure 预防性措施PCE （pyrometric cone equivalent）标准锥相当值principal crystalline phase主晶相pyrometric reference cone标准测温锥preliminary figure 预计数字pig iron 生铁plasma flame 等离子火焰personnel constraint 人为因素permanent lining 永久衬pressing technology 成型技术preventive gunning 防护性喷补performance evalution自评报告resin bond 树脂结合refractory，refractory product，refractory material耐火材料reference temperature弯倒温度runner 铁钩ramming mix 捣打料redox resistance 抗氧化还原residual carbon 碳含量slag granulation plant 渣处理厂specific consumption 单耗sag flow 渣沟stave cooling 冷却壁spray cooling 喷射冷却silicate bonding硅酸盐结合secondary crystalline phase次晶相shut down 休风，停炉service life 使用寿命salamander 炉底结瘤skimmer 撇渣器slag line（zone）渣线selective line 综合炉衬stress corrosion cracking 应力腐蚀裂纹safety lining 永久衬secondary memllurgy二次精炼shaft kiln 竖窑structural spalling 结构式剥落special element 特殊原件slag stopper 挡渣室slag conditioner 造渣剂steel production unit 炼钢单元sidewall area 炉墙saturation level 饱和水平stopper rod 塞棒thermal diffusivity热扩散系数thermal expansion coefficient热膨胀系数thermal conductivity 热导率true porosity真气孔率true density真密度temperature of collapse弯倒温度transition zone 过渡区thermal shock resistance 抗热震性trouble free 流畅throughput 风量tar-epoxy resin 焦油环氧树脂tapping area 出钢区tannel kiln 隧道窑taphole filling material 引流砂tap-to-tap 冶炼周期threshold value 临界值toxic substance 有毒物质thixotropic 触变性unshaped product 不定形产品upper border 上限water absorption吸水率wear resistance 抗侵蚀working lining 工作衬。

一、提名项目：考虑非均匀结构效应的金属材料剪切带二、提名意见：该项目以颗粒增强金属基复合材料和非晶合金为模型系统，突破经典的热塑剪切带理论框架，发展了位错机制依赖的应变梯度本构，揭示了蕴含的非均匀结构通过应变梯度效应对热塑剪切带形成具有强烈驱动作用；建立了包含多过程耦合与时空多尺度的剪切带新理论，澄清了非晶合金剪切带形成机制长期广泛的国际争议，得到了剪切带失稳判据、协同演化、特征厚度以及诱致断裂机理等一系列原创性成果。

该项目8篇代表性论文共被《Nature Materials》《Physical Review Letters〉、《Progress in Materials Science等SCI重要刊物他人引用393次，引用者包括国内外科学院或工程院院士、权威杂志主编、领域知名学者等。

项目研究成果系统揭示了材料内禀非均匀结构效应如何影响甚至颠覆热塑剪切带的传统认知，显著推动了剪切带理论的发展，在国际上产生了重要的学术影响。

提名该项目为国家自然科学二等奖。

三、项目简介剪切带是一类广泛存在的塑性变形局部化失稳现象。

本征上，具有特征厚度的剪切带是一种远离平衡态的动态耗散结构，其涌现与演化是材料内部多种速率依赖耗散过程高度非线性耦合控制的时空多尺度问题。

传统金属材料剪切带经百余年研究，逐渐形成了以热软化为主控机制的热塑剪切带理论，并获得了广泛的应用。

随着人们对高性能材料的不懈追求，众多内蕴微纳尺度非均匀结构的新型金属材料不断发展，其中代表性的有微米尺度颗粒增强的金属基复合材料和纳米尺度结构非均匀的非晶合金。

由于不考虑材料结构效应，经典热塑剪切带理论在描述这些新型金属材料的剪切带行为时，遇到了前所未有的挑战。

为此，该项目团队以颗粒增强金属基复合材料和非晶合金为模型材料，研究了材料内禀的非均匀结构效应如何影响甚至颠覆热塑剪切带的传统认知，显著推动了剪切带理论的发展，形成了具有鲜明特色的系统性的原创研究成果。

Trans. Nonferrous Met. Soc. China 30(2020) 2590−2598Effect of rapid cold stamping on fracture behavior oflong strip S′ phase in Al −Cu −Mg alloyCai-he FAN 1,2, Ling OU 1,2, Ze-yi HU 1, Shu WANG 3, Jun-hong WANG 31. College of Metallurgy and Material Engineering, Hunan University of Technology, Zhuzhou 412007, China;2. Anhui Jianye Science and Technology Co., Ltd., Huaibei 235000, China;3. 208 Research Institute of China Ordnance Industries, Beijing 102202, ChinaReceived 6 January 2020; accepted 16 July 2020Abstract: High-angle annular dark-field scanning transmission electron microscopy and selected area electron diffraction techniques were used to study the mechanism that underlies the influence of rapid cold-stamping deformation on the fracture behavior of the elongated nanoprecipitated phase in extruded Al −Cu −Mg alloy. Results show that the interface between the long strip-shaped S′ phase and the aluminum matrix in the extruded Al −Cu −Mg alloy is flat and breaks during rapid cold-stamping deformation. The breaking mechanisms are distortion and brittle failure, redissolution, and necking. The breakage of the long strip S′ phase increases the contact surface between the S′ phase and the aluminum matrix and improves the interfacial distortion energy. This effect accounts for the higher free energy of the S′ phase than that of the matrix and creates conditions for the redissolution of solute atoms back into the aluminum matrix. The brittle S′ phase produces a resolved step during rapid cold-stamping deformation. This step further accelerates the diffusion of solute atoms and promotes the redissolution of the S′ phase. Thus, the S′ phase necks and separates, and the long strip-shaped S′ phase in the extruded Al −Cu −Mg alloy is broken into a short and thin S′ phase.Key words: Al −Cu −Mg alloy; rapid cold stamping; nanoprecipitate; fracture behavior; breaking mechanism1 IntroductionA heterogeneous microstructure that consists of hardened precipitates and dispersoids is the primary strengthening source of Al −Cu −Mg alloys [1,2]. The S′ phase is a key strengthening precipitate in high-strength Al −Cu −Mg alloys andexhibits an orthorhombic structure with space groupCMCM [3−6]. Strong plastic deformation can improve the precipitation phase characteristics, such as the morphology, size, distribution, andorientation relationship with the matrix, of aluminum alloy, improve its microstructure, and obtain microscale or even nanoscale fine-grainedstructures [6−11]. STYLES et al [12] studied the relationship between the resolution sequence of a supersaturated solid solution and the formation phase of the equilibrium S phase in Al −Cu −Mg alloy. The formation time of the S phase at high temperature is considerably shorter than that at low temperature. YANG et al [13] studied the effects of applied stress on the aging kinetics and precipitatedphase morphology of Al −Cu −Mg alloy and found that applied stress promotes the precipitation of the S′ phase by changing the force contrast between the S′ phase and the θ′ phase in competitiveprecipitation. NOURBAKHSH and NUTTING [14] found that the flaky θ′ phase is severely bent and broken after the large reduction cold-rolling ofA1−4%Cu alloy. MURAYAMA et al [15] observed a needle-like θ′ transition phase in the equal-Foundation item: Project (19A131) supported by Key Scientific Research Project of Hunan Province, China; Project (2019JJ60050)supported by the Natural Science Foundation of Hunan Province, ChinaCorresponding author:LingOU;Tel:+86-731-22183432;E-mail:****************DOI: 10.1016/S1003-6326(20)65404-8diameter extrusion deformation of A1−Cu alloy, which gradually decomposed into short-chain particles during several equal-angle extrusion processes until it redissolved into the matrix. HUANG et al [16] used cold rolling to change the dislocation cell structure in Al−Cu−Mg alloy with a low Cu/Mg mass ratio and obtained a high-density dispersed fine S′′ phase. ZHAO et al [17] found that the S phase continuously converts to the Ωphase with an increase in cold deformation degree in the cold rolling deformation of 2024 aluminum alloy. ZHANG et al [18,19] studied the effect of multidirectional compressive deformation on the precipitation phase of Al−Cu alloy and found that the precipitation sequence of the supersaturated solid solution formed by the redissolution of the precipitates under strong plastic deformation is related to heating temperature, deformation, and postdeformation grain size.It can be seen that the existing literature is mainly to study the characteristics of the precipitated phase and the re-dissolution phenomenon under the conditions of conventional plastic deformation. There is still a lack of rapid cold punching as the method of plastic deformation, and the impact of the nano-precipitated phase on the microstructure of the alloy is discussed. In order to study the mechanism that underlies the influence of rapid cold-stamping deformation on the fracture behavior of the elongated nanoprecipitated phase in extruded Al−Cu−Mg alloy, this experiment is based on the rapid solidification of fine-grained Al−Cu−Mg alloy billet prepared via spray forming. The fracture behavior of the long strip S′phase in Al−Cu−Mg alloy during rapid cold stamping is also studied. The analysis of the breaking mechanism of the S′phase lays a theoretical foundation for the systematic study of the evolutionary law and strengthening mechanism of the long-shaped S′phase.2 ExperimentalA rapidly solidified fine-grained Al–Cu–Mg alloy cylindrical ingot was prepared on a self-developed SD380 large-scale injection molding apparatus. Table 1 provides the chemical composition of the alloy. The cylindrical ingot was extruded into a d30 mm round bar on a 1250 T extruder at 450 °C with an extrusion ratio of 15:1. The round bar was cut into a small cylinder with dimensions of d30 mm × 20 mm and rapidly cold-stamped for four passes at 25 °C. The sketch maps of rapid cold stamping are presented in Fig. 1. The small cylindrical sample was first placed in the mold, and the first small punch was used to rapidly cold-stamp along the center of the sample to complete the first pass of rapid cold forming. Then, a large-diameter punch was used for cold-stamping along the first pass forming hole, and four passes of rapid cold-stamping were completed in sequence. The parameters of rapid cold stamping are listed in Table 2.Table 1 Composition of alloy studied (wt.%)Cu Mg Mn Si Fe Al3.91 1.63 0.38 <0.03 <0.03 Bal.The samples for microstructure observation were obtained from the center wall of the cylinder fabricated through rapid cold forming. The morphology, size, and distribution of the nanoprecipitated phase of the sample were analyzed through JEM-F200 transmission electronFig. 1 Sketch maps of rapid cold punching (a—Sample; b—Drawing die; c—Punch)Table 2 Technological parameters of rapid cold punching Pass number Diameter/mm Velocity/(mm·s−1) One 10 30Two 14 25Three 20 20Four 27 15 microscopy (TEM). The transmission samples were mechanically prethinned to approximately 80 μm and then subjected to twin-jet electropolishing. The electrolytes were nitric acid and methanol (volume rati o of 1:3) at a temperature of less than −25 °C. The electron microscopy parameters of the high-angle annular dark-field scanning transmission electron microscope (HAADF–STEM) were as follows: acceleration voltage of 200 keV, electron beam half convergence angle of 10 mrad, and beam spot diameter of 0.20 nm.3 Results3.1 As-extruded sampleFigure 2 shows a HAADF−STEM images of the as-extruded Al−Cu−Mg alloy. The elongated S′phase in the as-extruded alloy is regularly distributed in the aluminum matrix; the lengthwise dimension of the elongated S′phase is less than 300 nm, and a number of dislocations are distributed around the precipitated phases (Fig. 2(a)). Figure 2(b) shows the interface between the elongated S′phase and the aluminum matrix observed along the [001]Al direction. The shape of the S′phase is regular, and the interface with the aluminum matrix along the length direction of the S′phase is flat.3.2 Rapid cold punching sampleFigure 3 depicts a HAADF−STEM image of the extruded Al−Cu−Mg alloy after two passes of rapid cold-stamping deformation. The number of S′phases increases drastically, and the S′phases are irregularly distributed in the aluminum matrix. The length of the S′phase after two passes of rapid cold-stamping deformation is the same as that of the extruded sample (Fig. 3(a)). The morphology of most S′phases has changed, and the interface between the S′phase and the aluminum matrix is not flat. The S′phase distorts and breaks during rapid cold-stamping deformation (Fig. 3(b)). The high-magnification field TEM image shows that the long S′phase is broken, the broken interface is visible, and one S′ phase is broken into several parts (A, B, C, and D). Each part undergoes drastic twisting, B is rotated by the α1 angle with respect to A, C is rotated by the α2 angle with respect to B, and D is rotated by the α3angle with respect to C (Fig. 3(c)).Figure 4 shows a HAADF–STEM image of the extruded Al–Cu–Mg alloy after three passes of rapid cold-stamping deformation. As shown in the figure, the number of S′phases in the alloy is reduced significantly compared with those in the specimens after two passes of rapid cold-stamping deformation. However, a large number of short rod- shaped S′ phases are observed, and the S′phasesFig. 2 HAADF−STEM images of S′ precipitates in as-extruded Al−Cu−Mg alloy viewed along [001]Al directionFig. 3HAADF−STEM images of S′precipitates in Al−Cu−Mg alloy undergoing two passes viewed along [001]Al directionremain irregularly distributed in the aluminum matrix (Fig. 4(a)). As the degree of the deformation increases, the morphology of the S′phase changes, the degree of distortion is aggravated, and the broken part of the S′ phase begins to separate from Fig. 4HAADF−STEM images of S′precipitates in Al−Cu−Mg alloy undergoing three passes viewed along [001]Al directionthe base metal (Fig. 4(b)). The high-magnification TEM image shows that the region of the long strip- shaped S′phase is necked, and L1is considerably larger than L2. The ends of necked region tend to separate as the value of L1/L2increases (Fig. 4(c)).Cai-he FAN, et al/Trans. Nonferrous Met. Soc. China 30(2020) 2590−2598 2594Figure 5 shows a HAADF−STEM image of the extruded Al−Cu−Mg alloy after four passes of rapid cold-stamping deformation. This figure shows that the number of S′phases in this sample has increased compared with those in the sample after three passes of rapid cold-stamping deformation.Fig. 5HAADF−STEM images of S′precipitates in Al−Cu−Mg alloy undergoing four passes viewed along [001]Al direction The S′phase is short and thin, and remains irregularly distributed in the aluminum matrix (Fig. 5(a)). After four passes of rapid cold-stamping deformation, the long strip-shaped S′ phases in the extruded sample are nearly broken into short rods, resulting in a considerable reduction in the size of the S′ phase in the aluminum matrix (Fig. 5(b)). The high-magnification field TEM image shows that the necked component of the long strip-shaped S′ phase has disappeared. The ends of the necked region have separated, and the distance between the two parts is L. The S′phase has broken into several independent parts (Fig. 5(c)).4 Analysis and discussion4.1 Breaking mechanism of nano-precipitatedphase during rapid cold stampingIn this experiment, rapid cold stamping deformation has three major characteristics: low deformation temperature (25 °C), large strain, and high strain rate. In accordance with the general criteria of PUGH and related research, the Al2CuMg phase is a typical brittle phase [20,21]. Therefore, combined with the results of this experiment, the breaking mechanisms of the long strip-shaped S′phase in the extruded Al−Cu−Mg alloy induced by rapid cold-stamping deformation are as follows.(1) Distortion and brittle fractureThe long strip-shaped S′phase acts as the major strengthening phase of the extruded Al−Cu−Mg alloy and is regularly distributed in the aluminum matrix, and the S′phase is flat with the aluminum matrix (Fig. 2). During rapid cold stamping deformation, the soft aluminum matrix is subject to severe shear deformation, which causes the hard S′phase to distort under the strong shear deformation force. The torsion angle between parts C and D is α3. When the deformation force reaches a certain level, i.e., when the torsion angle reaches a certain value (e.g. α1 or α2), brittle S′ phase fractures and interfaces are formed between A and B or B and C (Fig. 3(c)). With the distortion and brittle fracture of the long strip-shaped S′phase, the drastic increase in brittle interfaces considerably increases the contact surface between the precipitated phase and the aluminum matrix. The interfacial distortion energy between the precipitated phase and the aluminum matrix is improved, and diffusionconditions for the solute atoms are created. TheCai-he FAN, et al/Trans. Nonferrous Met. Soc. China 30(2020) 2590−2598 2595redissolution of the S′ phase during rapid cold-stamping deformation accelerates. The distortion and brittle fracture model of the long strip-shaped S′ phase is shown in Fig. 6.(2) Resolving and neckingFigure 7 presents a HAADF–STEM image showing the resolved step of the S′ phase and a model of redissolution and necking in the sample subjected to three passes of rapid cold stamping. The resolved step is clearly observable in the S′ phase, and the arrows in the figure are the resolving direction. The resolved step is the product of the twisting and brittle fracture of the S′ phase during strong plastic deformation. Given that the distortion and breakage of the long S′ phase considerably improve the interfacial distortion energy between the S′ phase and the aluminum matrix, the solute atoms located at the region of the S′ phase with high distortion energy are highly likely to be redissolved. The local accelerated redissolution of the S′ phase results in a significant reduction in the size of the S′ phase in this region, i.e., the formation of a neck (Fig. 4(c)), which further increases the interfacial distortion between the precipitated phase and the aluminum matrix and can accelerate redissolution at the necked position. Necking disappears when the value of L 1/L 2 is infinite, and the disappearance of the neck causes the separation of the two parts of the precipitated phase (Fig. 5(c)).The presence of high numbers of resolved steps in the S′ phase is associated with the large contact surface of the broken particles within the aluminum matrix. The diffusion of solute atoms is facilitated. In accordance with the classical thermodynamic view [22], a certain equilibrium vacancy concentration in the crystal exists at anyFig. 6 Distortion and brittle fracture model of long strip-shaped S′ phase: (a) As-extruded; (b) Two-passFig. 7 HAADF −STEM image (a) and necking model (b) of precipitated phase in Al −Cu −Mg alloy undergoing three passestemperature and changes positions continuously. The existence and movement of vacancies create conditions for atomic diffusion. The dislocation density increases significantly during rapid cold deformation. Dislocations will generate a large number of vacancies in the matrix during delivery and movement, and each time, a rapid cold shear deformation will produce a large number of new vacancies. Simultaneously, the long strip-shaped S′phase is constantly twisted and broken during rapid cold-stamping deformation, and the contact interfaces between S′phase and aluminum matrix increase and are accompanied by a large number of vacancies. The number of vacancies undergoing strong deformation increases sharply [19,23], the diffusion rate of solute atoms can be increased by five orders of magnitude, and volume diffusion can be increased by eight orders of magnitude. The alloy was prepared through spray-forming rapid solidification technology and had fine grains and uniform structure. The average grain size was app roximately 5 μm. During rapid cold-stamping deformation, the grains are further refined into nanoscale, and thus, the grain boundary area increases. It also provides an atomic diffusion channel for the redissolution of the precipitated phase, and this channel considerably promotes the redissolution of the precipitated phase.4.2 Fracture mechanism of nano-precipitatedphase during rapid cold stampingUnder the condition of rapid cold-stamping deformation, the long strip-shaped S′phase in the extruded Al−Cu−Mg alloy undergoes intense distortion, brittle fracture, redissolution, necking, and separation (Figs. 3−5). Consequently, the long strips of S′ phases transform into short rods or even redissolve and disappear. In contrast with that of the extruded sample, the brittle fracture of the long strip S′ phase remarkably increases the contact surface of the S′ phase and the matrix, and interface distortion can be drastically improved. This phenomenon results in the free energy of the S′phase being higher than that of the matrix. The disturbance of the energy balance between the reprecipitated phase and the matrix creates conditions for the redissolution of the solute atoms into the matrix. The binding energy (E coh) and the formation enthalpy of ∆H[20,21] between the Al2Cu and Al2CuMg phases are calculated using the first principle equations:AB A Bcoh total A atom A atom[(1)]E E x E x E=−+− (1)AB A Btotal A solid A solid[(1)]H E x E x E∆=−+− (2)where ABtotalE is the average energy per atom ofeach intermetallic compound; AatomE and BatomE are the energy of free atoms A and B, respectively;AsolidE and BsolidE are the average energy per atom in stable elemental A and B, respectively; x A represents the mole fraction of atom A in the compound.V ASILS et al [24] believed that two types of processes will occur when precipitates are completely dissolved into the matrix. The first process is lattice transformation, wherein the lattice of the precipitated phase changes into the matrix lattice. The second process is the diffusion of the redissolved solute atoms in the matrix; this phenomenon increases the uniformity of the supersaturated solid solution. The results show that the S′ phase in the as-extruded alloy redissolves into the aluminum matrix after three passes of rapid cold-stamping deformation and becomes refined. The rapid cold-stamping deformation temperature of the alloy in this test is low, and if solute atoms uniformly disperse into the matrix by diffusion to redissolve the precipitate phase, the redissolution rate must be extremely slow.In accordance with the solid-state phase transition theory [22], the critical size of the nucleus (r c) is expressed as follows:cr f UGσ=−∆(3)where σis the surface free enthalpy, ∆G is the difference in volume free enthalpy between the parent phase and the precipitated phase, and U is the strain energy produced by the two phases. With the fragmentation of the precipitated phase, σ, ∆G, and U change significantly, resulting in an increase in the critical size (r c) of the precipitated phase. The redissolution amount and rate of the precipitated phase increase.In summary, the redissolution of nano- precipitates during strong plastic deformation plays an important role in structural evolution. The distortion and brittle fracture of the S′ phase during rapid cold-stamping deformation results in a resolved step, which accelerates the redissolution ofCai-he FAN, et al/Trans. Nonferrous Met. Soc. China 30(2020) 2590−2598 2597the S′ phase. The partial redissolution of the phase promotes the further necking of the S′phase and results in the separation of various parts of the S′phase, resulting in the fragmentation of the elongated S′phase in the extruded alloy and its transformation into a short and thin S′ phase.5 Conclusions(1) The interface between the strip-shaped S′phase and the aluminum matrix in the extrudedAl−Cu−Mg alloy is flat and breaks during rapid cold-stamping deformation. The fracture mechanisms are distortion, brittle fracture, redissolution, and necking.(2) The brittle fracture of the long S′phase increases the contact surface between the S′phase and the aluminum matrix, improves the interfacial distortion energy, and accounts for the higher free energy of the S′ phase than that of the matrix. The disturbance of the energy balance between the S′phase and the matrix creates diffusion conditions for the redissolution of the solute atoms into the matrix.(3) The redissolution of nanoprecipitates during strong plastic deformation plays an important role in structural evolution. The break of the local S′phase accelerates its redissolution, resulting in the fragmentation of long S′phase into short and thin S′phase.References[1]CHENG S, ZHAO Y H, ZHU Y T, MA E. Optimizing thestrength and ductility of fine structured 2024Al alloy bynano-precipitation [J]. Acta Mater, 2007, 55(17): 5822−5832. 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时效处理对2024-T3搅拌摩擦焊接头组织及性能的影响乔文广1，杨新岐1，董春林2，吴海亮11天津大学材料学院，天津 (300072)2北京赛福斯特技术有限公司，北京 (100024)E-mail：qiaowgsx@摘要：本文对AA2024-T3进行了搅拌摩擦焊，并焊后对接头进行了时效处理，其中时效℃，45/12h℃。

采用了金相、硬度、拉伸、塑性弯曲试验规范为：自然时效35天，180/10h方法对处理过的接头经行了评定。

研究结果表明，三种时效规范的接头，其各区域的组织微℃的焊缝析出的S相观上也未能发现明显的区别，但是经过X-Rad衍射分析，发现45/12h℃的分布趋势一样，只是比其它两种的多。

从接头整体硬度分布来看，自然时效与45/12h℃的接头硬度分布毫无规律，但热机影响区的硬度升高并大于后者整体比前者高；180/10h焊缝硬度。

由于接头存在缺陷，导致三种时效制度的接头抗拉强度、延伸率下降。

关键词：2024铝合金，搅拌摩擦焊，时效处理，组织及性能中图分类号：TG41.引言搅拌摩擦焊是一种新的金属焊接技术，解决了铝合金等低熔点材料的焊接，获得了无气孔、裂缝等缺陷的高质量焊缝[1-2]。

目前，搅拌摩擦焊已经被广泛的应用与航空航天、船舶、铁路、汽车等领域。

国内外评定铝合金FSW接头时，主要考察焊态条件下接头区域微观组织结构、力学性能和疲劳性能，很少考虑焊后时效热处理对FSW接头使用性能的影响。

铝合金的特点是具有时效强化效应，即使在自然放置状态下其性能也会随时间而变化。

因此，研究时效工艺对FSW接头性能的影响具有重要工程意义。

但是，目前，关于基体材料的时效行为在一些国外文献中有介绍[3-13]，但是关于搅拌摩擦焊接头的特别少，尤其是关于2024的，几乎没有。

2.试验方法试验材料为轧制状态的AA2024-T3，其化学成分如表1。

材料原始厚度为2.6mm，为了减少包铝对接头的影响，对材料进行了化学洗涤，化洗后材料的厚度为2.4mm，其机械性能如表2。

第38卷第4期2023年8月安 徽 工 程 大 学 学 报J o u r n a l o fA n h u i P o l y t e c h n i cU n i v e r s i t y V o l .38N o .4A u g.2023文章编号:1672-2477(2023)04-0008-06收稿日期:2022-11-03 基金项目:芜湖市揭榜挂帅基金资助项目(2021h g 08)作者简介:鲍正浩(1997-),男,安徽六安人,硕士研究生㊂通信作者:黄仲佳(1979-),男,广西贵港人,教授,博士㊂430不锈钢激光选区增材件的组织及抗拉性能鲍正浩,黄仲佳*,张晨阳,陶 靖(高性能有色金属材料安徽省重点实验室,安徽芜湖 241000)摘要:为了探明430不锈钢增材制造件的打印态和固溶态的抗拉性能,采用激光选区增材制造制备430不锈钢增材样品,研究打印态和固溶处理态的组织㊁拉伸性能㊂结果表明:430不锈钢具有良好的激光选区增材成形性能;打印态的晶界析出物较多,固溶处理消除了晶界析出;打印态的抗拉强度高达700M p a ,固溶态的抗拉强度降低至550M p a ,但固溶态具有更好的塑性;断口形貌表明,打印态拉伸断裂受到剪切力影响较大,固溶处理态为典型韧性断裂㊂关 键 词:激光选区增材制造;不锈钢;组织;拉伸性能中图分类号:T G 142.23 文献标志码:A 增材制造技术是通过计算机三维模型控制逐层智能制备而获得越来越多的研究[1]㊂增材制造技术具有自身独特的技术优势,如与铸造㊁锻造等传统制备技术相比较,具有更少的加工工序和更高的材料利用率[2]㊂其中激光选区熔覆工艺(S L M )是增材技术中比较成熟的一种,与其他增材制造技术相比,S L M 技术较为成熟且已研制出商用设备[3-4]㊂激光选区增材制备可以获得比铸造和锻造更优越的结构和机械性能[5]㊂430不锈钢是一种铁素体不锈钢钢种,具有高强度㊁抗高温氧化及优秀的抗腐蚀性能,被广泛应用于工具㊁航空㊁海洋㊁汽车㊁核电及国防领域[6]㊂在430不锈钢的应用领域中,有很多复杂结构的工件,如汽车发动机排气歧管,其传统成型工序多,有些复杂结构传统工艺实习难度大㊂而增材制造在复杂件制备中具有很大的优势,采用增材制备430不锈钢复杂件将是一种可行的方法㊂不锈钢的增材制造受到了研究者们的广泛关注,目前430不锈钢㊁316不锈钢㊁304不锈钢等都有研究报告㊂金成嘉等[7]研究发现激光冲击后430铁素体不锈钢的表面发生了塑性变形并产生晶粒细化;强化处理后材料表面的显微硬度得到明显提升;激光冲击后,A I S I 430铁素体不锈钢表面出现残余压应力层㊂K r a k h m a l e v 等[8]研究表明,420马氏体不锈钢在选择性激光熔融过程中打印件上层的硬度为750H V 0.2,组织中包含(21±1.2)v o l .%的奥氏体相;最终的大块打印件微观结构由热分解马氏体组成,硬度为500~550H V 0.2,奥氏体含量异常高达到(57±8)v o l .%㊂丝状电弧增材制备的不锈钢零件与不锈钢锻造件相比,增材制备件具备无缺陷的冶金结合的微观结构以及更好的机械性能[9]㊂对于激光选区熔融法制备的不锈钢金属基复合材料样品,T i N 的加入对密度和硬度有很大的影响,当T i N 含量增加时,打印件的密度和硬度都快速增加,这主要归因于其对粉末的激光吸收率和液态金属的润湿性的影响[10]㊂K a n g 等[11]研究增材制造316L 片材发现弹性各向同性的弹性模量㊁泊松比㊁极限应力和伸长率,但塑性力学各向异性,表现出较高的极限抗拉强度和较好的延展性㊂W a n g 等[12]采用电弧增材制备316不锈钢件,样品极限抗拉强度沿水平方向制造的试样大于540M p a ,高于轧制件㊂为了探明430不锈钢增材制造工艺性及其抗拉性能,采用激光选区增材制造制备430不锈钢增材样品,研究打印态和固溶处理态的组织㊁拉伸性能;分析打印件的断口形貌,探明固溶处理对430增材件的组织及拉伸性能的影响㊂1 实验方法1.1 实验材料及样品制备实验材料为430不锈钢球形粉末,粒度范围:15~53μm ,真空雾化法制备㊂其化学成分如表1所示㊂表1 430不锈钢粉末化学成分(质量分数,%)元素C r M n N i S i C O S P F e 含量16.000~18.000≤1.000≤0.600≤0.750≤0.120≤0.039≤0.030≤0.040余量在打印制备样品之前,使用烘箱进行不锈钢粉末烘干,烘干温度为100℃,时间为2h ㊂再使用S L M 125金属3D 打印机制备尺寸为100mm×50mm×20mm 不锈钢金属块,打印工艺参数主要调整激光功率(P )㊁扫描速度(v )㊁层厚(t )及扫描间距(h ),层间转角固定选用35°及扫描方式为条带扫描㊂打印样品的固溶热处理实验在真空/气氛管式电炉中进行,设备型号为S X 2-4-10A ㊂固溶处理工艺为:860℃保温20m i n ,水冷至室温㊂ 图1 拉伸试样示意图1.2 样品检测及表征采用洛氏硬度计测量样品的硬度;采用万能拉伸实验机进行拉伸实验,检测打印件的力学性能,拉伸试样示意图如图1所示㊂采用光学显微镜㊁扫描电镜S E M (型号E M -30A X )进行微观组织形貌和断口形貌分析㊂2 结果和讨论2.1 增材制备工艺参数对空隙率的影响打印工艺参数及样品的密度如表2所示㊂由表2可以看到,打印工艺参数扫描速度(v )㊁激光功率(P )㊁层厚(t )和扫描间距(h )变化对密度均有影响㊂提高激光功率和适当降低速度均可以提高打印件的致密性,但考虑到打印效率问题,应尽快提高打印速度㊂表2 打印工艺参数及样品的密度序号激光功率/W 扫描速度/(mm /s )扫描间距/μm 层厚/μm 密度/%1300900110509623001025110509533309001105098序号激光功率/W 扫描速度/(mm /s )扫描间距/μm 层厚/μm 密度/%43301025110509753509001105099635010251105098探究不同工艺参数对打印件的影响,430不锈钢打印件的气孔分布形貌如图2所示㊂图2a 对应激光功率350W 和扫描速度900mm /s 的工艺(见表2中的第5组工艺),由图2a 可知,气孔较少,这也是该工艺致密度高的另一种表现㊂图2b 对应激光功率350W 和扫描速度1025mm /s 的工艺(见表2中的第6组工艺),图2b 的气孔率比图2a 有所增加㊂图2c 对应激光功率330W 和扫描速度900mm /s 的工艺(见表2中的第3组工艺),图2c 的气孔率与图2b 相当㊂图2d 对应激光功率330W 和扫描速度1025mm /s 的工艺(见表2中的第4组工艺),与图2c 相比,气孔率又进一步增加㊂图2e 对应激光功率300W 和扫描速度900mm /s 的工艺(见表2中的第1组工艺),气孔率也有所增加㊂图2f 对应激光功率300W 和扫描速度1025mm /s 的工艺(见表2中的第2组工艺),气孔率达到最高㊂由图2可知,随着激光功率增加和扫描速度降低,气孔率呈逐步降低的趋势㊂除了打印工艺之外,气孔的形成还受到粉末表面的水分及粉末中残留气体的影响[13],因此气孔率也不会与致密度成正比增加㊂较高的激光功率或者较低的激光扫描速度致使了较高的熔池温度,利于粉末完全熔化获得较为致密的打印件㊂此外,低熔点合金元素的蒸发也可以形成气孔,熔池的高凝固速率导致低熔点元素的蒸发无法有足够的时间从熔池逸出,因此只能存留在熔池内形成气孔[14],气孔在制件中随机分布,很难完全消除㊂综上,430不锈钢的打印工艺中,表2中的第5组工艺参数打印出来的样块,气孔最少且密度最高,因此,提高功率,降低扫描速度,打印件成形更好㊂2.2 打印态和固溶处理态的组织分析430不锈钢增材件底面(即激光扫描面)打印态和固溶处理态的光学显微组织如图3所示㊂由图3a 可见,430不锈钢激光熔融打印态的组织较细小,激光熔融的道与道之间的熔化线清晰可见;由图3b 可见,其固溶处理态的组织经过固溶处理,光学显微组织尺寸㊁形貌等没有明显变化㊂430不锈钢增材件扫描面打印态及固溶处理态的扫描电镜形貌图如图4所示㊂打印态在晶界处析出相较多,在图中显示为白色相分布,由图4a ㊁b 可见,这些晶界处的白色相是富铬的脆性相[15],会导致不锈㊃9㊃第4期鲍正浩,等:430不锈钢激光选区增材件的组织及抗拉性能㊃01㊃安 徽 工 程 大 学 学 报第38卷钢件的塑性降低㊂由图4c㊁d可见,经过固溶处理,430不锈钢增材件的晶界析出已经融入固溶体中,晶界的析出相几乎看不到㊂此外,对比图4b㊁d可知,相组织的尺寸也变大,说明固溶处理使得430打印件的相相应增加且晶界处的析出相融入固溶体中㊂图2 430不锈钢打印件的气孔分布形貌图3 430不锈钢增材件底面的光学显微组织430不锈钢增材件侧面的打印态及固溶处理态的形貌图如图5所示㊂由图5a㊁b可见,增材件侧面的打印态组织为鱼鳞状微观组织,图5a上的鱼鳞状结构特点说明大量金属颗粒经过激光轰击熔化后,处在短暂熔融状态的金属小液滴在重力作用下有向下滴落的运动趋势,其迅速凝固后,保留了这一特点,所以每一熔积层结构的下端会有近圆形的形貌呈现出来,凝固后的试样组织存在明显的各向异性,形成了类似鱼鳞状的组织形态,这是典型激光选区熔融打印态的组织㊂经固溶时效热处理可以看出,鱼鳞形状组织的边界形貌有所改变,鱼鳞状变得不明显,这是固溶处理过程中,组织重结晶导致组织形貌发生变化形成的㊂2.3 增材制造件的拉伸性能430不锈钢的S L M打印件的拉伸曲线如图6所示,图6中虚线曲线为打印态的拉伸曲线㊂由图6可知,打印态的抗拉强度达到700M p a,没有出现明显的屈服现象㊂图6中实线曲线为固溶处理态的拉伸曲线,固溶态具有明显的屈服现象;固溶态的抗拉强度比打印态有所降低,抗拉强度大约550M p a,固溶处理态比打印态具有更大的延伸率,表明固溶态具有更优良的塑性㊂打印态的强度高是由于S L M成形过程中形成的具有较高的位错密度的细小胞状组织结构使得打印件的强度指标提高[16-18];此外,打印态的晶界析出相也会提高不锈钢的拉伸强度,降低塑性[19];固溶处理有效消除了高位错密度,因此固溶态比打印态的抗拉强度有所降低㊂打印态拉伸断口形貌如图7所示㊂由图7a可见,打印态断口的低倍数的宏观断口形貌呈灰暗,剪切断裂的斜度较大,断裂的宏观表面有部分垂直于最大正应力方向,也有平行于切应力方向的,故正断和切图4 430不锈钢增材件底面扫描电镜形貌图图5 430不锈钢增材件侧面组织形貌图 图6 增材制造件的拉伸曲线断都有;高倍断口形貌呈暗灰色纤维状,由图7b ㊁c 可知,高倍断口形貌具有韧性断裂特征,故可认定为韧性断裂㊂增材件固溶处理态的断口形貌如图8所示㊂由图8a 断口宏观形貌可知,断口较为平齐,没有剪切断裂特征;图8b ~d 为断口的高倍扫描形貌图,断口表面有许多小韧窝,是典型的韧性断裂形貌㊂可见与打印态相比较,固溶处理态降低了样品的抗拉强度,提高了塑性㊂3 结论430不锈钢具有良好激光选区增材制造成型性能,打印工艺参数直接影响打印件的致密度和气孔率,当打印㊃11㊃第4期鲍正浩,等:430不锈钢激光选区增材件的组织及抗拉性能㊃21㊃安 徽 工 程 大 学 学 报第38卷图7 430不锈钢增材制造件的打印态断口形貌图8 430增材制造件的固溶热处理态的断口形貌的激光功率为350W,扫描速度为900mm/s,扫描间距为110mm和层厚为50mm时,打印件获得较高的致密度及较低的气孔率㊂打印态的晶界析出物明显,固溶处理消除了晶界析出;打印态的抗拉强度高达700M p a,固溶态的抗拉强度降低至550M p a,但具有更好的塑性;打印态样品的断裂受到剪切的影响较大,固溶处理态为典型韧性断裂㊂参考文献:[1] 熊华平,郭绍庆,刘伟,等.航空金属材料增材制造技术[M].北京:航空工业出版社,2019.[2] K R U T HJP,F R O Y E N L,V A E R E N B E R G H JV,e t a l.S e l e c t i v e l a s e rm e l t i n g o f i r o n-b a s e d p o w d e r[J].J o u r n a l o fm a t e r i a l s p r o c e s s i n g t e c h n o l o g y,2004,149(1):616-622.[3] G U DD,Z HA N G H M,C H E N H Y,e t a l.L a s e r a d d i t i v em a n u f a c t u r i n g o f h i g h-p e r f o r m a n c em e t a l l i c a e r o s p a c e c o m-p o n e n t s[J].C h i n e s e j o u r n a l o f l a s e r s,2020,47(5):0500002.[4] Z HA N GJL,L IFL,Z HA N G HJ.R e s e a r c h p r o g r e s so n p r e p a r a t i o no fm e t a l l i cm a t e r a l sb y s e l e c t i v e 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t r u c t u r e /p r o p e r t y (c o n s t i t u t i v ea n ds p a l l a t i o nr e s p o n s e )o fa d d i t i v e l ym a n u f a c t u r e d316Ls t a i n l e s s s t e e l [J ].A c t am a t e r i a l i a ,2017,138:140-149.[17]C A S A T IR ,L E MK EJ ,V E D A N IM.M i c r o s t r u c t u r ea n df r a c t u r eb e h a v i o ro f 316La u s t e n i t i cs t a i n l e s ss t e e l p r o d u c e db y s e l ec t i v e l a s e rm e l t i n g [J ].J o u r n a l o fm a t e r i a l s s c i e n c e a nd te c h n o l o g y ,2016,32(8):738-744.[18]S A L M A N O O ,G A MM E RC ,C H A U B E YAK ,e t a l .Ef f e c t o f h e a t t r e a t m e n t o nm i c r o s t r u c t u r e a n dm e c h a n i c a l p r o pe r t i e s of 316Ls t e e l s y n t h e s i z e db y s e l e c t i v e l a s e rm e l t i ng [J ].M a t e r i a l s s c i e n c e&e n g i n e e r i n g A ,2019,748:205-212.[19]R O YS ,S I L WA LB ,N Y C ZA ,e t a l .I n v e s t i g a t i n g th e e f f e c t o f di f f e r e n t s h i e l d i n gga sm i x t u r e so n m i c r o s t r u c t u r e a n d m e c h a n i c a l p r o p e r t i e s o f 410s t a i n l e s s s t e e l f ab r ic a t e dv i a l a r g e s c a l e ad d i t i v em a n u f a c t u r i n g [J ].A d d i t i v em a n u f a c t u r -i n g ,2021,38:101821.M i c r o s t r u c t u r e a n dT e n s i l eP r o p e r t i e s o f 430S t a i n l e s s S t e e l F a b r i c a t e db y L a s e r S e l e c t e dA d d i t i v eM a n u f a c t u r i n gB A OZ h e n g h a o ,HU A N GZ h o n g J i a *,Z H A N GC h e n y a n g ,T A OJ i n g(A n h u iK e y L a b o r a t o r y o fH i g h -P e r f o r m a n c eN o n -f e r r o u sM e t a lM a t e r i a l s ,W u h u241000,C h i n a )A b s t r a c t :I no r d e r t o i n v e s t i g a t e t h e t e n s i l e p r o p e r t i e so f 430s t a i n l e s s s t e e l p a r t s f a b r i c a t e db y a d d i t i v e m a n u f a c t u r i n g i n p r i n t e d a n d s o l i d s o l u t i o n s t a t e ,t h e 430s t a i n l e s s s t e e l s a m p l e sw e r e p r e p a r e db y l a s e r s e l e c t i v e a d d i t i v em a n u f a c t u r i n g ,t h e n t h em i c r o s t r u c t u r e a n d t e n s i l e p r o p e r t i e s o f a s p r i n t e d a n d a s s o l u -t i o n t r e a t e dw e r e i n v e s t i g a t e d .T h e r e s u l t s s h o wt h a t 430s t a i n l e s s s t e e l h a s g o o d f o r m a b i l i t y o f l a s e r s e -l e c t i v e a d d i t i v em a n u f a c t u r i n g .T h e r ea r em a n yg r a i nb o u n d a r yp r e c i p i t a t e s i nt h e p r i n t e ds a m p l e s ,a n d g r a i nb o u n d a r yp r e c i p i t a t e s h a v e b e e n e l i m i n a t e db y s o l i d s o l u t i o n t r e a t m e n t .T h e t e n s i l e s t r e n g t ho f t h e p r i n t e d s a m p l e s i s u p t o 700M P a ,a n d t h a t o f t h e s o l i d s o l u t i o n s a m p l e s i s r e d u c e d t o 550M P a .H o w e v -e r ,t h e p l a s t i c i t y o f s o l u t i o nt r e a t e ds a m p l e sw e r eb e t t e r t h a nt h a t o f t h e p r i n t e ds a m p l e s .T h e f r a c t u r e m o r p h o l o g y s h o w s t h a t t h e t e n s i l e f r a c t u r eo f p r i n t e ds a m p l e i s g r e a t l y a f f e c t e db y t h e s h e a r f o r c e ,a n d t h e s o l u t i o n t r e a t e d s a m p l eh a s a t y p i c a l d u c t i l e f r a c t u r e .K e y w o r d s :l a s e r s e l e c t i v e a d d i t i v em a n u f a c t u r i n g ;s t a i n l e s s s t e e l ;m i c r o s t r u c t u r e ;t e n s i l e p r o p e r t y ㊃31㊃第4期鲍正浩,等:430不锈钢激光选区增材件的组织及抗拉性能。

Ž.Materials and Design21200027᎐30Effects of temperature and loading rate on fracturetoughness of structural steelsCun-Jian Li UWelding Research Department,Central Iron&Steel Research Institute,No.76Xueyuan Nanlu,Haidian District,Beijing100081,ChinaReceived19April1999;accepted1June1999AbstractŽ.The fracture characteristics of body-centered-cubic BBC metallic materials is highly dependent on temperature and strain rate.To study the effects of temperature and loading rate on the dynamic fracture behavior of a high-strength low-alloy steel,a series of fracture toughness experiments were carried out using three-point-bend specimens and impact specimens.The results show that with the increase of temperature and decrease of loading rate,dynamic fracture toughness rises.Furthermore,the study indicates that dynamic fracture toughness is controlled by dislocation thermal activation motion and may be quantitatively expressed by introducing the Arrhenius equation.ᮊ1999Elsevier Science Ltd.All rights reserved.Keywords:Ferrous metals and alloys;Fracture toughness;Fracture behavior1.IntroductionIt has been well known that among the factors af-fecting material fracture characteristics,temperature and strain rate are two very important external ones. Most brittle fracture accidents with low stress occurr in steel structures due to low temperature and dynamic loading conditions.The mechanism of plastic deforma-tion and fracture of materials may change at differentw xtemperatures and strain rates1.At present it has been widely investigated that the dynamic yield strengthw x of steels varies with temperature and strain rate2. However,the research of dynamic fracture toughness of steels is mostly concentrated on the study of the influence of temperature under some loading condi-tion.Therefore,it is essential to systematically study the combined effects of temperature and load rate on dynamic fracture toughness of steels in ensuring the safety application of their structural parts.U Tel.:q86-10-6218-2783;fax:q86-10-6218-2581.2.Materials and experimental procedureŽA785-MPa grade steel designated as HQ785C in.Chinese steel standards plate of20-mm thickness was employed for this investigation with its chemical com-position being shown in Table1.The steel plate was received as quenched and tempered with its mi-crostructure being tempered martensite.Three-point-bend specimens with dimensions130=28=14mm and instrumented impact specimens with dimensions 55=10=10mm were cut from the steel plate,respec-tively.The preparation of the specimens is according to ASTM E399-91and ASTM E636-83,respectively. Fracture toughness evaluating tests were carried out at a series of temperatures ranging from77to333K on a MTS-880materials testing machine and on a TINIUS-T84instrumented impact machine,respec-tively.The load-point displacement rates of the three-point-bend test were,respectively,0.01,0.7and50 mm r s;the instrumented impact rate was5540mm r s.˙When the load-point displacement rates,⌬,were0.01 mm r s and0.7mm r s,an X᎐Y coordinate recorder was0261-3069r00r$-see front matterᮊ1999Elsevier Science Ltd.All rights reserved.Ž.PII:S0261-30699900042-4()C.Li r Materials and Design 21200027᎐3028Table 1Ž.Chemical composition of the investigated steel wt.%C Si Mn P S Cr Mo V Cu B 0.120.291.050.0170.0130.920.380.050.240.0015Žapplied to record the P y ⌬load ᎐load point displace-.ment curve.When the rates were 50mm r s and 5540mm r s,a computerized data acquisition processing sys-tem was used.During the test,ethyl alcohol and liquid nitrogen were used to adjust the required temperatures below ambient for the experiment and an autoheating device was used to control the temperature over am-bient.K values in the lower-shelf region were de-1d termined according to the method of ASTM E 399-91;in the ductile-to-brittle transition and upper-shelf re-gions,J ,the critical J -integral values were measured 1d w x based on the following relation given by Rice et al.3:2U Ž.J s11d Ž.B W y a where B is specimen thickness,W is specimen width,a is crack depth,U is the area under the load ᎐load point Ž.displacement P y ⌬curve until the crack extension point.In instrumented impact tests,J values were mea-1d sured over all temperature ranges according to the method of ASTM E 636-83,J was then converted 1d into K by the following equation:1d 1r 22Ž.Ž.K s EJ r 1y ␯21d 1d where E s Young’s modulus and ␯s Poisson’s ratio.3.Results and discussionFig.1shows the experimental results of dynamic fracture toughness for HQ785C steel at different tem-peratures and loading rates.From Fig.1,it can be shown that with the decrease of temperature or in-crease of loading rate,K values usually reduce;when 1d temperature is lowered to a certain critical tempera-˙ture,approximately 77K for ⌬s 50mm r s and 5540mm r s,K values tend to hold a constant value at the 1d lower-shelf region.This value is not related to temper-ature and loading rate and can be defined as the athermal activation component K ,which is 38.61dn MPa m 1r 2by measurement from Fig.1.In the transi-tion region,the part relative to temperature and load-ing rate in K values can be defined as the thermal 1d activation component,K .Then the total dynamic 1da fracture toughness K can be expressed as:1d Ž.K s K q K 31d 1dn 1da˙when temperatures are over 253K for ⌬s 0.01mm r s ˙and 293K for ⌬s 0.7mm r s,K values have little to 1d do with temperature,i.e.reach the upper-shelf region.In the transition region,according to the Arrhenius w x equation 4:Ž.Ž.␧s ␧exp y ⌬G r␬T 4˙˙0f where ⌬G is the activation energy,␧is a frequency˙f 0Ž.Fig.1.Results of experimental K values the values calculated are shown in the dotted lines .1d()C.Li r Materials and Design 21200027᎐3029˙y 1Fig.2.Relation between ⌬and T at the different K values.1d a Žy 5factor,␬is Boltzmann’s constant ␬s 8.6112=10.ev r K ,T is Kelvin’s temperature,and ␧is the strain˙˙rate relative to l ,r ,t and ⌬,and can be calculated by w x the following formula 5:˙r и⌬Ž.␧s5˙t иlwhere r is the radius of slip band with the specimen Žbeing bent,t is the slip and width,l is the span the˙.distance between specimen supports ,and ⌬is the load-point displacement rate.Ž.Ž.Substituting Eq.5into Eq.4,one can obtaint иl˙Ž.Ž.Ž.⌬s ␧exp y ⌬G r ␬T s X exp y ⌬G r ␬T 6˙0f f r t иlwhere X s ␧.˙0r ˙Ž.Eq.6can be plotted in ln ⌬y 1r ␶coordinate by Žthe data from Fig.1at constant K levels see Fig.1da .2.With the same K values at the logarithmic1da ˙y 1coordinate,the relation between ⌬and T is basically a linear relation.Furthermore,the lines are approxi-mately interparallel at different K levels.Measuring 1da the slope in Fig.2,one can obtain the activation energy ⌬G s 1.02ev.As ln X is the intercept of the lines,Fig.f 2also indicates that ln X decreases with the increase of K .Hence we presume that there is an index rela-1da tionship between K and X for example:1da mŽ.Ž.X s X K r K 701da 1do where X and m are two material parameters irrele-0vant to fracture toughness,K is a unit of fracture 1do toughness.Fig.3is a plot of X and K by data from 1da Fig.2.At double the logarithmic coordinate,as a good linear relation between X and K is shown,the1da Ž.correctness of the assumption of Eq.7is also verified.Measuring the slope and intercept of the line in Fig.3gives m sy 12.7,X s 4.17=1048mm r s.From Eqs.0Ž.Ž.Ž.7,6and 3,the total dynamic fracture toughness K can be expressed as:1d 1r m⌬G ˙⌬fŽ.K s K q K exp81d 1dn 1dož/X m ␬TŽ.K values calculated by Eq.8at different loading 1d rates and temperatures are indicated with the dotted lines in Fig.1.They are in good agreement with the Ž.experimental results.Hence Eq.8quantitatively de-scribes the relationship between the dynamic fracture toughness and temperature and loading rate for HQ785C steel.This also indicates that dynamic frac-ture toughness is controlled by thermal activation mo-tion of dislocationsat intermediate temperatures and strain rates.Fig.3.Relation between X and K .1d a()C.Li r Materials and Design21200027᎐30 30When temperature reduces to a certain critical value under some loading rate,the thermal activation process of material dislocation motion is restrained,so that it is difficult for dislocations to move away form crack tip. Thus cracks can only extend along the direction of the maximum principal stress on the plane with the maxi-mum strain energy density,resulting in brittle fracture. With the increase of temperature,the crystal lattice resistance of dislocation motion,i.e.P-N force and the resistance of point imperfection against dislocation source decrease,so that plastic deformation easily oc-cur,but dislocations still need thermal activation en-w xergy to overcome the barrier6.The higher the load-ing rate,the shorter the time of dislocations over-coming the barrier by thermal activation becomes. Therefore,dynamic fracture toughness rises with tem-perature increasing and loading rate decreasing.When temperature increases to a certain critical value at some loading rate,i.e.the so-called upper-shelf region, the probability of dislocation being activated increases, and dislocations have sufficient time to complete the thermal activation process,so that dynamic fracture toughness becomes insensitive to temperature and loading rate.4.ConclusionThe range of K values varying with temperatures1dand strain rates can be divided into three regions,i.e. upper-shelf region,thermal activation region and athermal activation region.In the thermal activation region,the dynamic fracture behavior of metallic mate-rials is controlled by thermal activation motion of dislo-cations.With the increase of temperature or decrease of loading rate dynamic fracture toughness rises and may be quantitatively expressed by introducing the Arrhenius equation.AcknowledgementsThe author would like to thank De-Gen Wang and Qiao-Tian Wu of the Department of Physical Metal-lurgy at Central Iron&Steel Research Institute for their assistance in doing tests.Thefinancial support of the China Natural Science Foundation is also gratefully acknowledgedReferencesw x1Radone JG,Ruben DJ.The effect of temperature and strain rate on yield strength of high strength steels.High TemperatureŽ.Technol1994;152:216᎐228.w x2Weisheng L,Mei Y.Quantitative description of temperature and strain rate dependence of yield strength of structural steels.Ž.Eng.Fracture Mechanics1996;534:633᎐643.w x3Rice JR,Paris PC,Merkle JG.Some further results of J-integral analysis and estimates.Progress in Flaw Growth and Fracture Toughness Testing.ASTM STP536,1973:231᎐245.w x4Bassani JL.Macro and micro mechanical aspects of creep frac-ture.In:Wang SS,Penten WJ,editors.Advances in aerospace structures and materials.New York:ASME,1981.w x5Chi C,Qigong C,Renzhi W.Engineering fracture mechanics.Beijing:National Defense Industry Press,1977:275᎐276.w x6Xianhe G.Effect of solid solution softening on dynamic me-chanical behaviors of structural steels.PhD.thesis,Tsinghua University,1991.。