Microstructure Development and Mechanical
Microstructural Development during Hot Working of Mg-3Al-1Zn
Microstructural Development during Hot Working of Mg-3Al-1ZnA.G.BEER and M.R.BARNETTThe microstructural evolution is examined during the hot compression of magnesium alloyAZ31for both wrought and as-cast initial microstructures.The influences of strain,tempera-ture,and strain rate on the dynamically recrystallized microstructures are assessed.Both thepercentage dynamic recrysallization(DRX)and the dynamically recrystallized grain size werefound to be sensitive to the initial microstructure and the applied deformation conditions.Lower Z conditions(lower strain rates and higher temperatures)yield larger dynamically re-crystallized grain sizes and increased percentages of DRX,as expected.The rate with which thepercentage DRX increases for the as-cast material is considerably lower than for the wroughtmaterial.Also,in the as-cast samples,the percentage DRX does not continue to increase towardcomplete DRX with decreasing Z.These observations may be attributed to the deformationbecoming localized in the DRX fraction of the material.Also,the dynamically recrystallizedgrain size is generally larger in as-cast material than in wrought material,which may beattributed to DRX related to twins and the inhomogeneity of deformation.Orientation maps ofthe as-cast material(from electron backscattering diffraction(EBSD)data)reveal evidence ofdiscontinuous DRX(DDRX)and DRX related to twins as predominant mechanisms,withsome manifestation of continuous DRX(CDRX)and particle-stimulated nucleation(PSN).DOI:10.1007/s11661-007-9207-5ÓThe Minerals,Metals&Materials Society and ASM International2007I.INTRODUCTIONR EDUCING the weight of vehicles,for increased fuel efficiency,is a high priority for the automotive industry. Due to its low density,magnesium is a potential material for a range of automotive components.Most of the consumption of magnesium alloys by the automotive industry has been in the form of die castings.[1]However, wrought magnesium products have the advantage over castings of higher strength and ductility,and so are more suited to structural applications.Wrought magnesium alloys can be deformed at elevated temperatures using primary fabrication meth-ods such as rolling,extrusion,and forging,but as the workability of the material is limited,production rates are slow and hence thefinal product is comparatively expensive.Deforming magnesium and its alloys at elevated temperatures is of great metallurgical impor-tance because not only is the workability improved but also thefinal grain size and,to a great extent,thefinal properties of the material are altered.The operation of dynamic recrystallization(DRX)during hot deforma-tion of magnesium is of particular importance,because it reduces theflow stress during deformation and controls thefinal grain size.[2]Therefore,to be able to successfully control the microstructural evolution dur-ing bulk working operations,an understanding of hot deformation behavior of magnesium,particularly the operation of DRX,is essential.Magnesium has a high stacking-fault energy and it might thus be expected to dynamically soften by dynamic recovery(DRV)instead of DRX.However, the early work by Humphreys and co-workers[2,3] showed that DRX was in fact an important mechanism during the high-temperature deformation of magne-sium.The presence of DRX was attributed to the constraints imposed by the lack of easily activated slip systems of magnesium,rather than its stacking-fault energy.The operation of DRX in magnesium may also be linked to its high grain boundary diffusion rate.[4] The literature presents a complex picture of the mechanisms by which DRX operates in magnesium.It appears that different types of DRX take place under different deformation conditions,e.g.,References5,6, and7.Twinning also seems to play a significant role in the nucleation of dynamically recrystallized grains,e.g., References5,8,and9.The mechanism of conventional discontinuous DRX (DDRX)has been identified in magnesium alloys by a number of workers.[3,5–7,10–12]This mechanism involves the development of high-angle grain boundaries via the nucleation and growth of new grains.This typically initiates at high-angle boundaries:original grain bound-aries,the boundaries of dynamically recrystallized grains,or boundaries created during straining.[13]The bulging of grain boundaries is frequently observed prior to DRX,and it is assumed that a mechanism closely related to strain-induced boundary migration,which is observed during annealing of cold-worked metals, operates.[13,14]The continuous mechanism of DRX(CDRX)has been identified in aluminum alloys(e.g.,References15 and16)and also in some magnesium alloys.[2–7,11,12,17–19]A.G.BEER,Research Academic,and M.R.BARNETT,QEII Research Fellow,are with the CRC for Cast Metals Manufacturing (CAST),School of Engineering and Technology,Deakin University, Geelong VIC3217,Australia.Contact e-mail:aiden.beer@.au Manuscript submitted:June30,2006.Article published online July13,2007.This mechanism differs from that of conventional DRX in that the development of high-angle grain boundaries during hot deformation does not involve the nucleation and growth of grains at pre-existing boundaries.Due to the high efficiency of DRV,new grains are formed progressively within the deformed original grains from the continuous increase of subgrain boundary misorien-tations.[20]The influence of twinning on the dynamically recrystallized behavior of magnesium is important in coarse-grained alloys.Many studies have reported the nucleation of dynamically recrystallized grains at twins,e.g.,References2,5,9,11,17,and21through 23.The influence of twinning on DRX has been shown to be of importance for other hcp metals as well,particularly titanium.[24,25]Sitdikov et al.[8,12] observed DRX associated with twinning in pure magnesium and suggested that the nucleation of ‘‘twin’’DRX occurred by the intersection of primary twins(whereby a crystallite is formed,bordered by twin boundaries),by double twinning,and by the subdivision of primary twins by transverse low-angle boundaries(which are transformed to random high-angle boundaries upon further straining).Myshlyaev et al.[26]also observed the nucleation of dynamically recrystallized grains where twins intersected with grain boundaries.For enhanced mechanical properties,it is often desirable to obtain thefinest grain size possible.Know-ing the influence that the deformation conditions have on the dynamically recrystallized grain size is important for achievingfine grain sizes.When DRX(continuous or discontinuous)is operating,the size of the new grains is characteristic of the deformation conditions and is generally independent of the initial grain size.In magnesium,however,it has been observed that a smaller initial grain size yields a smaller dynamically recrystallized grain size,[27]although the reason for this phenomenon remains unclear.Several studies have examined the dynamically recrystallized grain size developed during the hot deformation of magnesium and developed equations that relate it to the deforma-tion conditions(either stress or Z).[5–7,12,27–31]While some researchers have shown that the relationship between the dynamically recrystallized grain size and the deformation conditions can be described using a power law(e.g.,References27and29through31), others have found that this relationship was not linear and attributed this to a change in the rate controlling dynamically recrystallized mechanism,i.e.,DDRX operative at high temperatures and CDRX operative at lower temperatures.[5–7,12]Understanding the degree to which DRX goes to completion is also important,because a homogenous microstructure is desired for optimalfinal properties. The percentage DRX,at constant and increasing strains, has received a small degree of attention(e.g.,References 4through6,12,and30through32),but comparisons between different initial starting conditions is lacking. Many studies have used magnesium alloys in the‘‘as-cast’’condition,but the deformation behavior of wrought magnesium is also of importance.The defor-mation of wrought,or preworked,material occurs in primary fabrication processes such as multistep forging and rolling operations and in the extrusion of billets that have been scalped and pre-extruded.The initial grain size will have a marked effect on recrystallization kinetics,with afine-grained material recrystallizing more rapidly than a coarse-grained material.[13]The higher density of grain boundaries in afiner-grained structure means that there is an increase in the nucle-ation sites for DDRX.On the other hand,the reduction in grain size reduces the number of inhomogeneities such as shear and deformation bands,which also act as sites for nucleation.Twinning is also suppressed at smaller grain sizes,and thus,one would expect a smaller contribution of‘‘twin’’dynamically recrystallized mech-anisms to the overall microstructural development of fine-grained magnesium alloys.To examine these issues further,the present article studies the effects of strain,deformation conditions,and initial state(wrought or as-cast)on the microstructures developed after the hot deformation of magnesium alloy AZ31.II.EXPERIMENTAL PROCEDUREThe material used in the current study was commer-cial grade magnesium alloy AZ31(Mg-3pct Al-1pct Zn-0.2pct Mn)in both the‘‘wrought’’and‘‘as-cast’’form.The wrought material was in the form of18-mm-diameter extruded rod,and the machined compression samples were orientated such that the compression direction coincided with the extrusion direction.The initial wrought microstructure has a strong texture,in which the basal planes were aligned parallel to the extrusion direction,[33]and an average grain size of22.5 l m.The as-cast material was received in the form of a direct chill(DC)cast billet that was130mm in diameter. Machined compression samples were subsequently homogenized for24hours at425°C and had an average grain size of316l m.Samples were covered with Polytetraflouroethylene(PTFE)tape and held at the test temperature for5minutes prior to testing.To examine the evolution of microstructure during hot deformation,uniaxial compression testing was conducted at a temperature of350°C and at a constant strain rate of0.01s–1.The samples were quenched in water following equivalent strains of0.2,0.4,0.6,0.8, and 1.0.Also,to examine thefinal microstructures developed,uniaxial compression testing was conducted at temperatures ranging between300°C and450°C and at constant strain rates ranging between0.01s–1and 1s–1.These samples were quenched in water following an equivalent strain of1.0.Specimens for metallograph-ic examination,taken from the cross section of deformed samples,were cold mounted in epoxy resin,fine ground with1200grit SiC paper,and diamond polished through6and3l m.Specimens were chemi-cally polished for45seconds in10pct nital and then etched for5seconds with acetic picral.Samples for examination using electron backscattering diffraction (EBSD)were polished as for optical metallography,after which they were polished with a colloidal silica suspension and then etched with a solution of10mL HNO3,30mL acetic acid,40mL H2O,and120mL ethanol for5seconds.[34]III.RESULTSA.Evolution of Microstructure with StrainThe influence of strain on the microstructures that are developed when the wrought and as-cast materials are deformed at a temperature of350°C and a strain rate of 0.01s–1can be seen in Figures1and2,respectively.The initial microstructure of the extruded material displays fairly uniform grain boundaries.As the strain is increased to0.2,the original grain boundaries become serrated and dynamically recrystallized grains are observed at prior grain boundaries.As the strain is further increased to0.4and0.6,dynamically recrystal-lized grains become more prevalent.At strains of0.8, DRX has proceeded further and consumed most of the original grains.The microstructure is almost completely DRX after a strain of1.0.In line with the observations during the compression of the wrought material,the original grain boundaries of the as-cast material also become serrated and a small number of dynamically recrystallized grains are ob-served at prior grain boundaries after deformation to a strain of0.2.In the deformation of the as-cast material, deformation twinning is evident.Twinning does not occur in all grains and slight serrations can be seen on some of the twin boundaries.At a strain of0.4,further dynamically recrystallized grains nucleate at the original grain boundaries forming a‘‘necklace’’type structure. In some areas,these necklaces have thickened to be several dynamically recrystallized grains wide,an exam-ple of which is the diagonal band offine grains running from the lower left corner to the upper right corner of the micrograph(Figure2(c)).The boundaries oftwins, Fig1—Microstructural evolution of wrought AZ31with increasing strain,deformed in compression at temperature of350°C and a strain rate of0.01s–1:(a)initial microstructure,(b)strain of0.2,(c)strain of0.4,(d)strain of0.6,(e)strain of0.8,and(f)strain of1.0.Fig2—Microstructural evolution of as-cast AZ31with increasing strain,deformed in compression at a temperature of350°C and a strain rate of0.01s–1:(a)initial microstructure,(b)strain of0.2,(c)strain of0.4,(d)strain of0.6,(e)strain of0.8,and(f)strain of1.0.which had formed at earlier strains,have since becomefurther serrated and DRX related to twinning is evident.The dynamically recrystallized grains are generallyobserved to have nucleated on the twin boundaries ofthick twins and at twin intersections.It is also clear thatdynamically recrystallized grains develop within,andfill,narrow twins.This can be more clearly seen in themicrograph at a strain of 0.6(Figure 2(d)),where a largegrain has been subdivided by a thin lenticular twin,which has than acted as a nucleation site for DRX.Atthis strain,the necklaces of dynamically recrystallizedgrains decorating the pre-existing grains,and thedynamically recrystallized grains associated with twins,have continued to broaden.At a strain of 0.8,thedevelopment of DRX is more extensive,but even afterstrains of 1.0,the structure remains partially DRX.Differences between the development of DRX for thewrought and as-cast material are highlighted in Fig-ure 3,where the percentage DRX and the dynamicallyrecrystallized grain size are plotted as a function ofstrain.For the as-cast material,the rate with which thepercentage DRX increases is considerably lower thanthat observed for the wrought material.The microstruc-ture is close to 65pct DRX at the highest strainexamined,and it is not obvious what level of strain willbe required to develop a completely dynamically recrys-tallized microstructure.For both initial microstructures,the size of dynamically recrystallized grains (distin-guished from pre-existing grains by their size andmorphology)remains virtually unchanged as deforma-tion proceeds (Figure 3(b))This finding is in line withthe observation of Sah et al.,[35]who examined theconventional dynamically recrystallized behavior duringthe hot deformation of nickel.The average dynamicallyrecrystallized grain size is approximately 7l m in thewrought material and 9l m in the as-cast material.B.Microstructure at a Strain of 1.0The influence of strain rate and temperature on themicrostructure formed when the wrought and as-castmaterials are deformed at a strain of 1.0is shown inFigures 4and 5,respectively.For the wrought samples,the material is not completely DRX for the tests carried out at 300°C.Twins and dynamically recrystallized grains decorating twins can be seen in some of the remnants of pre-existing grains.When the temperature is increased,the percentage of DRX and the size of the dynamically recrystallized grains are increased.A sim-ilar observation,albeit to a lesser degree,is observed for a reduction in strain rate.For temperatures of 350°C and above,many of the microstructures developed when the wrought material is deformed to a strain of 1.0are virtually completely DRX.A significantly different microstructure from that of the wrought material is observed in the as-cast samples (Figure 5).At 300°C,the material is far from being completely DRX.The dynamically recrystallized grains have developed on pre-existing grain boundaries and also within twins,the latter of which is particularly noticeable in the microstructure developed at the highest strain rate (Figure 5(c)).As with the deformation of the wrought material,as the temperature of deformation is increased,the dynamically recrystallized grain size developed in the as-cast material increases.Also,the percentage of DRX can be seen to increase at higher deformation temperatures.At the highest temperatures examined,the material is still far from being completely DRX,although this is not obvious in the micrographs due to the high magnification chosen to sufficiently display the dynamically recrystallized grain size at the low temperatures.Differences between the development of DRX,when the wrought and as-cast material is deformed to a strain of 1.0,are highlighted in Figure 6,where the percentage DRX and the dynamically recrystallized grain size are plotted as a function of the Zener–Hollomon parameter,Z (¼_e exp Q =R T ðÞ),where Q was taken as the activa-tion energy for self-diffusion,135kJ/mol.[36,37,38]For the wrought material,it can be seen that the percentage dynamic recrystallization is lowest at the highest Z values (high strain rate and low temperature)and increases toward being completely DRX at lower values of Z .It is also evident that the material neverreaches Fig 3—Influence of initial microstructure on the development of (a )the percentage DRX and (b )the dynamically recrystallized grain size (plot-ted as a function of strain for AZ31deformed at a temperature of 350°C and a strain rate of 0.01s –1).Lines are drawn to guide the eye.100pct DRX at a strain of 1.0.For a given value of Z ,the percentage dynamic recrystallization is considerablylower in the as-cast material as compared to the wroughtmaterial.The percentage DRX is lowest at the highest Zvalues in the as-cast material;however,at lower valuesof Z ,the percentage DRX does not keep increasingtoward being completely DRX.Instead,the percentageDRX saturates at a value of approximately 65pctDRX.Figure 6(b)reveals that the dynamically recrystallizedgrain sizes developed,after deformation to a strain of1.0,are larger in the as-cast material,particularly at lowvalues of Z .For both the wrought and as-cast material,the dynamically recrystallized grain sizes decrease withincreasing Z according to an inverse power law of theformd DRX ¼AZ Àn ½1 where d DRX is the dynamically recrystallized grain size,A is a constant,Z is the Zener–Hollomon parameter,and n is the power-law exponent.Equations of this form have been used to predict the dynamically recrys-tallized grain size developed during hot working of magnesium alloys.[27,29–31]Although a more accurate fit of the current data could be obtained by applying a calculated apparent activation energy,it is convenient to adopt a standard value of the activation energy,Q ,for the hot working of magnesium alloys (it has also been shown that Q varies with temperature [21,39–42]).We have therefore opted for the activation energy for self-diffusion (135kJ/mol).[36,37,38]The values of A and n determined in the present work are 57.5and 0.095,respectively,for the wrought material,and 165.5and 0.126,respectively,for the as-cast material.IV.DISCUSSION Both continuous and discontinuous DRX have been observed to operate in magnesium alloys.[2–7,10–12,17,18]However,they are expected to yield differing responses with respect to the deformation behavior and micro-structural development.Our data will therefore be examined in light of these possible mechanisms.A.Flow Stress and Dynamically Recrystallized GrainSizesOne characteristic that can help distinguish betweenthe operation of DDRX and CDRX is the stress-straincurve.In DDRX,the dislocation density is lowered bythe migration of high-angle grain boundaries that‘‘sweep’’out the stored dislocations.This results in areduction of the work hardening rate andappreciableFig 4—Developed microstructures of wrought AZ31hot deformed to a strain of 1.0in compression:(a )300°C,0.01s –1;(b )300°C,0.1s –1;(c )300°C,1s –1;(d )350°C,0.01s –1;(e )350°C,0.1s –1;(f )350°C,1s –1;(g )400°C,0.01s –1;(h )400°C,0.1s –1;(i )400°C,1s –1;(j )450°C,0.01s –1;(k )450°C,0.1s –1;and (l )450°C,1s –1.flow softening,leading to a distinct peak in the strain-strain curve.By contrast,CDRX involves very littleboundary migration as subgrains rotate and develophigh-angle grain boundaries.The dislocation density isnot significantly reduced as it is in DDRX and thus thestress-strain curve generally attains a maximum stressand is expected to show very little softening upon further straining,all else constant.The flow curves obtained from the compression of the wrought and as-cast material,at a temperature of 350°C and a strain rate of 0.01s –1,are displayed in Figure 7.The flow curves exhibit a distinct peak inthe Fig 5—Developed microstructures of as-cast AZ31hot deformed to a strain of 1.0in compression:(a )300°C,0.01s –1;(b )300°C,0.1s –1;(c )300°C,1s –1;(d )350°C,0.01s –1;(e )350°C,0.1s –1;(f )350°C,1s –1;(g )400°C,0.01s –1;(h )400°C,0.1s –1;(i )400°C,1s –1;(j )450°C,0.01s –1;(k )450°C,0.1s –1;and (l )450°C,1s –1.Fig 6—Influence of initial microstructure on (a )the percentage DRX and (b )the dynamically recrystallized grain size,for wrought and as-cast AZ31hot deformed to a strain of 1.0in compression.stress-strain behavior,as do many magnesium hotworking flow curves presented in the literature (e.g.,References 2,3,7,22,32,41,and 43through 45),whichtends to support the operation of a significant level ofDDRX.A higher strain to the peak flow stress with acoarser grain size in the as-cast material is also consis-tent with conventional DRX.The size of the new grains developed when DRX isoperating (either continuous or discontinuous)is char-acteristic of the deformation conditions.For DDRX,this is due to the simultaneous deformation occurringwhere,as the nucleated grains grow,work hardeningoccurs in the dynamically recrystallized grains and thedriving force for further growth is reduced.For CDRX,subgrains developed during DRV rotate and develophigh-angle grain boundaries.Because very little bound-ary migration occurs,the recrystallized grains aregenerally found to be a little larger than the priorDRV subgrains,the size of which is found to be stronglydependent on the deformation stress.[28]For both DDRX and CDRX,increasing Z will resultin a reduced grain size and this relationship is upheld inthe present results (as seen in Figure 6(b)).However,thesensitivity of the grain size to Z is expected to differ between the two mechanisms.Derby [28]compared theDDRX grain size for a wide selection of metals andminerals,including magnesium,and found that it isrelated to the deformation stress in a uniform manner.Asimilar relationship,albeit with a different sensitivity tochanges in flow stress,was found for the DRV subgrainsize (which can be expected to be slightly smaller thanthe CDRX grain size).The dynamically recrystallizedand DRV regions presented by Derby [28]are plotted inFigure 8with the current grain size data.In this plot,the mean steady-state grain size during DRX,normalized by a Burger’s vector of 3.21·10–10(m),[28]is plotted against the steady-state deformation stress (taken at a strain of 1.0),normalized by the shear modulus (equal to 16;6001þÀ0:49T À300ðÞ=924ðÞðÞ½ (MPa)[36]).It can be seen that the dynamically recrystallized grain size is more sensitive to changes in stress than the DRV subgrain size.However,a large degree of overlap between the DRX and DRV regions exists,and the present data fit within both regions.That is,no clear distinction between the operation of DDRX and CDRX can be made from the present grain size data.Figure 6(b)also reveals that the log-log plot of the dynamically recrystallized grain size against Z is approximately linear for both initial microstructures.Similar relationships developed by Kaibyshev and co-workers for pure magnesium and magnesium alloy ZK60found that the dynamically recrystallized grain size correlates with Z (or the normalized stress,r /G ),but the power-law relationship is only obeyed for specified temperature ranges.[5–7,12]This was attributed to changes in the rate-controlling dynamically recrystal-lized mechanism,i.e.,DDRX operative at high temper-atures and CDRX operative at lower temperatures.The current results,however,do not show such a sharp transition,and so it is unlikely that a dramatic change in the dominant dynamically recrystallized mechanism has occurred in the present tests.A gradual change in the dynamically recrystallized mechanism may,however,still result in a linear log-log plot of the dynamically recrystallized grain size against Z.Fig 7—Flow curves for the deformation of AZ31at a temperatureof 350°C and a strain rate of 0.01s –1.Fig 8—Mean steady-state grain size,normalized by the Burger’s vec-tor,plotted against the deformation stress (at a strain of 1.0),nor-malized by the shear modulus.The shaded regions correspond to the grain size (or subgrain size in the case of DRV)models developed by Derby [26].B.EBSD and Nucleation MechanismsIn an attempt to gain more insight into the mecha-nisms by which DRX operates in the present material,orientation maps were generated from EBSD data(Figures 9through 11).This was conducted on as-castsamples that were compressed to various strains at atemperature of 350°C and a strain rate of 0.01s –1.As-cast samples were chosen in preference to wroughtsamples as a clearer distinction between pre-existing anddynamically recrystallized grains can be made.In allorientation maps presented (apart from Figure 11(b)),thick black lines correspond to high-angle boundaries(misorientations greater than 15deg),while low-angleboundaries are represented by thin white lines (misori-entations greater than 2deg)and thin black lines(misorientations greater than 5deg).Orientation maps of as-cast AZ31,deformed to astrain of 0.2,reveal evidence of the operation of DDRX(Figures 9(a)and (b)).The bulging of original grainboundaries is frequently observed,as can be seen at thehigh-angle boundary between parent grains A and B inboth orientation maps.An example of where a bridginglow-angle boundary has developed behind a bulgedsection of grain A is indicated by C in Figure 9(a).In thecase of Figure 9(b),a bulged section of grain A hasalmost been completely ‘‘pinched off’’(indicated by D).Interestingly,Figure 9(b)reveals that new grains have developed in regions associated with particles.The E indicates a typical example (the dark regions in the bottom right-hand corner of Figure 9(b),where no indexing of the EBSD patterns occurred,are interme-tallic particles).Particle-stimulated nucleation (PSN)has been previously reported for several magnesium alloys.[46,47,48]However,in the case of the present alloy,this mechanism is rarely observed and is not believed to significantly contribute to the development of DRX.The CDRX is typically identified by an increase in misorientation from the center to the edge of pre-existing grains and subgrain development near the boundary.[13]This was not observed in orientation maps of the sample that had been deformed to a strain of 0.2.However,when the material was deformed to a strain of 0.6,microstructural features consistent with the opera-tion of CDRX could be observed.Figure 10(a)shows an orientation map of a large pre-existing grain bordered by a necklace of dynamically recrystallized grains.A plot of the cumulative misorientation along the dotted line from the interior of the grain (A)to the grain boundary reveals a large increase in misorientation (Figure 10(b)).The development of low-angle bound-aries immediately adjacent to the necklace can also be seen,as indicated by B and C,and it could be envisagedthat,upon further deformation,these mayprogressivelyFig 9—(a )and (b )Orientation maps revealing evidence of DDRX (grain boundary bulging and the development of bridging subgrain bound-aries)in as-cast AZ31,deformed in compression to a strain of 0.2at a temperature of 350°C and a strain rate of 0.01s –1(thin white line >2deg misorientation,thin black line >5deg misorientation,and thick black line >15deg misorientation).Evidence of PSN is also present in(b).Fig 10—EBSD measurements revealing evidence of CDRX (an increase in misorientation from the center to the edge of a pre-existing grain and low-angle boundaries developing in the grain boundary region).(a )Orientation map of as-cast AZ31,deformed in compression to a strain of 0.6at a temperature of 350°C and a strain rate of 0.01s –1(thin white line >2deg misorientation,thin black line >5deg misorientation,and thick black line >15deg misorientation)and (b )the cumulative misorientation along the dotted line from A to the grain boundary.。
Microstructure and mechanical properties of wrought Mg-4.1Li-2.5Al-1.7Zn-1Sn alloy
Microstructure and mechanical properties of wroughtMg-4.1Li-2.5Al-1.7Zn-1Sn alloyRuizhi Wu1,2, a, Dayong Li1,b, Xuhe Liu2,c, and Milin Zhang2,d1College of Materials Science & Engineering, Harbin University of Science & Technology, Harbin,P.R. China 1500802Key Laboratory of Superlight Materials & Surface Technology (Harbin Engineering University),Ministry of Education, Harbin, P.R. China 150001a ruizhiwu2006@,b dyli@,c liuxuhe@,d zhangmilin@Keywords: Mg-Li alloy, deformation, microstructure, mechanical properties.Abstract.An Mg-Li-Al-Zn-Sn alloy was prepared by vacuum melting. The actual content of the elements in the alloy was determined using inductively coupled plasma-atomic emission mass spectrometry (ICP-AEMS). The density of the alloy was detected using Archimedes’ method. Extrusion and rolling deformation were carried out on this alloy. Its microstructures and mechanical properties were then studied with an optical microscope (OM), scanning electronic microscope (SEM), X-ray diffractometer (XRD), energy dispersive spectrometer (EDS), and tensile tester. The extruded alloy was composed of α-Mg and Mg2Sn phases and had good strength and elongation properties as well as a good comprehensive performance. After further rolling deformation, an Al-Li phase appeared due to atomic diffusion during the hot rolling process. Strain-hardening and the strengthening effect of the Al-Li phase further improved the strength of the alloy but decreased its elongation capacity.IntroductionSince the Mg-Li alloy was discovered in 1910, it has attracted a lot of attention from researchers because of its low density, high specific strength, stiffness, good processing performance, and dimensional stability. With these properties, it has shown great potential for use in applications in the aerospace, automotive, electronics, and defense industries [1-3].The Mg-Li alloy is the lightest metal material. Its binary phase diagram shows that, when the lithium content is less than 5.7%, the alloy displays an α single-phase that resembles a close-packed hexagonal structure. When the lithium content is more than 10.3%, the alloy shows a β single-phase, which is a body-centered cubic structure. At lithium contents between 5.7-10.3%, the alloy shows a two-phase organization [4]. Studies have shown that the addition of lithium causes the length of the c-axis in the hexagonal close-packed (HCP) structure to decrease, thus bringing about a decrease in the axial ratio c/a and rendering the alloy more able to undergo dislocation slip. This factor improves the deformability of the alloy [5].Al and Zn are two other elements that are commonly used in alloys. Appropriate amounts of Al added to alloys not only increases their strength and hardness but also improves their ductility and corrosion resistance. The addition of Zn can improve the deformation capacity of alloys [6, 7].Xiang Qi et.al[8] studied the influence of Sn on the microstructure and mechanical properties of a Mg-Li-Al-Zn alloy. Their results indicated that the addition of Sn refined the alloy, thereby improving its strength due to the formation of an Mg2Sn strengthening phase. When the Sn content was 1%, the grain size of the alloy was at the minimum size.Extrusion and rolling are the main deformation methods used for Mg-Li alloys. Deformation not only eliminates some casting defects but also causes dynamic recrystallization under certainconditions [9, 10]. This allows for the formation of refined alloys with improved comprehensiveIn this paper, a Mg-Li-Al-Zn-Sn alloy was prepared. After being subjected to two deformation modes, the microstructure and mechanical properties of the alloy obtained were determined.1. Materials and MethodsPure Mg, Li, Al, Zn, and Sn were used as raw materials. The alloy was created by vacuum induction melting using Ar gas as the protection gas in low-carbon steel molds. Homogenization treatments were performed at 300 o C for 24 h. The alloy ingot was extruded from Φ56 mm to 14 mm at 350 o C and denoted as an as-extruded alloy. As a last step, the extruded alloy was reheated and rolled at 260 o C to a final thickness of 3 mm.The chemical composition of the alloy was tested by inductively coupled plasma-atomic emission mass spectrometry (ICP-AEMS). The density of the alloy was determined by Archimedes’ method. The specimens for microscopic examinations were prepared using standard metallographic sample preparation methods. In brief, the specimens were etched with 1 vol% of nitric acid alcohol for 5-10 s.A LEICA DMIRM and JSM-6480 scanning electron microscopes (SEM) were used to observe the surface and fracture morphology of the alloy. A TTR III X-ray diffractometer (XRD) was utilized to identify the different phases in the alloy. An Energy Dispersive Spectrdmeter (EDS) was used to analysis micro-area composition.The tensile specimens were prepared according to the ASTM E8M-04 standard procedure. Tensile tests were carried out on an Instron4505 electronic universal testing machine with a speed of 1.5 mm/min. Five samples for each test were subjected to analysis along the extrusion and rolling directions.2. Results and Discussion2.1 The composition and density of the alloyThe analysis determined that the alloy composition was made of Mg-4.1Li-2.5Al-1.7Zn-1Sn. The density of the alloy was found to be 1.57 g/cm 32.2 The microstructure of the alloyThe microstructure of the extruded Mg-Li-Al-Zn-Sn alloy is shown in Figures 1a and 1b. The alloy was composed of a single α-phase, although some black matter appeared to be distributed in the matrix material. The grain size was small and its shape was equiaxial. These observations are typical of a material that has undergone dynamic recrystallization. Thus, it can be said that dynamic recrystallization occurred during the extrusion process.The microstructure of the alloy after rolling is shown in Figures 1c and 1d. Except for the black material, there also existed some eutectic compounds in the crystals, which may impact the performance of the alloy. The grain size for as-rolled samples was bigger than that for extrusion alloys.Figure 1. The microstructure of the alloy: (a) As-extruded alloy, (b) Magnified as-extruded alloy, (c) As-rolled alloy, and (d) Magnified as-rolled alloy.3.3 Phase analysisbelonged to Mg 22Sn in the 2Sn exists as a phase in the alloy.There were also sections of the rolled alloy that2θ/(°)I nte nsi t y /a .u .Figure 2. The XRD patterns of the alloy.30µmElements Wt./% At./%Mg 33.68 68.77Al 1.07 2.43Zn 1.47 1.36Sn 63.78 27.47A3.4 Mechanical properties of the alloyThe stress-strain curves of the extruded and rolled alloy are shown in Figure 4. Furthermore, Table 1 lists the mechanical properties of the two deformation state of the alloy. It also lists the corresponding performance parameters of commercially available LA141 Mg-Li alloy.Compared with the LA141 alloy, the as-extruded Mg-4.1Li-2.5Al-1.7Zn-1Sn had a higher tensile and yield strength with a considerable elongation capacity (>20%). Its specific strength and modulus are significantly higher than those of the LA141 alloy. These differences may be due to the α-phase (i.e., the α-phase alloy has higher strength than the β-phase alloy). It is also possible that Mg 2Sn, which was extensively distributed throughout the matrix, hinders dislocation slips when the alloy is deformed, thus playing a role in second phase strengthening.After rolling, the strength of the alloy was further increased, and its tensile strength reached 290.26 MPa. The elongation capacity of the alloy decreased but was still above 10%. The increase in strength may be explained in part by several factors, including work-hardening, additional deformation processes, and an increase in internal dislocation density. The latter causes flow stress to increase and improves the strength of alloys. The existence of an Al-Li phase after rolling could also be another reason for the increase in alloy strength. These two strengthening mechanisms, however, contribute to a decline in alloy plasticity. The alloy grain size increased, leading to a decline in its plasticity too.0510152025050100150200250300A s-extruded A s-rolledT e nsi le stre s s/MP a Strain/%Figure 4. Stress-strain curves of the alloy.Table 1. Mechanical properties of the alloyCondition As-extruded As-rolled LA141Tensile strength, MPa 267.51 290.26 144.69Specific tensile strength, cm ×105 170.39 184.87 105.03Yield strength, MPa 161.27 192.19 124.14Specific Yield strength, cm ×105 102.72 122.41 90.12Elastic modulus, GPa — 57.3 42.1Specific Elastic modulus,×106 — 36.49 31.12Elongation, % 21 11 24Density, g/cm 3 1.57 1.57 1.353.5 Fracture microstructure of the alloyThe fracture microstructure of the alloy is shown in Figure 5. The fracture microstructure of the as-extruded alloy was composed of a large number of small dimples. In some individual dimples,mechanism of the as-extruded alloy was ductile in nature. The fracture microstructure of the as-rolled alloy consisted of cleavage planes and a small number of dimples, which indicate that the fracture mechanism of the as rolled alloy can be ascribed to quasi-cleavage fractures.Figure 5. Fracture microstructure of the alloy. (A) As-extruded alloy and (B) As-rolled alloy.Summary1) An ultra-light Mg-Li-Al-Zn-Sn alloy was prepared by vacuum melting, then it was extruded and rolled. The density of the resulting alloy is 1.57 g/cm 3.2) The Mg(4.1)-Li(2.5)-Al(1.7)-Zn-Sn alloy ingot was subjected to two kinds of deformation processes: extrusion and rolling. The as-extruded alloy was found to be composed of the α-Mg and Mg 2Sn phases. After further rolling deformation, however, the alloy was found to consist of the α-Mg, Mg 2Sn, and Al-Li phases.3) Rolling deformations could further improve the strength of the alloy, but this resulted in a decrease of the elongation capacity.AcknowledgmentsThis work was supported by the National Natural Science Foundation of China (No. 51001034), China Postdoctoral Science Foundation(No. 20100481016) and Heilongjiang Postdoctoral Science Foundation.References[1]. R.Z. Wu, M.L. Zhang: Rev. Adv. Mater. Sci. Vol. 24 (2010), p.14[2]. H.Y Wu, Z.W. Gao and J.Y. Lin: J. Alloys Compd. Vol. 474 (2009), p.158[3]. Z.K. Qu, X.H. Liu, R.Z. Wu and M.L. Zhang: Mater. Sci. Eng. A Vol. 527 (2010), p.3284.[4]. L.Y.Wei, G.L.Dunlop and H.Westengen: Mater. Sci. Technol. Vol. 12 (1996), p.741[5]. C.H.Chiu, H.Y.Wu and J.Y.Wang: J. Alloys Compd. Vol. 460 (2008), p.246[6]. R.Z Wu, M.L Zhang: Mater. Sci. Eng. A Vol. 520 (2009), p.36[7]. T.C.Chang, J.Y.Wang and C.L.Chu: Mater. Lett. Vol. 60 (2006), p.3272[8]. Q. Xiang, R.Z.Wu and M.L. Zhang: J. Alloys Compd. Vol. 477 (2009), p.832[9]. T.C. Chang, J.Y. Wang and C.L. Chu: Mater. Lett. Vol. 60 (2006), p.3272[10]. R. Ninomiya and K. Niyake: J. Jpn. Inst. Met. Vol. 10 (2001), p.509[11]. R.Z Wu, Y.S Deng and M.L Zhang: J. Mater. Sci. Vol. 44 (2009), p.4132[12]. D.K. Xu, L. Liu and Y.B. Xu: Scripta Mater. Vol. 57 (2007), p.285(a)Material and Manufacturing Technology IIdoi:10.4028//AMR.341-342Microstructure and Mechanical Properties of Wrought Mg-4.1Li-2.5Al-1.7Zn-1Sn Alloydoi:10.4028//AMR.341-342.31。
vctic超晶格薄膜的微结构及超硬效应
Trans. Nonferrous Met. Soc. China 25(2015) 2581−2586Microstructure and superhardness effect of VC/TiC superlattice filmsXue-chao DONG 1, Jian-ling YUE 1,2, En-qing WANG 1, Miao-lei LI 1, Ge-yang LI 31. School of Aeronautics and Astronautics, Central South University, Changsha 410083, China;2. State Key Laboratory of Powder Metallurgy, Central South University, Changsha 410083, China;3. State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200030, ChinaReceived 22 September 2014; accepted 23 April 2015Abstract: Vanadium carbide/titanium carbide (VC/TiC) superlattice films were synthesized by magnetron sputtering method. The effects of modulation period on the microstructure evolution and mechanical properties were investigated by EDXA, XRD, HRTEM and nano-indentation. The results reveal that the VC/TiC superlattice films form an epitaxial structure when their modulation period is less than a critical value, accompanied with a remarkable increase in hardness. Further increasing the modulation period, the hardness of superlattices decreases slowly to the rule-of-mixture value due to the destruction of epitaxial structures. The XRD results reveal that three-directional strains are generated in superlattices when the epitaxial structure is formed, which may change the modulus of constituent layers. This may explain the remarkable hardness enhancement of VC/TiC superlattices. Key words: superlattice films; carbide films; microstructure evolution; superhardness effect; epitaxial growth1 IntroductionHard films have been successfully applied for theprotection of materials and particularly to increase theefficiency and extend the lifetime of cutting tools sincethe 1970s [1]. With the development of modern high-speed cuttings, it is demanded for protective hardfilms that both hardness and wear resistance should becontinuously increased. It was found in 1987 that superlattices deposited alternately by two kinds of materials at nanometer-scale thickness, i.e., superlatticefilms, could exhibit an anomalous increase of hardness ascompared with the monolayer films [2]. Therefore, tosynthesize superlattice films has become an importantapproach to obtain the films with high hardness andexcellent wear resistance.For superlattice films, their constituent layers can bemetal layers, nitrides, carbides or oxides of differentmaterials or a combination of one layer made of nitride,carbide, or oxide of one metal and the second layer madeof another metal [3]. However, less attention has beenpaid to carbide superlattice films compared with thosesuperlattices consisting of nitrides, oxides ormetals [4−8]. Generally, the carbide films have the samecrystal structure as nitride films and even higher hardness, lower friction coefficient and superior wear resistance performance. This supplied an alternative kind of material to obtain the superlattice films with excellent properties. Although some carbide superlattice films, e.g., VC/TiC [9,10], TiC/NbC [11] and VC/SiC [12], have been reported, further studies on the relationshipbetween the microstructure and mechanical properties of these superlattices are still required. For instance, there exist different results regarding whether the superhardness effect can be obtained in the VC/TiCsuperlattices [9,10], indicating that the strengthening mechanism also needs to be further investigated. This study intends to clarify the relationship between the microstructure and the superhardness effect of VC/TiC superlattices. A series of VC/TiC superlattice films with different modulation periods were prepared by magnetron sputtering method. Then, the microstructure and mechanical properties of these superlattice films were investigated. TiC/VC superlattices were chosen in this study for two reasons. First, both TiC and VC have already been applied practically due to their excellent properties. Second, they both exhibit the same structure (NaCl-type) and have smaller lattice mismatch (a VC =0.4160 nm, a TiC =0.4328 nm), which is believed toFoundation item: Project (51201187) supported by the National Natural Science Foundation of China Corresponding author:Jian-lingYUE;Tel:+86-731-88877495;E-mail:*****************.cn DOI: 10.1016/S1003-6326(15)63878-XXue-chao DONG, et al/Trans. Nonferrous Met. Soc. China 25(2015) 2581−2586 2582be favorable to obtaining high hardness increment by forming epitaxial structures.2 ExperimentalVC/TiC superlattices and TiC, VC monolayer films were prepared by SPC−350 multi-target magnetron sputtering system. VC (99.8% in purity) and TiC (99.8% in purity) targets were placed on two radio frequency cathodes, while the metallic Ti target (99.99%) was placed on the DC cathode. Mirror polished stainless steel substrates were ultrasonically cleaned in acetone and absolute alcohol before being mounted on a rotatable substrate holder in the vacuum chamber and the distance between target and substrate was about 5 cm. The base pressure was pumped down to 2×10−4 Pa before deposition. Ar (99.99% in purity), whose pressure was kept at 0.6 Pa, was introduced into the chamber. The VC/TiC superlattice films were prepared at 300 °C by alternating deposition of TiC and VC layers. In order to improve the adhesion, Ti transition layer with 100 nm in thickness was deposited before depositing superlattices.A series of superlattices with different modulation periods but the modulation ratio of about 1:1 were prepared by controlling the target power and the resident time that substrates were exposed to the target. And the VC and TiC monolayers were deposited under the same conditions as that of VC/TiC superlattices. The total thickness of each specimen was about 2 μm.The composition of the films was determined by energy dispersive X-ray spectrometry analysis (EDXA). The microstructures of all the films were characterized by Rigaku D/max−2550/PC X-ray diffraction (XRD) with Cu Kα excitation radiation, and high-resolution transmission electron microscopy (HRTEM) using a Philips CM200−FEG. The hardness (HV) and elastic modulus (E) of all the samples were measured by a Fischerscope H100VP nanoindenter equipped with a Vicker’s indenter and the values for specimen were evaluated by an average of at least 20 measurements.3 Results and discussion3.1 Microstructure of VC/TiC superlattice filmsEDXA results show that the C/Ti mole ratio of the TiC monolayer films is 52.8:47.2, and the C/V mole ratio of the VC monolayers is 47.6:52.4. This indicates that both TiC and VC monolayers in this work are near stoichiometric.The low-angle XRD patterns of the VC/TiC superlattice films with different modulation periods (Λ) are shown in Fig. 1. The low-angle diffraction peaks, which clearly appear to be systematically distributed in the diffraction patterns, can be attributed to the modulation period Λ of the superlattices. This indicates an obvious layered structure for all superlattices. Correspondingly, the modulation period of the superlattices can be figured out from [13]()1sinsin2−−=nnΛθθλ(1)where λ is the X-ray wavelength, θn and θn−1 are the diffraction angles, n is diffraction series. Fitting of the 2θvalues for the diffraction peaks into Eq. (1) resulted in the Λ of superlattices, which are shown in Fig. 1.Fig. 1 Low-angle XRD patterns of VC/TiC superlattices with modulation ratio of 1:1Figure 2 displays the cross-sectional HRTEM image and selected area electron diffraction (SAED) pattern of the VC/TiC (Λ=5.2 nm) superlattice film. The SAED pattern on the top-left corner is indexed as a NaCl-type structure, and the low-magnification image on the bottom-left corner of Fig. 2 clearly shows that this superlattice specimen possesses a well-defined composition modulation structure consisting of VC (dark layers) and TiC (light layers). The high-magnification image on the right side shows that the lattice fringes continuously go through several modulation layers and interfaces. This indicates that the two constituent layers,Fig. 2 Selected area election diffraction (SAED) pattern, low-magnified and high-magnified cross-sectional HRTEM images of VC/TiC superlattice with Λ=5.2 nmXue-chao DONG , et al/Trans. Nonferrous Met. Soc. China 25(2015) 2581−2586 2583VC and TiC, grow coherently each other and the epitaxial structure is therefore formed in this superlattice film.Figure 3 shows the high-angle XRD patterns of the VC/TiC superlattices along with VC and TiC monolayers. As can be seen, NaCl-type VC and TiC monolayer films present relatively broad (111) diffraction peaks. By contrast, the VC/TiC superlattices with modulation period ranging from 1.9 to 12 nm only present one fundamental peak, which is located between TiC (111) and VC (111) peaks and thus can be regarded as VC/TiC superlattice (111) peak. Additionally, when Λ is less than 6.8 nm, the VC/TiC superlattice (111) peaks are much higher and sharper than those of VC or TiC monolayers. With the increase of Λ, the superlattice (111) peaks gradually decrease in the intensity and also broaden. After Λ exceeds 12 nm, the superlattice (111) peaks become quite diffusive and are divided into two peaks, as shown in the local amplification figure on the left side of Fig. 3.Combined with the HRTEM results in Fig. 2, the higher (111) peak intensity of VC/TiC superlattices with Λ less than 6.8 nm compared with that of VC or TiC monolayers, as shown in Fig. 3, indicates that these superlattices form an epitaxial structure and their crystal integrity is greatly improved. Similar phenomena have also been observed in some other superlattice systems such as ZrO 2/TiN [14], VC/AlN [15], HfN/Si 3N 4 [16] and TiAlN/AlON [17], and the corresponding reasons have also been explained from both thermodynamics and kinetics factors of film growth in these studies [14−17].Furthermore, the VC/TiC superlattice (111) peaks are located between the VC (111) and TiC (111) peaks when Λ is lower than 12 nm, which can be attributed to the fact that the two (111) peaks from VC and TiC sublayers in superlattices approach each other and then overlap when the epitaxial structure is formed. Correspondingly, the (111) interplanar spacings (d 111) of TiC and VC sublayers also approach each other. That is, d 111 of VC layers increases along the direction perpendicular to the superlattice interface (out-of-plane), while d 111 of TiC decreases along the same direction. Therefore, it can be inferred that the uniaxial tensile strains in VC and the uniaxial compressive strains in TiC sublayers along the out-of-plane direction are induced from epitaxial structure. When Λ increases to 25.1 and 26.9 nm, the superlattice (111) peaks are divided into two parts, i.e., VC (111) and TiC (111) peaks, suggesting that the strains start to be relaxed in superlattices due to the destruction of epitaxial structure.On the other hand, in order to grow epitaxially on the VC sublayers (a VC =0.4160 nm) in superlattices, the TiC sublayers with bigger lattice constant (a TiC = 0.4328 nm) will be subjected to the biaxial compressive strains in the (111) plane (i.e., parallel to the superlattice interface, in-plane). In contrast, the VC sublayers will be under the biaxial tensile in-plane strains. Combined with the above discussion for the out-of-plane strains in superlattices, the strain states of VC/TiC superlattices with epitaxial structure are shown in Fig. 4. It reveals that when the epitaxial structure is formed in superlattices, three directional strains are generated in the different sublayers, for example, the TiC sublayers are under the biaxial (in-plane) and uniaxial (out-of-plane) compressive strains, while the VC sublayers are under biaxial and uniaxial tensile strains, respectively.3.2 Hardness and elastic modulus of VC/TiCsuperlattice filmsFigure 5 demonstrates the nanoindentation hardness (HV) and elastic modulus (E ) of VC/TiC superlattices as a function of modulation period. The hardness and modulus of VC and TiC monolayers are also shown in Fig. 5. As can be seen, the hardness and modulus ofVC/TiC superlattices can be significantly increasedcompared with VC and TiC monolayers. With theincrease of modulation period (1.9 nm≤Λ ≤ 5.2 nm), the HV and E increase rapidly and reach peak values ofFig. 3 High-angle XRD patterns of VC, TiC monolayer and VC/TiC superlattice filmsXue-chao DONG , et al/Trans. Nonferrous Met. Soc. China 25(2015) 2581−25862584Fig. 4 Strain states in VC /TiC superlattices with epitaxial structureFig. 5 Dependence of hardness and elastic modulus of VC/TiC superlattices on modulation period41.9 GPa and 330 GPa when Λ is 5.2 nm. However, after Λ exceeds 5.2 nm, the HV and E begin to decrease by further increasing Λ and finally tend to be near the rule-of-mixture value.3.3 Strengthening of superlattice filmsSo far, there exist two main theories, the Koehler’s modulus-difference strengthening theory [18] and the alternating-stress strengthening theory [19], which are always used to explain the superhardness effects of superlattices. Both of these theories are based on the model of dislocations being blocked at interfaces.According to the Koehler’s strengthening theory, the modulus difference between two superlattice constituents is the key factor to restrict the dislocations from gliding across the interfaces and thus determine whether the superhardness effect could be obtained in superlattices. The hardness increment (∆H max ) of superlattices compared with the constituent material with lower hardness can be calculated by this theory asA max 3sin =8πRG H mθΔ (2)where θ is the smallest angle between the interface andthe dislocation slip plane of constituent material with smaller elastic modulus; m is the Taylor factor, calculated as 0.3 for TiC and VC [20]; R =(G B −G A )/(G B +G A ), where G A and G B are shear moduli of two constituents, respectively, and G B >G A , G =E /[2(1+υ)] (E is the elastic modulus and υ is the Poisson ratio). Fitting of the elastic modulus of TiC and VC (E TiC =255 GPa, E VC =245 GPa), and υ of 0.25 for TiC and VC [20], and θ of 45° for NaCl-type VC/TiC superlattices into Eq. (1), yields a ∆H max of 0.6 GPa compared with VC monolayer. This value is much lower than the measured ∆H max of about 12.6 GPa in this study, indicating that it is not enough for the Koehler’s strengthening theory to directly explain the hardness enhancement of VC/TiC superlattices.In reference to the alternating-stress theory [19], it is known that when the dislocations go across the interfaces in superlattices, they are resisted not only by the potential barriers coming from the change of line energy of dislocations due to the two different modulus layers (Koehler’s strengthening theory), but also by that from the alternating-stress field. The maximum hardness increase from the alternating-stress field can be estimated asmax max 10τ≅ΔH (3)where τmax =(2/π2)A ηE [21] is the maximum shear stress at the interface; A is the modulation amplifying factor that is closely correlated to structural parameters of superlattices such as modulation period, modulation ratio, roughness and width of interface; E is the weighted average elastic modulus of two constituents of superlattices; and η is the lattice mismatch between two layers of superlattices.Based on the studies from MIRKARIMI et al [22] and SHINN and BARNETT [23], A takes the value of 0.5 for calculation in this study. Using the data of E of TiC and VC, 250 GPa, and η of 0.04 (a VC =0.4160 nm, a TiC =0.4328 nm), we can roughly estimate ∆H max as ~10.1 GPa. However, this estimation is not enough to explain the actual hardness enhancement. It should beXue-chao DONG, et al/Trans. Nonferrous Met. Soc. China 25(2015) 2581−2586 2585noted that even the sum of the hardness increase from the alternating-stress field model and modulus-difference model, 10.7 GPa, is still lower than the measured value (12.6 GPa).Actually, the discrepancy between the estimated value from these two strengthening theories and the actual result is due to approximation made when not considering any changes in the modulus of constituent layers in the superlattices under epitaxial strains. It is noted that the modulus changes have an important influence on the hardness increment based on these two theories. Some studies [24,25] have demonstrated that an applied elastic strain can significantly change the in-plane biaxial modulus of thin films, that is, the biaxial modulus increases with compressive strains, and reduces with tensile strains. As shown in Fig. 5, the remarkable increase of elastic modulus of superlattices is closely related to the epitaxial strains when the modulation period is less than a critical thickness. Thus, the biaxial modulus of TiC sublayers with relatively high modulus in the presence of VC/TiC superlattices should increase, while the corresponding biaxial modulus of VC sublayers with lower modulus should decrease. This leads to the fact that the modulus difference, R, between TiC and VC sublayers in superlattices would be increased under the epitaxial strains. Therefore, the hardness enhancement in VC/TiC superlattices should increase in reference to the Koehler’s strengthening theory as well as the alternating-stress field theory, when taking into account the modulus change of constituent layers coming from the epitaxial strains.According to the (111) peak positions of TiC, VC monolayers and VC/TiC superlattices in Fig. 3, it can be calculated that the out-of-plane compressive strain in TiC sublayers is about 2.5%, and the out-of-plane tensile strain in VC sublayers is about 1.5%. Also, the calculations from Refs. [24,25] revealed that ~2.5% compressive strain could increase the biaxial modulus of materials by about 100%, and about 1.5% tensile strain could lead to 30% decrease in biaxial modulus. It is noted that these calculations agree well with the measured maximum modulus of VC/TiC superlattices, as shown in Fig. 5. If it is presumed that the in-plane strain is equal to the out-of-plane strain in constituent layers, we can accordingly obtain ΔH max of 9.7 GPa by the Koehler’s theory, and 13.8 GPa by the alternating-stress field theory in VC/TiC superlattices. Here, it can be found that the sum of hardness increase from the Koehler’s theory and alternating-stress field theory, 23.5 GPa, seems to be quite competent to explain the actual experimental results in this work.From the present VC/TiC superlattices, it can be deduced that in order to synthesize the superlattices with superhardness effect, the epitaxial structure should be formed primarily. On the other hand, the three- directional strains induced from the epitaxial structures should increase the modulus difference between two constituent layers in superlattices. Otherwise, the superhardness effect could not even be obtained in some superlattice systems with epitaxial structure, for example, our previously reported ZrO2/TiN superlattices [14], in which the modulus difference between two constituent layers was found to be decreased by the epitaxial strains.4 Conclusions1) The VC/TiC superlattices can form an epitaxial structure when the modulation period is less than a critical value (~5.2 nm), resulting in that their hardness and elastic modulus are significantly increased. Further increasing the modulation period, the hardness and elastic modulus of VC/TiC superlattice films gradually decrease to be close to the rule-of-mixture value due to the destruction of epitaxial structure.2) The hardness enhancement of the VC/TiC superlattices is closely related to the three-directional strains, which are induced from the epitaxial structure and may change the interplanar spacing and thus the modulus of constituent layers when the epitaxial structure is formed.3) It is found that when taking into account the modulus change of VC and TiC sublayers in superlattices under the three-directional strains, the remarkable hardness enhancement of VC/TiC superlattices could be interpreted by combining the Koehler’s theory and the alternating-stress field theory.References[1]KANG M S, WANG T G, SHIN J H, NOWAK R, KIM K H.Synthesis and properties of Cr−Al−Si−N films deposited by hybridcoating system with high power impulse magnetron sputtering(HIPIMS) and DC pulse sputtering [J]. 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Effect of superlattice layer elasticmoduli on hardness [J]. Appl Phys Lett, 1994, 64: 61−63.[24]JANKOWSKI A F, TSAKALAKOS T. The effect of strain on theelastic constants of noble metals [J]. J Phys F: Met Phys, 1985, 15:1279−1292.[25]CAMMARATA R C, SIERADZKI K. Eff'ects of surface stress on theelastic moduli of thin films and superlattices [J]. Physical ReviewsLetters, 1989, 62(17): 2005−2008.VC/TiC超晶格薄膜的微结构及超硬效应董学超1，岳建岭1,2，王恩青1，李淼磊1，李戈扬31. 中南大学航空航天学院，长沙 410083；2. 中南大学粉末冶金国家重点实验室，长沙 410083；3. 上海交通大学金属基复合材料国家重点实验室，上海 200030摘 要：采用磁控溅射工艺制备VC/TiC超晶格薄膜，并采用EDXA 、XRD、HRTEM和纳米力学探针研究调制周期对超晶格薄膜的微结构和力学性能的影响。
Microstructural evolution and mechanical propertie
International Journal of Minerals, Metallurgy and Materials Volume 25, Number 11, November 2018, Page 1294https:///10.1007/s12613-018-1682-8Corresponding author: Hamed Jamshidi Aval E-mail:h.jamshidi@nit.ac.ir© University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018Microstructural evolution and mechanical properties of friction stir-weldedC71000 copper–nickel alloy and 304 austenitic stainless steelHamed Jamshidi AvalDepartment of Materials and Industrial Engineering, Babol Noshirvani University of Technology, Shariati Avenue, Babol, 47148-71167, Iran(Received: 20 February 2018; revised: 29 May 2018; accepted: 11 June 2018)Abstract: Dissimilar joints comprised of copper–nickel and steel alloys are a challenge for manufacturers in modern industries, as these met-als are not thermomechanically or chemically well matched. The present study investigated the effects of tool rotational speed and linear speed on the microstructure and mechanical properties of friction stir-welded C71000 copper–nickel and 340 stainless steel alloys using a tungsten carbide tool with a cylindrical pin. The results indicated that a rotational-to-linear speed ratio of 12.5 r/mm did not cause any macro defects, whereas some tunneling defects and longitudinal cracks were found at other ratios that were lower and higher. Furthermore, chro-mium carbide was formed on the grain boundaries of the 304 stainless steel near the shoulder zone and inside the joint zone, directing carbon and chromium penetration toward the grain boundaries. Tensile strength and elongation percentages were 84% and 65% of the corresponding values in the copper–nickel base metal, respectively.Keywords: dissimilar friction stir welding; copper–nickel alloy; austenitic stainless steel; microstructure; mechanical properties1. IntroductionCopper–nickel alloys exhibit substantial corrosion resis-tance and anti-algae properties against biological sediments. Pure copper is not stable in oxygenated electrolytes, espe-cially in marine and chlorine ion environments where cop-per–nickel alloys are widely used, with copper as the main component [1]. The addition of nickel to copper improves the mechanical strength, durability, and resistance to corro-sion, abrasion, and cavitation in sea and polluted water. This alloy also exhibits significant stress corrosion cracking and corrosion fatigue resistance. Corrosion resistance can be in-creased by adding more nickel to copper–nickel alloys [2]. Since these alloys can be easily assembled and welded, they are prime candidates for plumbing systems, ship bodies, and other marine structures.Generally, stainless steel plays a major role in the modern world. Welding of austenite stainless steel is known for two important properties: maintenance of corrosion resistance and prevention of crack formation. Dissimilar joints of coatings on offshore platform insulators, achieved by different tech-niques, are among copper–nickel plate applications for corro-sion prevention. Other applications include the joining of copper–nickel pipes with steel flanges and/or direct joining of these pipes with steel pipes in marine industries [3].Nevertheless, welding of dissimilar metals is always challenging because of numerous factors. These include different melting points, thermal conductivity, and thermal expansion coefficients; galvanic corrosion; the high solidi-fication rate of molten copper; entry of molten copper into steel grain boundaries (especially in the heat-affected zone (HAZ)); formation of hot cracks; high copper oxidation at high temperatures; and type of filler metal [4−9]. It is essen-tial to select the appropriate filler metal and welding para-meters for dissimilar-metal fusion welding of copper–nickel and stainless steel alloys in order to reduce probable defects (e.g., cavitation and gas cavities).Recent developments in solid-state welding have made it an alternative to fusion welding. In comparison to other welding techniques, friction-stir welding is a solid-state technique with an outstanding combination of high speed, precision, and variety. Among different welding methods, friction-stir welding of dissimilar alloys is important due to the ability to join alloys with different properties. In addition,H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1295)different welding configurations in this method (e.g., lap and butt joints) make it applicable in different situations.Few studies have been conducted on friction-stir welding of dissimilar copper and stainless steel alloys. In this regard, Imani et al. [10] investigated a pure copper and stainless steel joint with a thickness of 3 mm using friction-stir weld-ing. It was found that the tool offset toward the copper side played a significant role in eliminating defects in the joints. In addition, Ramirez et al. [11] examined the effects of tool offset on the microstructure and mechanical properties of joints in friction-stir welding of pure copper and 316 stain-less steel with a thickness of 2 mm. They studied 0, 0.6, and 1.6 mm offsets relative to the joint interface. When a major part of the tool was on the steel side, the joint efficiency was 55% of copper base metal. Maximum joint efficiency, i.e., 87% of copper base metal, was reported in the 0.6-mm off-set relative to the joint interface.Furthermore, Najafkhani et al. [12] studied the joint of pure copper and 316 stainless steel with a thickness of 5 mm using friction-stir welding. In their study, all joints cracked from the heat-affected zone of the copper base metal. The highest tensile strength and elongation percentage were 220 MPa and 7%, respectively. In addition, Shamsujjoha et al. [13] studied the lap joint of pure copper with 1018 carbon steel using friction-stir welding. They found that the joining process at the interface was both mechanical and metallur-gical. Jafari et al. [14] also studied the friction-stir welding of pure copper and 304 stainless steel with a thickness of 3 mm. The heat input from the welding increased the grain size in the heat-affected zone and decreased joint ductility by increasing the number of welding passes.According to the literature, there are no studies on the friction-stir welding of copper−nickel and austenite stainless steel alloys. Accordingly, the present study investigated the effects of process parameters on the microstructure and me-chanical properties of friction stir-welded C71000 cop-per−nickel and 304 stainless steel alloys using a tungsten carbide tool with a cylindrical pin. Optical microscopy and scanning electron microscopy (SEM) were used to study the microstructure and detect the created phases in different zones. The mechanical properties of joints were also eva-luated by tensile and microhardness tests.2. ExperimentalIn the present study, C71000 copper−nickel and 304 aus-tenite stainless steel plates with thicknesses of 2 mm were used. Both plates were cut perpendicular to the rolled metal direction and had a dimension of 50 mm × 100 mm. The chemical compositions and mechanical properties of alloysare listed in Tables 1 and 2. The plates were welded in a buttjoint configuration. The copper−nickel alloy was on the re-treating side, while the stainless steel alloy was on the ad-vancing side. According to the literatures [10−11], 0.75 mmof the tool axis was offset to the copper−nickel alloy relativeto the joint interface. Fig. 1 shows the schematic of the tool offsetting procedure. A tungsten carbide tool with a cylin-drical pin with a height of 1.8 mm was used for welding. Fig.2 demonstrates the dimensions and geometry of the appliedtool in welding and Table 3 indicates the welding parameters.The present study selected two rotational speeds of 800 and1000 r/min and three linear speeds of 40, 60, and 80mm/min.Table 1. Chemical composition of alloy wt%ZnMnFeCrCuCNiAlloy0.90.010.05―Base0.0519.12C71000―1.20Base18.500.440.058.10SS304Table 2. Mechanical properties of alloysAlloyUltimate tensilestrength / MPaYieldstrength / MPaMicro-hardness,HV0.1Elonga-tion / %C71000338 110 9032 SS304585 210 15242Fig. 1.Schematic illustration of friction stir butt welding.Fig. 2. Tool geometrical characteristics.1296Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Table 3. Friction stir welding process parametersRotational-to-linear speed ratio / (r ⋅mm –1)Linear speed / (mm ⋅min –1)Rotational speed / (r ⋅min –1)Sample No.20.00 40 800113.33 60 800 2 10.00 80 800 3 25.00 40 1000 4 16.65 60 1000 5 12.508010006The samples were transversely cut for metallographic studies. A marble solution was used for etching the micro-structure on the stainless steel side after sanding and polish-ing, whereas a nitric acid and distilled water solution was used for the copper–nickel alloy. Scanning electron micro-scopy (SEM) and X-ray diffraction (XRD) were used to evaluate the joint interface and examine the distribution and type of intermetallic compounds in the joint cross section. The mechanical properties of the joint were investigated us-ing a tensile test according to the ASTM E8-M03 standard. The tensile test was carried out at a crosshead speed of 1 mm/min. A Vickers microhardness testing machine with a load of 3 N and test time of 15 s was used to evaluate the hardness distribution of a joint cross section.3. Results and discussion3.1. Weld appearanceThe qualitative test of the welded samples indicated that samples No. 1–5 had defects. Longitudinal cracks on the copper–nickel side or tunneling defects on the stainless steel side were observed in all defected samples. Fig. 3 shows the effects of rotational and linear speed on the appearance of welding samples No. 1, 3, 4, and 6, as representative sam-ples containing cracks, tunneling defects, and defect-free welds. It is generally difficult to explain the causes of de-fects in the samples; however, the heat input may be an in-fluential factor. Many researchers have introduced various analytical, numerical, and empirical models in order to evaluate the relationship between rotational and linear tool speed and heat input and to examine their effects on the temperature distribution in the friction-stir welding proce-dure. With a proper estimation, the rotational-to-linear speed ratio can be considered a measure of welding heat input.In this study, samples No. 3 and 4 received the least and most heat input, respectively. The lower temperature of sample No. 3 caused insufficient material flow into the stir zone. After the tool was moved forward, the flow of material stopped before arriving at the advancing side. Therefore, there was inadequate material to fill the hole on the advanc-ing side (stainless steel). The tunnel hole led to the loss of joint strength in this sample, and the two parts were easily separated.Fig. 3. Surface appearance of welded samples: (a) No. 4; (b) No. 1; (c) No. 6: (d) No. 3.Fig. 3 presents the longitudinal cracks because of a tunneling defect in sample No. 3. Fig. 4 shows the effect of welding heat input on longitudinal crack length. It can be seen that by increasing the rotational-to-linear speed ratio (increasing welding heat input), the maximum temperature in the joint increased, which led to the higher temperature gradient in the welded samples. The significant difference in thermal conductivity of copper–nickel and stainless steel al-loys (thermal conductivity of copper–nickel is 2.8 times higher than that of stainless steel) [15–16] at a high temper-ature gradient produced longitudinal cracks as a result of thermal stress in the joint. According to the visual inspection of welded samples, a tunneling defect developed in the joint at a rotational-to-linear speed ratio of less than 10.00 r/mm.H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1297)On the other hand, at the rotational-to-linear speed ratio of 13.33 r/mm or higher, longitudinal cracks were formed at the joint interface.3.2. Macrostructure and microstructureThe evaluation of mechanical and metallurgical proper-ties was only carried out for sample No. 6 because it had no defects. The macrostructure of the joint and microstructure of different zones are shown in Figs. 5 and 6. The micro-structure of stainless steel included austenite and δ-ferrite with a grain size of (40 ± 5) µm (Fig. 5(b)). Although the quantity of ferrite phase was not significant, the presence of δ-ferrite could improve the formation of the sigma phase inalloys during friction-stir welding [17].Fig. 4. Effect of the rotational-to-linear speed ratio on cracklength.Fig. 5. Optical images of different zones of sample No. 4: (a) macrostructure of welded sample No. 4; (b) base metal of AISI 304; (c) base metal of C71000; (d) TMAZ in AISI 304 side; (e) SZ in AISI304 side; (f) TMAZ in C71000 side as marked by zone I in (a).1298 Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 6. Microstructure of stir zone of sample No. 4: (a) microstructure of zone II in Fig. 5(a); (b) SZ in C71000 side; (c) SEM image of zone I in (a); (d) SEM image of zone II in (a).The copper–nickel microstructure had a grain size of (50 ± 4) µm and an average particle size of (10 ± 3) µm in the grain boundaries. The results of energy dispersive X-ray spectroscopy (EDS) indicated that these particles were nick-el-rich oxides with iron and zinc (Fig. 7). The stir zone mostly consisted of copper–nickel alloy, which was likely due to the lower flow stress of copper–nickel alloy [18] and location of the main part of the tool on the copper–nickel side. Different behaviors of the two alloys in the etchant so-lution confirmed this finding.Fig. 7. Element mapping result of base metal C71000 alloy.As shown in Fig. 5(a), a steel layer was drawn from the advancing zone to the retreating zone (zone I). The joint cross section as a result of friction-stir procedure consisted of the stir zone (SZ), thermomechanically affected zoneH. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1299)(TMAZ), and heat-affected zone (HAZ). The mechanical behaviors of the welding materials, especially the welding zone hardness, were affected by high plastic deformation and high temperature during the friction-stir welding. The stir zone microstructure in the friction-stir weld had smaller and equiaxed grains in comparison with the base metal due to high plastic deformation and stir resulting from the tool pin.As presented in Fig. 5(f), grains in the thermomechani-cally affected zone on the copper–nickel side were elon-gated, which is exclusive to this zone [19]. The steel layer on the copper–nickel side contained recrystallized cop-per–nickel grains (Fig. 5(f)). On the other hand, the stir zone microstructure on the copper–nickel side contained equiaxed grains with a size of (15 ± 4) µm as a result of dynamic re-crystallization in this zone (Fig. 6(b)). The oxide particles observed in the copper–nickel base metal are shown in this figure. These particles were mainly at grain boundaries with a size of (5 ± 2) µm and prevented the growth of stir-zone grains.The EDS results showed that zinc and iron concentra-tions in the oxide particles increased (Figs. 7 and 8). The stir zone on the steel alloy side contained small recrystallized grains with a size of (5 ± 1) µm (Fig. 5(e)). Clearly, the grain size in the copper–nickel stir zone was greater than that of the steel-stir zone. The temperature and deformation rate in the friction-stir procedure had inverse effects on the grain size of the stir zone. In fact, an increase in the defor-mation rate led to a reduced grain size, and a rise in temper-ature increased the grain size in the stir zone [20].Fig. 8. Element mapping result of stir zone of C71000 side.The advancing side showed the highest temperature and deformation [21]. According to the stir zone microstructure results, the deformation effect was dominant on the steel side, and the grain size of stir zone reduced relative to the copper–nickel alloy side. On the contrary, elongated grains did not exist in the thermomechanically affected zone on the steel side (Fig. 5(d)). However, annealing twins were found across the base metal, whereas there were fewer twins in the thermomechanically affected zone of the steel. There were no twins in the stir zone on the steel side. An interesting point in the microstructure study was the occurrence of a specific layer-by-layer structure at the interface between copper–nickel and steel alloys near the tool shoulder (Fig. 6(a)). The SEM images of different zones in Fig. 6(a) are presented in Figs. 6(c) and 6(d).According to the line scan analysis presented in Fig. 9, the layer-by-layer structures consisted of copper-rich layers adjacent to iron-rich layers. Based on the comparison of the chemical composition of the copper-rich layer and cop-per–nickel base metal, this zone belonged to the cop-per–nickel base metal. However, the iron-rich layer did not match the chemical composition of steel base metal. The highest mass percentages of copper and chromium in the iron-rich layer were 9% and 30%, respectively. The iron-rich layer had a higher copper percentage, which in-creased to 31wt% in some layers.The high percentages of nickel and copper as austenite stabilizers could promote the formation of austenite phase. Generally, welding of austenite stainless steel can cause de-fects, including formation of the brittle phase, hot cracks, and carbide–chrome in grain boundaries. Copper, as an auste-nite-forming element, eliminates the δ-ferrite and sigma phases. Furthermore, the copper–nickel alloy limits the sigma phase by increasing the cooling rate from 600 to 800°C [22].1300Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 9. SEM image and line scans of chemical elements at the layer structure: (a, b) zone I in Fig. 6(a): (c, d) Fig. 5(f).The sigma phase is very hard and brittle. Its value in-creases by increasing the percentage of chromium, molyb-denum, and silicon, but decreases by increasing the nitrogen, nickel, and carbon contents. Prevention of sigma phase for-mation in stainless steel is difficult when the chrome per-centage is about 20wt%. When the chrome percentage is less than 20wt%, the sigma phase is not observable in auste-nite stainless steels. Due to the very low amount of chrome (up to 9wt%) in the layered structure, formation of sigma phase is not expected.Fig. 10 shows the XRD analysis of the iron-rich zone in the layered structure (point A in Fig. 6(c)); the austenite phase is the only existing phase in this zone. The high per-centage of nickel and copper prevented the formation of sigma phase as expected. The line scan analysis (Fig. 9) in-dicated that nickel concentrations reduced in layer bounda-ries but increased in the iron-rich layers due to nickel migra-tion from the interface to iron-rich layers.According to the EDS results (Fig. 11(a)) regarding point A in Fig. 6(c), the nickel and copper percentages were 24wt% and 21wt%, respectively, indicating the diffusion ofFig. 10. XRD pattern of iron rich layer structure.nickel and copper from the copper–nickel alloy at the inter-face of steel alloy due to the close proximity of this region to the tool shoulder and high temperature of the zone. Ac-cording to the EDS results (Fig. 11(b)), regardless of the in-creased percentage of copper and nickel in the grain boun-daries of the recrystallized zone on the steel side, the high percentage of chrome indicates the increased effect of this element by moving toward the stir zone of the stainless steel. Carbon present in the grain boundaries indicates chrome carbide formation at the joint interface near the tool shoulder.H. Jamshidi Aval, Microstructural evolution and mechanical properties of friction stir-welded C71000 copper–nickel (1301)The chemical compositions of these spots indicate that chrome and carbon move toward high-energy zones and form chrome carbide. Formation of carbide and a chrome-free zone around the grain boundary can severely degrade corro-sion resistance of the joint. Analysis of point C (Fig. 11(c)) indicates that this zone belongs to 304 stainless steel. The transient zone in the joint interface can affect the mechanical properties of the joint. Partial diffusion and formation of iron- and copper-rich layers, as shown in Fig. 9, are alsoobserved in zone I of Fig. 5(a).3.3. Hardness evaluationThe joint microhardness profile at the mid-thickness of the weld cross section is presented in Fig. 12. Hardness of the stir zone increases by moving from the steel base metal. According to the Hall-Petch equation, smaller grains have greater hardness; accordingly, hardness increases by de-creasing the grain size and increasing the particle boundary density. Hardness near the interface fluctuates considering the layer-by-layer structure. This structure produces impor-tant features, such as non-uniform hardness profiles and stress concentration zones. The stir zone on the cop-per–nickel side had a more uniform hardness profile and lower quantity. Hardness gradually decreased to the level of copper–nickel base metal by moving toward the cop-per–nickel base metal.3.4. Tensile properties and fractographyThe stress–strain curves for the base metals and joint are shown in Fig. 13. The yield strength and tensile strength of the joint are 103 MPa and 285 MPa, respectively, while elongation is 21%; these values are significantly lower than the corresponding values in the base metals. Tensile strength and elongation of joint were 84% and 65% of the corres-ponding values, respectively in the copper–nickel base metal. It should be noted that fracture occurred in the weld nugget and at the interface of steel and copper–nickel. The hardness profile shows sudden fluctuations, which cause stress con-centrations and joint strength degradation.Fig. 12. Microhardness profiles of cross-section of joint No. 6.Fig. 11. EDS analysis of points A (a), B (b), and C (c) in Fig. 6(c).1302 Int. J. Miner. Metall. Mater ., Vol. 25, No. 11, Nov. 2018Fig. 13. Stress–strain curve of base metals and welded sample No. 6.The fractured cross section was investigated by SEM af-ter the tensile test. Fig. 14 shows the fractured section and SEM image. The SEM image of the fracture zone shows a brittle cleavage fracture, along with plastic deformation and small uniform holes on the surface. In the brittle cleavage fracture, the crack propagation corresponds to the successive and repeated breaking of atomic bonds along specific crys-tallographic planes. The fracture surface has a faceted tex-ture because of different orientations of the cleavage planes in the grains. In this type of fracture, no substantial plastic deformation occurs and the crack propagates very fast, nearly perpendicular to the direction of the applied stress. In the ductile fracture mode, spherical dimples correspond to microvoids initiating crack formation. Each dimple is half the size of the microvoid, which is formed and then sepa-rated during the fracture process. In the welded sample, brit-tle and ductile failures simultaneously occurred, which could be attributed to the transient zone (Fig. 12) and sud-den fluctuations in the hardness of the sample.Fig. 14. SEM image of fracture surface of the joint No. 6.4. ConclusionsThe present study investigated the friction-stir welding of C71000 and AISI304 stainless steel with a cylindrical pin tool and the following results were obtained.(1) Lack of proper material flow occurred as a result of low temperature at a rotational-to-linear speed ratio of 10 r/mm; therefore, there was not adequate material to fill the hole as the tool traveled forward on the advancing side (stainless steel). In case of rotational-to-linear speed ratio of greater than 20 r/mm, the high heat input produced a higher temperature gradient and resulted in the formation of longi-tudinal cracks as a result of thermal stress in the joint sec-tion.(2) The grain size on the copper–nickel side was larger than that of the stainless steel side. The stirring phenomena during friction-stir welding eliminated annealing twins in the stainless steel base metal and a uniform microstructure with small equiaxed grains formed in the stir zone. Tensile strength and elongation of joint were 84% and 65% of the corresponding values, respectively in the copper–nickel base metal. The fracture surface indicated brittle cleavage and plastic deformation behaviors.(3) Heat and plastic deformation caused element diffu-sion at copper- and iron-rich layers in the stir zone. Nickel and copper, as austenite stabilizers, led to the formation of austenite phase in the iron-rich layers. Chrome and carbon were transferred to grain boundaries, which were high-energy zones, and formed chrome carbide. The layer-by-layer structure and precipitation at the interface made the hardness profile non-uniform and formed possible stress concentra-tion zones.AcknowledgementThe author acknowledges the funding support of Babol Noshirvani University of Technology (No. BNUT/370167/97).References[1] M. Metikoš-Hukovi ć, R. Babi ć, I. Škugor, and Z. 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Guillot, Improving frictionstir welding between copper and 304L stainless steel, Adv.Mater. Res., 409(2012), p. 263.[11] A.J. Ramirez, D.M. Benati, and H.C. Fals, Effect of tool off-set on dissimilar Cu–AISI 316 stainless steel friction stir welding, [in] Proceeding of the Twenty-first International Offshore and Polar Engineering Conference, Maui, Hawaii, USA, 2011, p. 548.[12] A. Najafkhani, K. Zangeneh-Madar, and H. Abbaszadeh,Evaluation of microstructure and mechanical properties of friction stir welded copper/316L stainless steel dissimilarmetals, Int. J. ISSI, 7(2010), No. 2, p. 21.[13] M. Shamsujjoha, B.K. Jasthi, M. West, and C. Widener, Mi-crostructure and mechanical properties of FSW lap joint be-tween pure copper and 1018 mild steel using refractory metal pin tools, [in] Friction Stir Welding and Processing VII,TMS, San Antonio, Texas, 2013, p. 151.[14] M. Jafari, M. Abbasi, D. Poursina, A. Gheysarian, and B.Bagheri, Microstructures and mechanical properties of fric-tion stir welded dissimilar steel–copper joints, J. Mech. Sci.Technol., 31(2017), No. 3, p. 1135.[15] Copper Development Association Inc., Copper–NickelWelding and Fabrication, Copper Development Association Inc., McLean, Virginia [2013-02-01]. / applications/marine/cuni/fabrication/welding_and_fabrication.html[16] Smiths Metal Centres, 304/304L Stainless Steel Data Sheet,Smiths Metal Centres, Clerkenwell, London [2007-03-05]./datasheets.htm.[17] S.H.C. Park, Y.S. Sato, H. Kokawa, K. Okamoto, S. Hirano,and M. Inagaki, Rapid formation of the sigma phase in 304 stainless steel during friction stir welding, Scripta Mater.,49(2003), No. 12, p. 1175.[18] Y.V.R.K. Prasad, K.P. Rao, and S. Sasidhara, Hot WorkingGuide: A Compendium of Processing Maps, ASM Interna-tional, Materials Park, Ohio, 2015, p. 168.[19] Y. Sun and H. 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Interfacial microstructure and mechanical properties of
Interfacial microstructure and mechanical properties of aluminium –zinc-coated steel joints made by a modifiedmetal inert gas welding –brazing processH.T.Zhang a,⁎,J.C.Feng a ,P.He a ,H.Hackl baState Key Laboratory of Advanced Welding Production Technology,Harbin Institute of Technology,Harbin 150001,Heilongjiang Province,PR ChinabFronius.Internation GMBH,A4600Wels-Thalheim,AustriaReceived 10May 2006;accepted 4July 2006AbstractThe microstructure and properties of aluminium –zinc coated steel lap joints made by a modified metal inert gas CMT welding –brazing process was investigated.It was found that the nature and the thickness of the high-hardness intermetallic compound layer which formed at the interface between the steel and the weld metal during the welding process varied with the heat inputs.From the results of tensile tests,the welding process is shown to be capable of providing sound aluminium –zinc coated steel joints.©2006Elsevier Inc.All rights reserved.Keywords:Welding –brazing;Heat input;Intermetallic compound1.IntroductionIn order to reduce pollution and save energy,it is attractive to make car bodies lighter by introducing some aluminium parts as substitutes for the previous steel structures [1,2].Therefore,joining aluminium to steel has become a major problem,requiring resolution.Direct solid-state joining can be used to make these dissimilar metal joints by controlling the thickness of the interme-tallic compound layer that develops within a few micrometers of the joint interface [3–9].However,the shape and size of such solid-state joints are extremely restricted.Thus,the joining of aluminium to steel byfusion welding methods has been widely studied.As is well known,the joining of aluminium to steel by fusion welding is difficult because of the formation of brittle interface phases which can deteriorate the mechanical properties of the joints.However,Kreimeyer and Sepold [10]have shown that if the layer is less than 10μm thick,the joint will be mechanically sound.In addition,the authors also deem that the existence of a zinc coating increases the wettability of the Al to the steel substrate.As another approach,Achar et al.[11]reported that the thickness of the intermetallic compound layer formed during TIG arc welding of Al to steel is decreased by the use of an Al alloy filler metal containing Si.Murakami et al.[12]and Mathieu et al.[13]both point out that the temperature probably determines the thickness of the intermetallic compound layer of the joint and recom-mended the use of lower heat input to obtain a sound joint.Materials Characterization 58(2007)588–592⁎Corresponding author.Tel.:+8645186412974;fax:+8645186418146.E-mail address:hitzht@ (H.T.Zhang).1044-5803/$-see front matter ©2006Elsevier Inc.All rights reserved.doi:10.1016/j.matchar.2006.07.008The cold metal transfer process,identified here as CMT,is a modified metal inert gas welding process which invented by the Fronius Company.The principal innovation of this method is that the motions of the welding wire have been integrated into the welding process and into the overall control of the process.Every time the short circuit occurs,the digital process-control both interrupts the power supply and controls the re-traction of the wire.The wire retraction motion assists droplet detachment during the short circuit,thus greatlydecreasing the heat input during welding.In this study,we selected the CMT process to join aluminium to zinc-coated steel using a lap geometry. The main purpose of this effort was to reveal the rela-tionship between heat input and the microstructure of the joint.Hardness testing was also used to characterize the phases formed during the welding process.In ad-dition,the quality of the joints was assessed by tensile testing.2.ExperimentalDeep drawn sheets of hot-dip galvanized steel and sheets of pure Al1060with thickness of1mm were used in the welding experiments.An Al sheet was lapped over a Zn-coated steel sheet on the special clamping fixture, and the ending of the weld wire was aimed at the edge of the aluminium sheet,as shown in Fig.1.The MIG welding–brazing was carried out using the CMTwelding source with an expert system and1.2-mm-diameter Al–Si filler metal wire.Argon was used as the shielding gas at a flow rate of15L/min.The surface of the samples was cleaned by acetone before welding.Two sets of welding parameters of different heat inputs were selected,as shown in Table1.The heat input,J,is calculated using the equation:J=(60×UI)/v,where U is the mean welding voltage,I is the mean welding current and v is the welding speed.Typical transverse sections of the samples were observed using optical microscopy(OM)and scanning electron microscopy(SEM).The composition of the intermetallic compound layer at the interface between the steel and the weld metal was determined by energy dispersive X-ray spectroscopy(EDX).Hardness values were obtained using a microindentation hardness tester with a load of10g,and a load time of10s.In addition, the samples were cut in10mm widths,and transverse tensile tests(perpendicular to the welding direction) were used to measure the joint tensilestrength.Fig.1.Schematic plan of the welding process.Table1The welding parametersSamplenumberMeanweldingcurrent(A)Meanweldingvoltage(V)Wire feedrate(m/min)Weldingspeed(mm/min)Weldheatinput(J/cm)Sample A6611.8 3.9762613.2Sample B11013.3 5.4913961.5Fig.2.Front(upper)and back(lower)appearances of typical jointswith different heat inputs:(a)Sample A;(b)Sample B.589H.T.Zhang et al./Materials Characterization58(2007)588–5923.Results and discussion 3.1.Macro-and microstructuresThe appearance of the weld seams for different heat inputs are shown in Fig.2.For all welding cases,a smooth weld seam was made.The molten metal wetted the steel better when using lower heat input,i.e.,compare Sample A at lower heat input to Sample B.This may be related to the different degree of evapo-ration of the zinc coating at different heat inputs.While improving the heat input,the greater evaporation of zinc reduces the wettability of the molten metal on the steel.Fig.3shows a typical cross-section of the joints.Higher heat input (Sample B)resulted in a decrease in the contact angle between the steel and the weld metal.Meanwhile,a special zone with lighter colour at the toe of the weldments can be found (designated by white arrows in Fig.3).Optical micrographs shows that a visible intermetallic compound layer has formed be-tween the steel and weld metal during the welding process,Fig.4.The thickness of the intermetallic com-pound layer changes not only with the location within a given joint but also with the varying heat input between different joints.The thickness of the intermetallic compound layer in the center is greater than at the edge of the seam within one joint.For Sample A,the maximum thickness of the compound layer is about 10μm but is 40–50μm for Sample B.The microstructure of the intermetallic compound is shown in greater detail in the SEM micrographs inFig.5.At lower heat input (Sample A),the inter-metallic compound presents a serrated shape oriented toward the weld metal.When the heat input was increased (Sample B),the compound layer became much thicker and grew into the weld metal with tongue-like penetrations.Anisotropic diffusion is a possible explanation for this irregularity.The intermetallic compounds that form under these conditions generally have an orthorhombic structure (see below).Because of the high vacancy concentration along the c -axis of the orthorhombic structure,Al atoms can diffuse rapidly in this direction and cause rapid growth of the inter-metallic compound.EDX analysis was used to determine the phases of the intermetallic compound layer.The results show that the intermetallic compound layer of the joint made by lower heat input consists entirely of Fe 2Al 5.But when the heat input is increased,the intermetallic compound layer consists of two different phases,the FeAl 2phase near the steel surface and a FeAl 3phase which penetrates toward the weld metal.Thus it is clearthatFig.4.Optical microstructures of interface between steel and weld metal:(a)Sample A;(b)SampleB.Fig.3.Cross-section image at limit of penetration in the joint,showing change in contact angle with increased heat input.Arrows point to an intermetallic compound at the tip of the weld metal:(a)Sample A;(b)Sample B.590H.T.Zhang et al./Materials Characterization 58(2007)588–592the intermetallic compound layer that forms is closely related to the heat input during the welding process.With regard to the special zone designated by white arrows in Fig.3,dendritic-appearing structures can be distinguished on a high-magnification SEM micrograph (Fig.6).EDX analysis results show that such dendrite-shaped crystals of an Al-richα-solid solution containing residual zinc routinely formed at this location.3.2.Hardness measurementsHardness testing results also confirm the presence of a hard intermetallic compound layer.The hardness of the interface layer is much higher than that of the base metal and the weld metal and is found to vary for the corresponding intermetallic compound phases.For the high heat input weld(Sample B)the hardness is much higher,Fig.7.Fig.8.The location where the fracture occurred during tensile testing (designated by white arrows):(a)Sample A;(b)SampleB.Fig.7.Microindentation hardness test results of the joints made using different heatinputs.Fig.6.Dendrite crystal structure at the toe of the weldment(SampleB).Fig.5.SEM micrograph of interface between steel and weld metal:(a)Sample A;(b)Sample B.591H.T.Zhang et al./Materials Characterization58(2007)588–5923.3.Tensile test resultsThe tensile tests were performed to provide a qualitative measure of the joint strength and behavior. These results show that the bond strength is excellent, with the fractures occurring in the HAZ of the Al even when the thickness of the intermetallic compound layer was greater than40μm,Fig.8.From a general view-point,the thickness of the intermetallic compound layer should be controlled to less than10μm in order to obtain a sound joint.This implies that the joint made with higher heat input should have a lower intrinsic strength than the other because of the thicker brittle intermetallic compound layer.However,the intrinsic strength of the joints cannot be determined when the fracture occurs in the HAZ of the pure Al.Nevertheless, according to the thickness of the compound layer,we can presume that the intrinsic strength of the joints should be decreased when increasing the welding heat input.4.ConclusionsBased on the experimental results and discussions, conclusions are drawn as follows1)Dissimilar metal joining of Al to zinc-coated steelsheet without cracking is possible by means of a modified metal inert gas(CMT)welding–brazing process in a lap joint.2)Fe–Al intermetallic compound phases were formedat the interface between the steel and the weld metal.The thickness and the composition of the interme-tallic compound layer varied with weld heat input.3)Despite the formation of the intermetallic compoundphases,the interface between steel and weld metal is not the weakest location of the joints.Tensile tests of the joints caused fractured in the Al HAZ,even when the intermetallic compound layer thickness exceeded 40μm.AcknowledgementsThe authors wish to acknowledge the financial support provided by the National Natural Science Foundation under Grant No.50325517for this work. References[1]Schubert E,Klassen M,Zerner I,Walz C,Sepold G.Light weightstructures produced by laser beam joining for future applications in automobile and aerospace industry.J Mater Process Technol 2001;115:2.[2]Schubert E,Zernet I,Sepold ser beam joining of materialcombinations for automotive applications.Proc SPIE 1997;3097:212.[3]Oikawa H,Ohmiya S,Yoshimura T.Resistance spot welding ofsteel and aluminium sheet using insert metal sheet.Sci Technol Weld Join1999;2:80.[4]Czechowski M.Stress corrosion cracking of explosion weldedsteel–aluminum joints.Mater Corros2004;6:464.[5]Fukumoto S,Tsubakino H.Friction welding process of5052aluminium alloy to304stainless steel.Mater Sci Technol 1999;9:1080.[6]Ochi H,Ogawa K,Suga Y,Iwamoto T,Yamamoto Y.Frictionwelding of aluminum alloy and steel using insert metals.Keikinzoku Yosetsu1994;11:1.[7]Shinoda T,Miyahara K,Ogawa M,Endo S.Friction welding ofaluminium and plain low carbon steel.Weld Int(UK) 2001;6:438.[8]Uzun H,Donne CD.Friction stir welding of dissimilar Al6013-T4to X5CrNi18-10stainless steel.Mater Des2005;1:41. [9]Adler L,Billy M,Quentin G.Evaluation of friction-weldedaluminum-steel bonds using dispersive guided modes of a layered substrate.J Appl Phys2001;12:6072.[10]Kreimeyer M,Sepold ser steel joined aluminium-hybridstructures.Proceedings of ICALEO'02,Jacksonville,USA;2002.[11]Achar DRG,Ruge J,Sundaresan S.Joining aluminum to steel,with particular reference to welding(III).Aluminum1980;4:291.[12]Murakami T,Nakata K.Dissimilar metal joining of aluminum tosteel by MIG arc brazing using flux cored wire.ISIJ Int 2003;10:1596.[13]Mathieu A,Mattei S,Deschamps A.Temperature control in laserbrazing of a steel/aluminium assembly using thermographic measurements.NDT&E Int2006;39:272.592H.T.Zhang et al./Materials Characterization58(2007)588–592。
Microstructure and mechanical properties
Microstructure and mechanical properties of ZrB 2–SiC nanocomposite ceramicQiang Liu,*Wenbo Han and Ping HuCenter for Composite Materials,Harbin Institute of Technology,Harbin 150001,ChinaReceived 28March 2009;accepted 30May 2009Available online 6June 2009A ZrB 2–SiC nanocomposite ceramic in which 20vol.%nanosized SiC powder was introduced into a ZrB 2matrix was fabricated by hot-pressing at 1900°C for 60min under a 30MPa uniaxed load.The composite microstructure showed intragranular nanostruc-tures that were peculiar to this material.Investigation of the mechanical properties revealed a flexural strength of 930±28MPa and a fracture toughness of 6.5±0.3MPa m 1/2.These improved mechanical properties were strongly dependent on the formation of the unusual intragranular nanostructures.Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Intragranular nanostructure;Mechanical properties;Microstructure;Fracture toughness;NanocompositeUltrahigh-temperature ceramics (UHTCs),suchas borides and carbides,were developed in the 1960s [1].Among UHTCs,zirconium diboride (ZrB 2)is a material of particular interest because of its excellent combination of high melting point,low theoretical den-sity,high electrical conductivity,good chemical inert-ness and superb wear resistance.These properties make it an attractive candidate for high-temperature applications such as refractory materials in foundries,electrical devices,nozzles and armor [2].Moreover,ZrB 2could be used for super-high-temperature struc-tural applications in aerospace [3,4].Its low mechanical properties,however,have long prevented this material from being used in a wide range of applications.Its sus-ceptibility to brittle fracture can lead to unexpected cat-astrophic failure,therefore its mechanical properties must be improved before the potential applications of ZrB 2can be fully realized.The introduction of a second phase of particles has been a successful strategy for improving the mechanical properties of monolithic diboride ceramics.With this aim,introduction of SiC particles [3–6]into ZrB 2yields a ZrB 2–SiC composite ceramic that is far stronger than monolithic ZrB 2.As a rule,however,improvement of mechanical properties is limited by the micro-sized par-ticles of the second phase.The mechanical properties of ceramics can be signif-icantly improved by introducing nanosized ceramic par-ticles into the ceramic-matrix grains or grain boundaries.The most significant achievements with this approach have been reported by Niihara and Nakahira [7–9],who first revealed that an introduction of 5vol.%of nanosized SiC particles into Al 2O 3increased the room-temperature strength of the composite from 350MPa to $1.0GPa (three-point flexure,30mm span).Similar improvements in strength have since been achieved in Al 2O 3–Si 3N 4,MgO–SiC and Si 3N 4–SiC composite systems.Materials constructed by these types of approaches are termed nanocomposite ceramics.At this point in time,however,there have been few attempts to create nanocomposite ceramics out of ZrB 2–SiC.Moreover,the effects of the composite micro-structure on the mechanical properties of ZrB 2–SiC nanocomposite ceramics have never been documented.Therefore,the aim of the present study was to investi-gate the microstructural features and effects on mechan-ical properties of a ZrB 2–SiC nanocomposite ceramic.The starting powders used in this study were:ZrB 2powder (Northwest Institute for Non-ferrous Metal Re-search,China),average particle size 2l m (>99%);and nanosized b -SiC powder (Kaier Nanotechnology Devel-opment Co.Ltd,China),average particle size 30nm (>98%).The nanosized SiC powder was first dispersed in ethanol,with 1h of ultrasonication.Then the powder mixture ZrB 2plus 20vol.%nanosized SiC particles were ball-milled using ZrO 2ball media and ethanol at1359-6462/$-see front matter Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.scriptamat.2009.05.041*Corresponding author.Tel./fax:+8645186402382;e-mail:dqz0402@Available online at Scripta Materialia 61(2009)690–692/locate/scriptamat180rpm for 12h.All ball-milling was performed in polyethylene bottles.After mixing,the resulting slurry was dried by rotary evaporation and then screened.The obtained powder mixtures were hot-pressed at 1900°C for 60min at a uniaxial pressure of 30MPa in Ar atmosphere.The microstructure of hot-pressed specimens was ob-served by using scanning electron microscopy (SEM,FEI Sirion,Holland)and transmission electron micros-copy (TEM,Hitachi H-9000,Japan)with an X-ray en-ergy dispersive spectroscopy (EDS,EDAX Inc.)analyzer attachment.Flexural strength (r )was tested in three-point bending on 3Â4Â36mm bars,using a 30-mm span and a crosshead speed of 0.5mm min À1.Each specimen was ground and polished with diamond slurries to a 1-l m finish.The edges of all the specimens were chamfered to minimize the effect of stress concen-tration resulting from machining flaws.Fracture tough-ness (K IC )was evaluated by a single-edge notched-beam test with a 16-mm span and a crosshead speed of 0.05mm min À1using 2Â4Â22mm test bars on the same jig used for the flexural strength.All flexural bars were fabricated with the tensile surface perpendicular to the hot-pressing direction.A minimum of five specimens was tested in each experimental condition.Figure 1shows the typical microstructural morphol-ogies of the ZrB 2–SiC nanocomposite ceramic under SEM (Fig.1a)and TEM (Fig.1b).As shown in Figure 1a,a number of submicron SiC particles (dark contrast)are located along the grain boundaries of the ZrB 2(gray contrast).Some smaller SiC particles also appear inside the ZrB 2grains (indicated by arrows);these are termed intragranular nanostructures.Higher magnification examination of the ZrB 2–SiC nanocomposite ceramic by TEM (Fig.1b)showed that the intragranular SiC particles (indicated by arrows)were approximately 100nm in size.The formation of the intragranular nanostructures was dependent on both the migration speed of ZrB 2ma-trix grain boundary and the migration speed of the SiC second phase [10,11].When the migration speed of the second phase was less than that of the matrix grain boundary,the nanosized SiC particles tended to be trapped within the ZrB 2grains during sintering.The fine ZrB 2particles would then coalesce around them,form-ing the intragranular nanostructures.Figure 2shows that the specimen fracture surface used for testing fracture toughness exhibited the typicalcharacteristics of a transgranular fracture.In monolithic ZrB 2ceramic,the predominant fracture mode would have been intergranular [12].There are two possible interpretations for this difference in fracture mode.The first is that the intergranular SiC particles in the ZrB 2–SiC nanocomposite ceramic were firmly bonded to the ZrB 2/ZrB 2interfaces.This rigid bonding could then have suppressed intergranular fracture [13].The other explanation is that there are differences in relaxation of the tensile residual stress around the SiC particles located between the intergranular and intra-granular.Because of the different thermal expansion coefficients between SiC and ZrB 2,a large internal stress will be generated during cooling after sintering.Assum-ing that a SiC particle is spherical,an internal tension will occur in a tangential direction to the ZrB 2matrix around the SiC particle.This will cause a crack to al-ways propagate towards the SiC particle.The internal tangential tension also would be relaxed by lattice and grain-boundary diffusion around the intragranular and intergranular particles,respectively.However,the tem-perature at which the grain-boundary diffusion is acti-vated would be lower than that required by lattice diffusion,thus the internal tangential tension around the intergranular SiC particles would be further relaxed during cooling.As a result,the internal tangential ten-sion around the intragranular SiC particles of the sin-tered body would always be greater than that around the intergranular particles.This would lead to a fracture surface that would always tend to be characteristic of a transgranular fracture.Thus,it is the intragranular nanostructures that predominantly induce the trans-granular fracture characteristic of the ZrB 2–SiC nano-composite ceramic.Examination of the mechanical properties of the ZrB 2–SiC nanocomposite ceramic revealed a fracture toughness that ranged from 6.4to 6.7MPa m 1/2.This represented an increase of approximately 83%over that of the monolithic ZrB 2(2.3–3.5MPa m 1/2)[2].In addi-tion,the flexural strength (920–945MPa)of this nano-composite ceramic was also significantly higher than that recently reported for the monolithic ZrB 2($565MPa)[4].The formation of the intragranular nanostructures appeared to play an important role in the improved mechanical properties of the ZrB 2–SiC nanocomposite ceramic,especially its increased fracture toughness and flexural strength.In order to investigate effects of the intragranular nanostructures on the mechanical properties oftheFigure 1.Typical microstructural morphologies of the ZrB 2–SiC nanocomposite ceramic:(a)SEM image of the sample and (b)TEM image of thesample.Figure 2.SEM image of the fracture surface of the ZrB 2–SiC nanocomposite ceramic.Q.Liu et al./Scripta Materialia 61(2009)690–692691ZrB 2–SiC nanocomposite ceramic,it is necessary to investigate a crack propagation behavior in this mate-rial.Figure 3shows TEM micrographs of crack propa-gation behavior in the ZrB 2–SiC nanocomposite ceramic.It was evident that the crack had never propa-gated in a straight line,but had been deflected,selecting the neighboring particles (Fig.3a).As stated previously,this deflection was caused by thermal internal stress in this material.It can be also seen in Figure 3a that a crack has penetrated through an intragranular particle (indicated by black arrow).The possible reason for this case is that the cracked particle may be an agglomera-tion composed of many fine SiC particles.Because the bond strength of this agglomeration is not high enough,it tends to fracture when a crack meets this kind of par-ticle.However,for other intragranular particles (<100nm),neither crack penetration through the intra-granular particles nor propagation along the particle/matrix interfaces was evident (Fig.3b).This phenome-non indicates that the intragranular particles bridged the crack,pointing to the existence of a particle-bridging mechanism.Based on the experimental observation above,a spe-cific explanation for this effect is as follows.When a pri-mary crack meets an intragranular nanosized SiC particle,it is normally impeded and thus bows (Fig.3a).The bowing crack bypasses the impenetrable particles and instead interacts with neighboring cracks.At this point,the bridging particles firmly pin the cracks and further prevent the crack from extending.As a re-sult,only by increasing the crack extension force can the crack further extend.In other words,it is by means of the particle-bridging mechanism that the strength and toughness of the ZrB 2nanocomposite ceramic are signif-icantly improved.Besides the explanation mentioned above,there is an-other one for the improvement in strength.After the for-mation of the intragranular nanostructures,there are many sub-interfaces within the ZrB 2matrix grains that belong to the interfaces between intragranular particles and matrix grains.As stated previously,moreover,be-cause of the difference in thermal expansion coefficients between the ZrB 2matrix and the SiC second phase,a large number of microcracks were formed around the intragranular particles,as shown in Figure 4.The for-mation of the sub-interfaces and microcracks can cause the matrix grains to be at a potential differentiation state,corresponding to the further grain refining.Thisthen improves the strength of this material according to the Hall–Petch equation [10].As discussed above,it is concluded that the formation of intragranular nanostructure is the fundamental rea-son for the significant increase in the mechanical proper-ties of this nanocomposite ceramic.In conclusion,a hot-pressed ZrB 2–SiC nanocompos-ite ceramic was fabricated by introducing nanosized SiC powder into a ZrB 2matrix.the intragranular nanostruc-tures were peculiar to this ceramic-based composite and induced a transgranular fracture characteristic.The mechanical properties of this nanocomposite ceramic,especially its flexural strength and fracture toughness,were much higher than those of monolithic ZrB 2.It is believed that the formation of intragranular nanostructures is a main reason for the improvements in mechanical properties of the ZrB 2–SiC nanocompos-ite ceramic.Intragranular particle bridging is believed to be the predominant toughening mechanism imparting the improved characteristics to this material.This work was supported by the NSFC(10725207),the Research Fund for the Doctoral Pro-gram of Higher Education (24403037)and National Natural Science Fund for Outstanding Youths (24402052).[1]E.V.Clougherty,R.L.Pober,L.Kaufman,Trans.Met.Soc.AIME 242(1968)1077.[2]F.Monteverde,S.Guicciardi,A.Bellosi,Mater.Sci.Eng.A 346(2003)310.[3]A.L.Chamberlain,W.G.Fahrenholtz,G.E.Hilmas,D.T.Ellerby,J.Am.Ceram.Soc.87(2004)170.[4]F.Monteverde,C.Melandri,S.Gicciardi,Mater.Chem.Phys.100(2006)513.[5]F.Monteverde,Appl.Phys.A 82(2006)329.[6]S.S.Hwang,A.L.Vasiliev,N.P.Padture,Mater.Sci.Eng.A 464(2007)216.[7]K.Niihara, A.Nakahira,in:P.Vincentini (Ed.),Advanced Structural Inorganic Composites,Elsevier Sci-ence Publishers,Trieste,Italy,1990,pp.637–664.[8]K.Niihara,A.Nakahira,Ann.Chim.16(1991)479.[9]K.Niihara,J.Ceram.Soc.Jpn.99(1991)974.[10]W.D.Kingery,H.K.Bowen,D.R.Uhlmann,Introduc-tion to Ceramics,Wiley,1976.[11]C.M.Wang,J.Mater.Sci.30(1995)3222.[12]S.Q.Guo,J.M.Yang,H.Tanaka,Y.Kagawa,Compos.Sci.Technol.68(2008)3033.[13]I.A.Ovid’ko,A.G.Sheinerman,Scripta Mater.60(2009)627.Figure 3.TEM micrographs of crack propagation behavior in the ZrB 2–SiC nanocomposite ceramic:crack propagation is from upper right to lowerleft.Figure 4.TEM micrograph of microcracks around an intragranular particle.692Q.Liu et al./Scripta Materialia 61(2009)690–692。
Si含量对Al-6.5Cu-0.6Mn-0.5Fe合金组织演变及力学性能的影响
Trans.Nonferrous Met.Soc.China29(2019)1583−1591Microstructure evolution and mechanical properties ofAl−6.5Cu−0.6Mn−0.5Fe alloys with different Si additionsRui XU1,Bo LIN1,Hao-yu LI1,Hua-qiang XIAO1,Yu-liang ZHAO2,Wei-wen ZHANG31.School of Mechanical Engineering,Guizhou University,Guiyang550025,China;2.School of Mechanical Engineering,Dongguan University of Technology,Dongguan523808,China;3.School of Mechanical and Automotive Engineering,South China University of Technology,Guangzhou510640,ChinaReceived13November2018;accepted5June2019Abstract:The effect of Si content on the microstructures and mechanical properties of the heat-treated Al−6.5Cu−0.6Mn−0.5Fe alloy was investigated using image analysis,scanning electron microscopy(SEM),transmission electron microscopy(TEM),and tensile testing.The results show that the mechanical properties of Al−6.5Cu−0.6Mn−0.5Fe alloys decrease slightly when the Si content is below1.0%.This can be attributed to the comprehensive effect of microstructure evolution,including the increase of nano-sizedα-Fe,the coarsened grain size,and an increase in Al2Cu content at the grain boundary.When the Si content is1.5%,the mechanical properties of the Al−6.5Cu−0.6Mn−0.5Fe alloys decrease significantly,and this can be attributed to the agglomerated second intermetallics,which is resulted from the formation of excess Si particles.Key words:Al−Cu alloys;iron-rich intermetallics;Si;tensile properties1IntroductionWith the increasing requirements for environmental protection and green manufacturing,recycling Al alloys has become an important direction for development in the aluminum industry.However,recycled Al alloys contain many impurity elements,such as Fe,Si,Ni,Zn, Mg and Mn[1,2].Iron is the most harmful impurity element in recycled aluminum alloys.Because the solid solubility of Fe in Al alloys is limited,Fe usually precipitates in the form of hard and brittle Fe-rich intermetallics,such as Chinese scriptα-Fe (Al15(FeMn)3(SiCu)2)[3],Al6(FeMn)[4],Al m Fe[5], plate-like Al3(FeMn)[6],andβ-Fe(Al7Cu2Fe)[7]. Several methods have been effectively used to prevent the detrimental effects of iron in aluminum alloys.These methods include(1)reducing the formation of Fe-rich intermetallics by lowering the Fe levels to be as low as economically possible,and(2)modifying the Fe-rich intermetallics using chemical or physical approaches. In the chemical approaches,Mn and Si elements are added to break up the needle-like intermetallics or to transform them into block or Chinese script.The physical approaches include the use of superheated melt,solidification under high cooling rate,and melt treatment[8−11].Al−Cu alloys are widely used in aerospace and automobile manufacturing industry because of their excellent fatigue properties,high specific strength,and good heat resistance.However,Al−Cu alloys have poor casting properties,such as hot cracking susceptibility and fluidity in cast aluminum alloys[12].It has been found that Si and Cu alloying of Al−Cu alloys decreases hot cracking susceptibility[13−15].SABAU et al[15]found that Al−Cu alloys have the lowest thermal cracking tendency when the copper content is7.3%and8%. In our previous work,we found that nano-sizedα-Fe can be formed in Al−Cu alloys that have high Cu and Fe content,and this significantly improves the mechanicalFoundation item:Projects(51704084,51605106)supported by the National Natural Science Foundation of China;Project(2017M623068)supported by China Postdoctoral Science Foundation;Project(2015A030312003)supported by the Natural Science Foundation for Team Research ofGuangdong Province,China;Project(JC(2016)1026))supported by the Science and Technology Foundation of Guizhou Province ofChina;Project(KY(2017)101))supported by the Young Talent Growth Foundation of Education Department of Guizhou Province ofChina;Project(RC2017(5788))supported by the Science and Technology Plan of Guizhou Province of ChinaCorresponding author:Bo LIN;Tel/Fax:+86-851-83627516;E-mail:linbo1234@DOI:10.1016/S1003-6326(19)65065-XRui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−1591 1584properties[16].The nano-sizedα-Fe phase is commonly observed in3xxx and6xxx alloys with high Si content[17,18].Therefore,it can be expected that adding Si will promote the formation of the nano-sizedα-Fe phase in Al−Cu alloys.Also,Si content has a significant effect on the formation of Fe-rich intermetallics. However,Si is also an impurity element in Al−Cu alloys. As a result,the Si content is usually limited to below 0.1%for high-performance Al−Cu alloys[19].However, strictly controlling the Si content in aluminum alloys will make these alloys expensive and is not conducive to the recycling of waste aluminum.Therefore,the effects of different Si contents on the microstructure and properties of the Al−6.5Cu−0.6Mn−0.5Fe alloy were studied for the purpose of promoting the efficient use of recycled aluminum.2ExperimentalExperimental alloys with different Si contents were prepared using commercially pure Al(99.5%), Al−50%Cu,Al−10%Mn,Al−20%Si and Al−5%Fe master alloys.The compositions of the alloys were determined using optical emission spectrometry,and the results are listed in Table1.First,pure Al,Al−5%Fe and Al−20%Si master alloys were preheated in a clay-graphite crucible using an electric resistance furnace at400°C for1h to eliminate water vapor.The raw materials were then melted at730°C.Al−50%Cu and Al−10%Mn master alloys were added at730°C.Finally, the temperature of the melt was maintained at730°C for 30min.Approximately10kg of the melts were degassed using argon to minimize the hydrogen content.The melt was then poured into a cylindrical die,which had a size of80mm in height and50mm in diameter.The die temperature was set to be250°C,and the pouring temperature was730°C.T7heat treatment conditions were used in this study to stabilize the microstructures. The samples were then solution-treated at535°C for 12h before being quenched in warm water at100°C. The samples were then aged using T7conditions at 215°C for16h.For mechanical tests,samples with a diameter of10mm and a height of80mm were produced using a wire electrical discharge machining. The tensile test was carried out on an MTS CMT5105 standard testing machine,and the reported values are the averages of at least three samples.Samples for the micro-hardness test and metallographic observation were cut in the gauge length part from selected tensile specimens.The location for the micro-hardness test was restricted to the center of theα(Al)dendrite near the center of the etched specimens.The micro-hardness was measured on a tester equipped with a Vickers diamond indenter using an indentation load of100g.The value reported is the average of more than10readings.The samples used for metallographic observations were etched in0.5%HF solution for30s.The morphology, the Fe-rich phase,and the fracture surfaces were analyzed using SEM(Nova Nano SEM430)attached with EDS.Precipitates in theα(Al)matrix were analyzed using TEM(JEOL JEM−3010)at200kV.The area fraction of the intermetallics and precipitations of the alloys were quantitatively calculated by using image Pro Plus software.Table1Chemical composition of alloy(wt.%)Alloy Cu Mn Fe Si Al Al−6.5Cu−0.6Mn−0.5Fe−0Si(0Si)6.450.660.550.04Bal. Al−6.5Cu−0.6Mn−0.5Fe−0.5Si(0.5Si)6.470.620.540.58Bal. Al−6.5Cu−0.6Mn−0.5Fe−1.0Si(1.0Si)6.500.630.53 1.13Bal. Al−6.5Cu−0.6Mn−0.5Fe−1.5Si(1.5Si)6.520.630.56 1.50Bal. 3Results3.1Microstructure of as-cast alloysFigure1shows the microstructure of the as-cast Al−6.5Cu−0.6Mn−0.5Fe alloy with different Si contents. Table2gives EDS data of the Fe-rich phase in the as-cast Al−6.5Cu−0.6Mn−0.5Fe alloy with different Si contents.The Fe-rich phase shows Chinese script morphology in all of the experimental alloys.As seen in Table2,the Chinese script Fe-rich intermetallic isα-Fe (Al15(FeMn)3Cu2or Al15(FeMn)3(CuSi)2)[3,9].In addition,the amount of Si inα-Fe increases with an increase in Si content,and this is because the Al atom is substituted by the Si atom in the Fe-rich phase[20].On the basis of Fig.1,the area fraction ofα-Fe is quantitatively calculated.The area fraction ofα-Fe increases from3.9%to4.3%and then to4.5%when the Si content is increased from0to0.5%and then to1.0%. The results show that the area fraction ofα-Fe increases with an increase in the Si content,and this is similar to the results reported in the work by LIU et al[10].3.2Microstructure of heat-treated alloysFigure2shows the microstructure of the heat-treated Al−6.5Cu−0.6Mn−0.5Fe alloy with different Si contents.Table3gives EDS data of the Fe-rich phase in the heat-treated Al−6.5Cu−0.6Mn−0.5Fe alloy with different Si pared with the as-cast alloy, Al2Cu is completely dissolved into theα(Al)matrix after the T7heat treatment in the0Si alloy.Also,Cu-richβ-Fe transforms fromα-Fe after solution heat treatment in the0Si alloy.The amount of Al2Cu increases,andα-FeRui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−15911585Fig.1Microstructures of as-cast Al−6.5Cu−0.6Mn−0.5Fe alloys with different Si contents:(a)0Si;(b)0.5Si;(c)1.0Si;(d)1.5Si Table2Compositions of Fe-rich intermetallics in as-cast condition(at.%)Alloy Phase Al Cu Mn Fe Si0Siα-Fe78.99±2.068.62±2.05 4.22±0.618.17±1.92−0.5Siα-Fe77.71±0.72 3.59±0.37 3.41±0.247.00±0.448.28±0.191.0Siα-Fe74.70±2.42 2.67±0.29 5.48±1.048.08±0.939.07±1.561.5Siα-Fe73.30±1.92 4.25±2.48 4.50±2.0910.55±1.927.41±0.88Fig.2Microstructures of heat-treated Al−6.5Cu−0.6Mn−0.5Fe alloys with different Si contents:(a)0Si;(b)0.5Si;(c)1.0Si;(d)1.5SiRui XU,et al/Trans.Nonferrous Met.Soc.China 29(2019)1583−15911586remains with an increase of the Si content.With a further increase in the amount of added Si to 1.5%,Si particles appear in the alloys.Also,excessive Si particles agglomerate with the β-Fe phase and Al 2Cu phase at the grain boundary.Figure 3shows the effects of different Si contents on the grain size of the Al−6.5Cu−0.6Mn−0.5Fe alloy.The grain size of the alloy increases with an increase in the Si content.The grain sizes of alloys with different Si contents are quantitatively analyzed,and the results are shown in Fig.3(f).The grain size of the 0Si alloy is estimated to be 540μm,much less than that of the 1.5Si alloy (2350μm).Figure 4shows the morphology and EDS data of the second phase in the α(Al)matrix with different Si contents.According to EDS result of the 0.5Si alloy,the second phase precipitations in the α(Al)matrix are T phase and α-Fe.Furthermore,the amount of T phaseTable 3Compositions of Fe-rich intermetallics in heat-treated condition (at.%)Alloy Phase Al Cu Mn Fe Si 0Si β-Fe 73.82±2.5117.78±1.80 2.09±0.02 6.32±0.70−0.5Si α-Fe 74.15±0.75 4.52±0.43 4.66±0.198.13±1.298.54±0.191.0Si α-Fe 70.25±2.91 4.07±0.22 5.57±0.519.17±1.4610.94±2.001.5Siα-Fe71.47±1.543.23±0.636.12±0.817.64±0.6811.54±1.23Fig.3Macro-etched cross-section morphology of samples (a),micro-images of grain size for 0Si (b),0.5Si (c),1.0Si (d)and 1.5Si (e)alloys,and quantitatively-analyzed grain sizes (f)Rui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−15911587 Fig.4TEM images(a,b)and EDS results(c,d)of alloys with different Si contents:(a,c)0Si;(b,d)0.5Sidecreases and that ofα-Fe increases in alloys with an increase in Si content.This indicates that adding Si promotes the formation of theα-Fe phase in theα(Al) matrix.These results are also consistent with the previous finding that adding Si promotes the formation of theα-Fe phase during T7heat treatment in the3xxx and6xxx alloys[17,18].On the basis of Fig.4,the area fractions of T phases andα-Fe phases are quantitatively calculated.The area fractions of T phases andα-Fe phases in the0Si alloy are12.6%and1.3%,respectively. And,the area fractions of T phases andα-Fe phases in the0.5Si alloy are5.9%and10.2%,respectively.These results further prove that the Si content promotes the formation of theα-Fe phase in theα(Al)matrix.Figure5shows micro-hardness of alloys with different Si contents.Adding Si obviously increases the micro-hardness of the matrix.The micro-hardness of the 0Si alloy(about HV92.7)is much lower than that of the 1.0Si alloy(about HV110.5).These results can be attributed to the increase in the amount of nano-sized α-Fe phase with Si addition.3.3Mechanical properties of alloysFigure6shows the mechanical properties of the heat-treated Al−6.5Cu−0.6Mn−0.5Fe alloy with different Si contents.The strength and elongation decrease as the Si content is increased.The mechanical properties of the Al−6.5Cu−0.6Mn−0.5Fe alloys decrease slightly when the Si content is below1.0%.The ultimatetensile Fig.5Micro-hardness of alloy with different Sicontents Fig.6Mechanical properties of Al−6.5Cu−0.6Mn−0.5Fe alloys with different Si contentsRui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−1591 1588strength(UTS),yield strength(YS),and elongation of the heat-treated1.0Si alloy are recorded as280MPa, 253MPa,and2.5%,respectively.When the Si content is 1.5%,the mechanical properties of the Al−6.5Cu−0.6Mn−0.5Fe alloys decrease significantly,and this can be attributed to the precipitation of Si particles and the agglomeration of the coarse second phase.3.4Fracture surfaces of alloysFigure7shows the fracture surfaces of alloys with different Si contents.There are some dimples on the fracture surfaces,and this indicates that the fractures of alloy exhibit certain ductile fracture characteristics. When the Si content is less than1.0%,the alloy fracture surfaces are not significantly different,but when the Si content is1.5%,brittle fracture characteristics become obvious.In addition,the morphology of the second phase of the alloys with different Si contents is obviously different on the fracture surfaces.When the Si content is 0%,there are some cracks in the cylindricalβ-Fe (Al7Cu2Fe),and this indicates thatβ-Fe acts as crack initiation sites and leads to quasi-cleavage fracture (Fig.7(a)).With a further increase in the Si content, cracks are also found in theα-Fe and Al2Cu phases on the fracture surfaces,and this indicates that the brittle α-Fe and Al2Cu act as potential cleavage initiators (Figs.7(b,c)).However,the amount of cracks inα-Fe is less than that inβ-Fe,and this indicates thatα-Fe is less harmful than the cylindricalβ-Fe[21].In the1.5Si alloys,the agglomerated second intermetallic phases are clearly seen when the fracture surfaces undergo cleavage fracture(Fig.7(d)).The fracture surfaces clearly show that the plasticity of the alloy does not obviously change when the Si content is below 1.0%.However,the plasticity of the alloy decreases sharply with an increase in the Si content to1.5%.Figure8shows the longitudinal fracture morphology of alloys with different Si contents.When the Si content is0%,there are manyβ-Fe phases on the fracture surface of the alloy.This indicates that the brittle β-Fe acts as potential cleavage initiators(Fig.8(a)).With a further increase in the Si content(below1.0%),the α-Fe particles and the Al2Cu phase are observed at or beneath the fracture surface in the alloy,and this means that theα-Fe particles and Al2Cu act as potential cleavage initiators(Figs.8(b,c)).A large amount of agglomerated second intermetallics precipitate on the fracture surface of the 1.5Si alloy,and a crack of agglomerated second intermetallics results in the formation of secondary cracks(Fig.8(d)).4DiscussionFigure6shows that the UTS,YS and elongation of the heat-treated alloy decrease with an increase in Si content,and this is mainly because of microstructure evolution.First,adding Si promotes the transformation from T(Al20Cu2Mn3)phase to nano-sizedα-Fe after T7heat treatment.This phenomenon can be explained by theSi Fig.7Fractographs of alloys with different Si contents:(a)0Si;(b)0.5Si;(c)1.0Si;(d)1.5SiRui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−15911589 Fig.8Longitudinal fracture morphologies of alloys with different Si contents:(a)0Si;(b)0.5Si;(c)1.0Si;(d)1.5Siatom replacing the Al atom in the Fe-rich phase.As a result,adding Si promotes the formation of theα-Fe phase[17,18].This result is consistent with the previous finding that an increase in Si content can promote the formation ofα-Fe phase in the heat-treated3xxx and 6xxx aluminum alloys.It has been reported that the formation of nano-sizedα-Fe phase is more beneficial than the large T phase for the properties of alloys[22].In addition,Si content obviously increases the grain size of alloys(Fig.3).These results can be attributed to the decreased amount of T phases in the matrix with an increase in Si content.Many researchers have reported that the T phases prevent grain growth of Al−Cu alloys[8,23].Finally,an increase in the Si content stabilizesα-Fe and inhibits the transformation ofα-Fe toβ-Fe during heat treatment.In Si-free alloys,α-Fe transforms into copper-richβ-Fe((Al15(FeMn)3Cu2+4Al2Cu→3Al7Cu2(FeMn)+2α(Al))[20,24,25].However,α-Fe remains stable and reduces Cu consumption in alloys with high Si content,and this leads to an increase in the Al2Cu phase at the grain boundaries.The brittle Al2Cu phase acts as a potential cleavage initiator,and this indicates that an increase in the Al2Cu phase in alloys with high Si content deteriorates the performance of the alloy.In conclusion,the decrease in mechanical properties can be attributed to the comprehensive effect of microstructure evolution,including an increase in nano-sizedα-Fe,the coarsened grain size,and an increase in Al2Cu at the grain boundary.When the Si content is1.5%,the alloy properties decrease sharply because of excess Si particles and the agglomeration of brittle second intermetallics at the grain boundaries. These results demonstrate that the Si content should be appropriately controlled in the Al−Cu alloys.In this study,acceptable mechanical properties are achieved by controlling the Si content to be below1.0%.Moreover,it is feasible to extend the Fe and Si contents for the purpose of using recycled aluminum alloys,which greatly reduces manufacturing costs.Also,an increase in nano-sizedα-Fe is beneficial to improving elevated mechanical properties[26−28].Also,an increase in Al2Cu at the grain boundary is beneficial to improving elevated mechanical properties[29].Therefore,it can be expected that the Al−6.5Cu−0.6Mn−0.5Fe alloy with high Si content will be beneficial to the mechanical properties at elevated temperature.Further work will be carried out to evaluate the mechanical properties at elevated temperature for industrial applications of Al−6.5Cu−0.6Mn−0.5Fe with high Si content.5Conclusions(1)An increase in Si content in the Al−6.5Cu−0.6Mn−0.5Fe alloy promotes the transformation of the T(Al20Cu2Mn3)phase to nano-sizedα-Fe in the matrix.A decrease in the T(Al20Cu2Mn3)phase in alloys with high Si content weakens the inhibition ability of grain growth, and this leads to coarsening of grain size.(2)Adding Si stabilizesα-Fe and inhibits the transformation fromα-Fe toβ-Fe after heat treatment, and this leads to an increase in the Al2Cu phase in the alloy.Rui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−1591 1590(3)Mechanical properties of the Al−6.5Cu−0.6Mn−0.5Fe alloy decrease with an increase in Si content.Mechanical properties of Al−6.5Cu−0.6Mn−0.5Fe alloys decrease slightly when the Si content is below 1.0%.However,the mechanical properties of Al−6.5Cu−0.6Mn−0.5Fe−1.5Si alloys decrease significantly.This can be attributed to the agglomerated second intermetallic that is resulted from the formation of excess Si particles.References[1]BELOV N A,ESKIN D G,AKSENOV A A.Iron in aluminiumalloys:Impurity and alloying element[M].Florida,USA:CRC Press, 2014.[2]ZHANG L F,GAO J W,DAMOAH L N W,ROBERTSON D G.Removal of iron from aluminum:A review[J].Mineral Processing and Extractive Metallurgy Review,2012,33(2):99−157.[3]LIU K,CAO X,CHEN X G.Solidification of iron-rich intermetallicphases in Al−4.5Cu−0.3Fe cast alloy[J].Metallurgical and Materials Transactions A,2011,42(7):2004−2016.[4]ZHANG W W,LIN B,LUO Z,ZHAO Y L,LI Y Y.Formation ofFe-rich intermetallic compounds and their effect on the tensile properties of squeeze-cast Al−Cu alloys[J].Journal of Materials Research,2015,30(16):2474−2484.[5]LIU K,CAO X,CHEN X G.A new iron-rich intermetallic−Al m Fephase in Al−4.6Cu−0.5Fe cast alloy[J].Metallurgical and Materials Transactions A,2012,43(4):1097−1101.[6]LIU K,CAO X,CHEN X G.Precipitation of iron-rich intermetallicphases in Al−4.6Cu−0.5Fe−0.5Mn cast alloy[J].Journal of Materials Science,2012,47(10):4290−4298.[7]TSENG C J,LEE S L,WU T F,LIN J C.Effects of Fe content onmicrostructure and mechanical properties of A206alloy[J].Materials Transactions JIM,2000,41(6):708−713.[8]TSENG C J,LEE S L,TSAI S C,CHENG C J.Effects of manganeseon microstructure and mechanical properties of A206alloys containing iron[J].Journal of Materials Research,2002,17(9): 2243−2250.[9]KAMGA H K,LAROUCHE D,BOURNANE M,RAHEM A.Solidification of aluminum−copper B206alloys with iron and silicon additions[J].Metallurgical and Materials Transactions A,2010, 41(11):2844−2855.[10]LIU K,CAO X,CHEN X G.Effect of Mn,Si,and cooling rate onthe formation of iron-rich intermetallics in206Al−Cu cast alloys[J].Metallurgical and Materials Transactions B,2012,43(5):1231−1240.[11]WANG Q L,GENG H R,ZHANG S,JIANG H W,ZUO M.Effectsof melt thermal-rate treatment on Fe-containing phases in hypereutectic Al−Si alloy[J].Metallurgical and Materials Transactions A,2014,45(3):1621−1630.[12]WANG L,WANG N,PROVATAS N.Liquid channel segregation andmorphology and their relation with hot cracking susceptibility during columnar growth in binary alloys[J].Acta Materialia,2017,126: 302−312.[13]KAMGA H K,LAROUCHE D,BOURNANE M,RAHEM A.Hottearing of aluminum–copper B206alloys with iron and silicon additions[J].Materials Science and Engineering A,2010, 527(27−28):7413−7423.[14]KANG B K,SOHN I.Effects of cu and Si contents on the fluidity,hot tearing,and mechanical properties of Al−Cu−Si alloys[J].Metallurgical and Materials Transactions A,2018,49(10): 5137−5145.[15]SABAU A S,MIRMIRAN S,GLASPIE C,LI S,APELIAN D,SHYAM A,HAYNES J A,RODRIGUEZ A F.Hot-tearing assessment of multicomponent nongrain-refined Al−Cu alloys for permanent mold castings based on load measurements in a constrained mold[J].Metallurgical and Materials Transactions B, 2018,49(3):1267−1287.[16]LIN B,XU R,LI H Y,XIAO H Q,ZHANG W W,LI S B.Development of high Fe content squeeze cast2A16wrought Al alloys with enhanced mechanical properties at room and elevated temperatures[J].Materials Characterization,2018,142:389−397. [17]ALEXANDER D T L,GREER A L.Solid-state intermetallic phasetranformations in3XXX aluminium alloys[J].Acta Materialia,2002, 50(10):2571−2583.[18]BIROL Y.The effect of homogenization practice on themicrostructure of AA6063billets[J].Journal of Materials Processing Technology,2004,148(2):250−258.[19]GREEN J A S.Aluminum recycling and processing for energyconservation and sustainability[M].Ohio,USA:ASM International, 2007.[20]KAMGA H K,LAROUCHE D,BOURNANE M,RAHEM A.Mechanical properties of aluminium–copper B206alloys with iron and silicon additions[J].International Journal of Cast Metals Research,2012,25(1):15−25.[21]LIU K,CAO X,CHEN X G.Tensile properties of Al−Cu206castalloys with various iron contents[J].Metallurgical and Materials Transactions A,2014,45(5):2498−2507.[22]LIN B,ZHANG W W.Evolution of iron-rich intermetallics andelevated temperature mechanical properties in gravity die cast2A16 Al alloy[J].International Journal of Cast Metals Research,2018, 31(4):222−229.[23]FENG Z Q,YANG Y Q,HUANG B,LI M H,CHEN Y X,RU J G.Crystal substructures of the rotation-twinned T(Al20Cu2Mn3)phase in2024aluminum alloy[J].Journal of Alloys and Compounds,2014, 583:445−451.[24]LIU K,CAO X,CHEN X G.Solid-state transformation of iron-richintermetallic phases in Al−Cu206cast alloys during solution heat treatment[J].Metallurgical and Materials Transactions A,2013, 44(8):3494−3503.[25]LIN B,ZHANG W W,ZHAO Y L,LI Y Y.Solid-statetransformation of Fe-rich intermetallic phases in Al−5.0Cu−0.6Mn squeeze cast alloy with variable Fe contents during solution heat treatment[J].Materials Characterization,2015,104:124−131. [26]LIU K,CHEN X G.Improvement in elevated-temperature propertiesof Al−13%Si piston alloys by dispersoid strengthening via Mn addition[J].Journal of Materials Research,2018:33(20): 3430−3438.[27]SHAHA S K,CZERWINSKI F,KASPRZAK W,FRIEDMAN J,CHEN D L.Ageing characteristics and high-temperature tensile properties of Al−Si−Cu−Mg alloys with micro-additions of Mo and Mn[J].Materials Science and Engineering A,2017,684:726−736.[28]FARKOOSH A R,CHEN X G,PEKGULERYUZ M.Interactionbetween molybdenum and manganese to form effective dispersoids in an Al−Si−Cu−Mg alloy and their influence on creep resistance[J].Materials Science and Engineering A,2015,627:127−138.[29]WANG E R,HUI X D,CHEN G L.Eutectic Al−Si−Cu−Fe−Mnalloys with enhanced mechanical properties at room and elevated temperature[J].Materials&Design,2011,32(8−9):4333−4340.Rui XU,et al/Trans.Nonferrous Met.Soc.China29(2019)1583−15911591 Si含量对Al−6.5Cu−0.6Mn−0.5Fe合金组织演变及力学性能的影响许锐1，林波1，李浩宇1，肖华强1，赵俞亮2，张卫文31.贵州大学机械工程学院，贵阳550025；2.东莞理工学院机械工程学院，东莞523808；3.华南理工大学机械与汽车工程学院，广州510640摘要：采用定量分析、扫描电镜、透射电镜和拉伸性能测试研究Si含量对高铁含量铝铜合金组织演变及力学性能的影响。
基于高通量的原位制备网状结构TiC增强TC4复合材料的组织与性能
第27卷第1期粉末冶金材料科学与工程2022年2月V ol.27 No.1 Materials Science and Engineering of Powder Metallurgy Feb. 2022DOI:10.19976/ki.43-1448/TF.2021070基于高通量的原位制备网状结构TiC增强TC4复合材料的组织与性能杜康鸿，柳中强，张建涛，温利平，肖志瑜(华南理工大学国家金属近净成形工程技术研究中心，广州510640)摘要：以不同粒度的TC4合金粉末为基体材料，以VC作为碳源，采用高通量热压烧结工艺，原位制备具有不同网状结构尺寸和不同TiC体积分数(分别为2%、4%和6%)的TiC/TC4钛基复合材料，研究TiC含量和TC4粉末粒度对复合材料组织与性能的影响。
结果表明，TiC/TC4复合材料中的TiC增强颗粒呈网状分布。
与TC4合金相比，TiC/TC4复合材料的组织明显细化。
随TiC含量增加，TiC网状结构的厚度增大，材料的抗拉强度与伸长率先升高后下降，TiC含量为2%的复合材料综合性能最优。
随TC4粉末粒度减小，TiC/TC4复合材料中的基体组织逐渐细化，基体的连通性提高，材料抗拉强度与伸长率同时提高。
采用粒度为40～80 μm的TC4合金粉末为原料制备的2%TiC/TC4复合材料，网状结构尺寸小，综合性能最优，屈服强度、抗拉强度和伸长率分别为946 MPa、1058 MPa和18.1%，较TC4合金分别提高29.6%、31.6%和118.1%。
关键词：钛基复合材料；原位碳化钛；高通量制备；网状结构；显微组织；力学性能中图分类号：TB331文献标志码：A 文章编号：1673-0224(2022)01-56-10Microstructure and properties of high-throughput in situ networkAll Rights Reserved.structure TiC reinforced TC4 composite materialsDU Kanghong, LIU Zhongqiang, ZHANG Jiantao, WEN Liping, XIAO Zhiyu(National Engineering Research of Net-Shape Forming for Metallic Material,South China University of Technology, Guangzhou 510640, China)Abstract: Using TC4 alloy powders with different particle sizes as the matrix material, using VC as the carbon source,using high throughput hot press sintering process, TiC/TC4 composites were prepared with different network structuresizes and TiC volume fractions (2%, 4%, 6%). The effect of TiC content and TC4 powder particle sizes on themicrostructure and properties of composite materials were studied. The results show that the TiC reinforcement particlesin TiC/TC4 titanium matrix composites are distributed in a network. Compared with TC4 alloy, the microstructure ofTiC/TC4 composite materials are significantly refined. As the TiC content increases, the thickness of the TiC networklayer increases, and the tensile strength and elongation of the material first increase and then decrease. The material with2%TiC has the best overall performance. As the particle size of TC4 decreases, the microstructure of the TiC/TC4composite is gradually refined, the connectivity of the matrix increases, and the tensile strength and elongation of thematerial increase at the same time. 2%TiC/TC4 composite material prepared by TC4 alloy powders with a particle size of40−80 μm has a small network structure and the best overall performance. The yield strength, tensile strength andelongation reach 946 MPa, 1058 MPa and 18.1%, respectively, which are 29.6%, 31.6%, and 118.1% higher than TC4alloy.Keywords:titanium matrix composites; in-situ titanium carbide; high-throughput preparation; network structure;microstructure; mechanical properties基金项目：广东省重大科技攻关项目(2019B010942001)；国家自然科学基金资助项目(51627805)收稿日期：2021−08−17；修订日期：2021−10−26通信作者：肖志瑜，教授，博士。
混凝土技术的发展与展望
第39卷第1期硅酸盐通报Vol.39No.1 2020年1月BULLETIN OF THE CHINESE CERAMIC SOCIETY J<u<y,2°2°\特邀综述i**««««««««««*混凝土技术的发展与展望缪昌文V,穆松2(1.东南大学,材料科学与工程学院，南京2111892高性能土木工程材料国家重点实验室，南京211103)摘要:混凝土是基础设施建设的主体建筑材料。
继钢筋混凝土、预应力钢筋混凝土与纤维混凝土之后，基体微结构调控与性能提升将成为未来混凝土技术第四次飞跃的主要驱动力。
本文从近年来混凝土技术发展现状与趋势两个方面入手，重点介绍了混凝土的工作性调控、裂缝控制、力学性能提升、耐久性提升四个方面的主要技术现状，同时对上述混凝土技术的主要发展趋势进行了展望。
关键词:混凝土；微结构；性能；工作性；收缩裂缝；力学性能；耐久性中图分类号:TB321文献标识码:A文章编号：1001-1625(2020)01-0001-11 Development and Prospect of Concrete TechnologyMIAO Changwen'2,MU Song2(1.School of Materials Science and Engineering,Southeast University,Nanjing211189,China；2.State Key Laboratory of High Performance Civit Engineering Materials,Nanjing211103,China)Abstrac):Concrete i t the used buVding material for infrastructure construction.Following the reinforced concrete, prestressed reinforced concrete and finey concrete,the microstructure control and performancc improvement of the concrete will become the main driving force for the fourth leap of concrete technology in the future.The recenr developmene of workabilim regulation,shrinkaae and cracking mitigation,mechanica.properties improvement and durabilim improvement are reviewed.The prospective and main development trends are also introduced.Key wordt:concrete；microstructure；property；workabilim；shrinkaae crack；mechaniccl property；durabilith0引言作为传统硅酸盐材料之一，水泥混凝土诞生至今已有200年历史。
TA12A钛合金热处理过程中等轴和片层α相演变行为研究
Ti 穀臧Vol. 38 No. 1February 2021第38卷第1期2021年 2月TA12A 钛合金热处理过程中等轴和片层！相演变行为研究陈飞1,周瑜2,王柯2(1.陕西宏远航空锻造有限责任公司，陕西 咸阳713801)(2.重庆大学材料科学与工程学院，重庆400044)摘要：对近a 型TA12A 钛合金进行热处理实验，利用扫描电子显微镜(SEM )和电子背散射衍射！ EBSD )技术对热处理后的微观组织进行观察，研究了两相区固溶温度和冷却速率对微观组织的影响。
研究表明：TA12A 钛合金在980和1000 ＞保温后冷却时，"相向a 相转变，一方面可以使得等轴a 相长大，另一方面也可析出片层a 相。
等轴 a 相长大过程中，大的区域与初始 a 相存在成分差异，从 大的区域在背散射电子成像模式下表现出a 环状组织。
空冷时，因冷却速率较快，使得片层a 相快速大，从而抑制了等轴a 晶粒的长大。
但是，固溶温度对 a 晶粒的长大和 a 相的 行为影响 。
同时，冷却速率显著影响 a 相的 ，这与等轴a 相的含量密切相关$关键词：TA12A 钛合金；热处理；等轴a 晶粒；片层a 相中图分类号：TG166. 5 ； TG146. 23文献标识码：A 文章编号：1009-9964(2021 )01B001-05Evolution of Equiaxed and Lamellar a Phase during Heat Treatment of TA12A Titanium AlloyChen Fei 1，Zhou Yu 2，Wang Ke 2(1. Shaanxi Hongyuan Aviation Forging Company Ltd.，Xianyang 713801，China )(2. School of Materials Science and Engineering ，Chongqing University ，Chongqing 400044，China )Abstract : Heat treetmeni expermenis werr conducted on the neer-9 titanium Hoy TA12A. The microstructurr was observed using sccnning electron microscopy ( SEM) and electron backscattered difraction ( EBSD) method, and theeffect of solution Omperature in tmo phase region and cooling rate on the microstructure evolution wat 0x 630164. The resultt show that ，during the cooling proceso after holding at 980 and 1000 ＞，the transformation from " phase te aphase could make equiaxed a grain grow up and Umellar a phase precipimm. During the growth of equiaxed a grain ， the composition of the new-9rowth region is dnferent from that of the initial equiaxed a grain ，which makes the new-growth region show an a eng structure under the SEM observation. The air cooling makes the lamellar a phase nucleate and grow rapidly ，thus inhibiting the growth of equiaxed a graine. However ，the solution temperature has little effecOon thegoowth otequoaeed a goaon and thepoecopotatoon otaameaao a phaee.Fuotheomooe ， thecooaongoatehaeaeognotocanafect on the nucleation site of lamellar a phase ，which is closely related te the content of equiaxed a grain.Key words ： TA12A titanium Cloy ； heat treetment ； equiaxed a grain ; lamellar a phase钛合金具有比强度高、断裂韧性好、高温性能， 航空航一 重要的金属材料。
固溶处理过程中变质AlSi10MnMg压铸铝合金的组织演化及其对力学性能的影响
Trans. Nonferrous Met. Soc. China 29(2019) 919−930Microstructure evolution of modified die-cast AlSi10MnMg alloy during solution treatment and its effect on mechanical propertiesZi-hao YUAN1, Zhi-peng GUO1, Shou-mei XIONG1,21. School of Materials Science and Engineering, Tsinghua University, Beijing 100084, China;2. State Key Laboratory of Automobile Safety and Energy, Tsinghua University, Beijing 100084, ChinaReceived 7 May 2018; accepted 7 January 2019Abstract: To optimize the solution treatment process of a modified high-pressure die-cast AlSi10MnMg alloy, the influence of the solution treatment on the microstructure, mechanical properties and fracture mechanisms was studied using OM, SEM, EBSD and tensile test. The experimental results suggest that the solution treatment could be completed in a shorter time at a temperature much lower than the conventional practice. Surface blistering could be avoided and substantial strengthening effect could be achieved in the following aging process. Prolonging solution treatment time and elevating solution temperature would be meaningless or even harmful. The rapid evolution of eutectic silicon during solution treatment, especially at the early stage, affected the way of interaction among α-Al grains during plastic deformation, and changed the ultimate mechanical properties and fracture mode.Key words: AlSi10MnMg alloy; die casting; solution treatment; microstructure evolution; mechanical properties; process optimization1 IntroductionHigh-pressure die casting (HPDC) is one of the most efficient casting processes for mass production of shaped components. Due to extremely high interfacial heat transfer coefficient [1], the melt solidifies at a high cooling rate, leading to a very refined microstructure and better mechanical properties. However, the complicated flow pattern [2] for thin-walled castings would lead to air entrapment or gas pores, which significantly decreases the ultimate mechanical properties of the component. Besides, the sub-surface gas pores would cause surface blisters during conventional solution treatment (4−10 h at 540 °C for Al−Si−Mg alloys [3−5]), which follows the principle of high temperature and long time to homogenize the alloy elements.Both vacuum-assisted and pore-free die casting could reduce the gas porosity significantly [6,7]. Recent studies [8−11] have also demonstrated that a low temperature solution treatment could reduce surface blisters, and the subsequent artificial aging could improve the tensile strength significantly. Mostly processed by HPDC, the modified AlSi10MnMg alloy has been widely used in automobile industry due to its excellent castability and mechanical performance. Mg could promote the formation of strengthening phases, e.g., Mg2Si at the inter-granular boundary of the as-cast microstructure. In order to further improve the mechanical properties, the Mg2Si phase must dissolve into the α-Al matrix and precipitate as nano-size meta-stable phases inside the grains [12−15].Extensive studies have been performed to investigate the effect of solution treatment on the microstructure and mechanical properties of the cast Al−Si−Mg alloys. The best solution treatment process parameters varied within a wide range, which strongly depend on the alloy composition [16,17], morphology of eutectic silicon [17−19] and manufacture process [4,5]. And the solution treatment affected the final mechanical performance through dissolving and homogenizing alloy elements [16,19,20], as well as changing silicon phase [19,21].Due to the advantages and disadvantages of the die- cast alloys, the optimized solution treatment parameters should be revised, following different principles fromFoundation item: Project (U1537202) supported by the National Natural Science Foundation of China; Project (BA2015041) supported by the Special Funding Program on Transformation of Scientific and Technological Achievements in Jiangsu Province, ChinaCorresponding author: Shou-mei XIONG; Tel: 86-10-62773793; Fax: 86-10-62770190; E-mail: smxiong@DOI:10.1016/S1003-6326(19)65001-6Zi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 920that of conventional heat treatment. Very refined microstructure makes it possible to shorten the solution treatment process at a relatively low temperature, which also reduces the risk of blisters. For the heat treatment of a modified AlSi10MnMg high-pressure die-cast alloys, detailed studies need to be performed to optimize the process parameters. In this work, quantitative analysis was performed on the microstructure evolution and mechanical properties during the solution treatment. In-situ tensile deformation together with fractography analysis was performed to reveal the fracture mode and crack expansion. The process−microstructure−property relationship was investigated, together with the discussion on the optimization of process parameters.2 ExperimentalThe composition of the modified AlSi10MnMg alloy used in this study is listed in Table 1. Si would improve the flowability of the alloy and reduce the risk of hot cracking. The content of Mg is 0.301 wt.% to ensure a distinct strengthening effect, and all the β-Mg2Si phases could dissolve into the α-Al matrix at a relatively low solution temperature. Sr could modify the morphology of the eutectic silicon. The content of Fe was as low as 0.109 wt.% to reduce the brittle needle-like iron-rich β-Al5FeSi during solidification, but at the risk of die soldering. As a compensation, Mn was added into the alloy to reduce die soldering.The HPDC components were produced by using a TOYO BD−350V5 cold chamber die casting machine with key processing parameters listed in Table 2. Oil was circulated inside the dies to stabilize the temperature (at ~120 °C). Figure 1 shows the configuration of the die casting, which comprised 3 bars and 1 plate of 170 mm ×30 mm ×2.5 mm. Detailed configuration of the cylindrical tensile specimen is shown in Fig. 1(b).Table 1 Composition of modified AlSi10MnMg alloy (wt.%)Si Mg Mn Fe Cu Sr Ti Zn Al 10.14 0.301 0.639 0.109 0.016 0.025 0.073 0.007 Bal.Table 2 Key processing parameters of die castingParameter ValueMelt temperature/°C 700Slow shot speed/(m∙s−1) 0.15Gate speed/(m∙s−1) 57.9 Intensification pressure/MPa 79Vacuum degree/kPa ~90The tensile testing bars were put into an electric- resistance muffle furnace at 490 °C for 5, 15, 30, 60, 120, 300, 720 and 1440 min for heat treatment followed by an immediate water quenching. As a comparison, some bars were solutionized at 520 and 540 °C for 30 min, respectively. Artificial aging process was performed at 200 °C for 2 h immediately after quenching to achieve the maximum strength (T6 state).Fig. 1 Configuration of die casting with 3 bars and 1 plate (a), dimensions of tensile testing bar (b) and tensile plate (c) for in-situ observation (unit: mm)Zi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 921 Several sections of samples were polished andobserved using OM (optical microscope). And the OM images were used to calculate the porosity of each sample with ImageProPlus software. To investigate the microstructure evolution, the die-cast and the water- quenched samples were ground, polished, etched with Keller’s etchant and observed usin g a JEOL 7001F SEM. The specimens for EBSD were prepared with the method of ion milling.Mechanical properties of the as-cast state, as-quenched state and/or T6 state alloys were investigated. The yield strength (YS), ultimate tensile strength (UTS) and elongation to fracture were measured at room temperature with a strain rate of 0.0005 s−1. A 50-mm extensometer was used to measure the elongation within the gauge length. To study the fracture mode transfer due to solution treatment, the fracture surfaces of the tensile specimens were investigated using a JEOL 7001F SEM after ultrasonic cleaning in acetone.In order to study the plastic deformation, in-situ tensile samples were prepared from the 2.5-mm thick plate (see Fig. 1(a)) using electrical discharge machining.A 0.3-mm notch was made on the sample to imitate the fatal defect inside the components during service, and facilitate the nucleation of crack. The die-cast and T6-state in-situ samples were ground, polished and etched. In-situ tensile deformation was performed in a SEM chamber with a grip moving speed of 0.1 μm/s.3 Results3.1 Microstructure evolution during solutiontreatmentFigure 2 shows the sections and appearances of tensile bars before and after solution treatment. The porosity was originally about 0.41% in the as-cast tensile bar (see Fig. 2(a)), which slightly went up to 0.52% after 300 min solution treatment (see Fig. 2(b)), and 0.56% after 720-min solution treatment (see Fig. 2(c)). The increase in porosity was generally attributed to the expansion of gas pores [8,9]. Tensile bars were bent because of the gravitational creep after 720-min solution treatment (see Fig. 2(d)). Due to good vacuum degree and relatively low solution temperature, no obvious surface blisters were found until 1440-min solution treatment (see Fig. 2(e)). The solution temperature of 540 °C, which was ordinarily used in conventional practice, caused blisters after only 30 min (see Fig. 2(f)).Figure 3 shows a typical microstructure of a modified high-pressure die-cast AlSi10MnMg alloy and a magnified view of modified eutectic Si. Because of the high cooling rate during solidification, most eutectic phases exhibited a divorced morphology. The externally solidified crystals (ESCs), which originally nucleated inside the shot sleeve, could be observed on the section of casting (see Fig. 3(a)). The α-Al grains, with a size ofFig. 2 Sections (a−c) and appearance (d−f) of tensile barsZi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 922Fig. 3 Typical microstructure of as-cast modified AlSi10MnMg alloy (a) and magnified view of modified eutectic Si (b)15−40 μm, exhibited a globular or dendritic morphology. The fibrous eutectic silicon exhibited a coarse surface (see Fig. 3(b)), which was mostly caused by the modification of Sr. The diameter of the eutectic Si was only 100−300 nm, while its length could be several microns. The Si fibers were deeply embedded into the α-Al matrix, and distributed at the boundaries of large α-Al grains. The β-Mg2Si phase was also found in the eutectic region, as marked by the circle in Fig. 3(b).The influence of the solution treatment time on the microstructure of the alloy and its average diameter as function of solution treatment time at 490 °C was shown in Fig. 4. A small part of β-Mg2Si phases were still not completely dissolved into the matrix after solution treatment at 490 °C for 5 min and fibrous Si changed into particles or short rods (see Fig. 4(a)). The dendrite arms of large α-Al grains coalesced, with Si particles packed in. Some particles forming clusters distributed in α-Al grains. Those particles would detach from the α-Al matrix when deeply etched, generating holes on the polished surface of the samples (see Figs. 4(a−e)). All the β-Mg2Si phases disappeared after solution treatment for 15 min (see Fig. 4(b)). Thus, the solution temperature of 490 °C was sufficient to dissolve all the β-Mg2Si phases in this alloy due to the low content of Mg.As coarsening occurred, the interface between the α-Al matrix and the silicon particles became smooth. During the solution treatment, the average diameter of silicon particles increased rapidly within the first 30 min, i.e., from 0.2 to 1.68 μm, after which it increased rather slowly, i.e., the average diameter increased to only ~2.8 μm after 1 d (1440 min) of solution treatment (see Fig. 4(e)).The grain boundaries could be recognized in the EBSD image quality maps (see Figs. 5(a−e)), according to which the average grain sizes of α-Al matrix could be calculated (see Fig. 5(f)). Over 500 grains at the center of the bar were taken for the statistical analysis in each sample. The statistical average grain area increased rapidly in the first 60 min of solution treatment. That was the result of the disintegration of eutectic silicon. Meanwhile, some dendrite grains still existed (see Fig. 5(b)). After that, the grain size grew slowly. Nearly all grains exhibited a granular morphology, which indicated the coalescence of dendrite arms that eliminated the inner boundaries.3.2 Mechanical propertiesThe samples were deflected and/or blistered after 720-min solution treatment, which would fail in obtaining the uniform uniaxial tensile stress during tensile test. Thus, the corresponding tensile tests were not performed. Figure 6 shows the mean values of YS, UTS and elongation as a function of solution treatment time. The dash line in each figure indicated the mechanical properties of the die-cast state. The first point of each curve corresponded to the solution time of 5 min. The dash line in each figure, i.e., 161 MPa for YS, 314 MPa for UTS, and ~7.5% for elongation, represented the values of the die-cast state.A 5-min solution treatment could reduce both YS and UTS and increase the elongation significantly. However, longer solution treatment produced negligible effect on the as-quenched yield strength which just fluctuated between 98 and 107 MPa (see Fig. 6(a)). The UTS of as-quenched samples increased from 228 MPa at 5 min, to the maximum of 245 MPa at 30 min, then decreased slowly to 232 MPa at 300 min (see Fig. 6(b)). The elongation of as-quenched samples increased from 13% at 5 min, to the maximum of 21% at 30 min, and then decreased slowly to 16% at 300 min (see Fig. 6(c)).A T6-treatment improved both YS and UTS significantly, but reduced elongation. The YS increased significantly from 192 to 257 MPa at treatment time from 5 to 15 min, after which the YS maintained rather stable with a slight decrease to 245 MPa at 300 min. The UTS of the T6 samples increased from 275 MPa at 5 minZi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 923Fig. 4 Microstructures of modified die-cast AlSi10MnMg alloy after solution treatment at 490 °C for different time (a −e) and average diameter as function of solution treatment time at 490 °C (f)to the maximum of 326 MPa at 15 min, and then decreased slowly to 302 MPa at 300 min (see Fig. 6(b)). The elongation increased from 10% at 5 min to the maximum of 13% at 30 min, and then decreased slowly to 9.5% at 300 min (see Fig. 6(c)). As a comparison, the YS, UTS and elongation of 520 °C-solutionized T6 bars were 257 MPa, 329 MPa and 11%, respectively, which have no distinct improvement over the 490 °C- solutionized ones.During the age process, the main microstructure transformation was the precipitation reaction [12,15,22], which turned the α-Al grains from homogeneous solid solution to the phase reinforced with precipitates inside. Those brittle precipitates could hinder the movement of dislocations during plastic deformation. Thus, the aging process improved the strength but lower the elongation compared with that of the as-quenched samples. However, the elongation of T6 state is still higher than that of as-cast state (see dash line in Fig. 6(c)). 3.3 FractographyFigure 7 shows the fracture surface observed using SEM after the tensile test. Without solution treatment, the fracture surface exhibited more brittle features. The modified fiber-like eutectic Si, embedded deeply in the α-Al matrix, and cracked during the tensile test (see Fig. 7(a)). The fracture mode changed as the solution treatment time increased. After 5-min solution treatment, some shallow dimples started to emerge on the fracture surface (see Fig. 7(b)). After solution treatment of 30 min, the fracture surface exhibited a uniform distribution of small dimples of ~2.5 μm (see Fig. 7(c)). Prolonged solution treatment, i.e., 300 min, only enlarged the dimple size to ~5 μm (see Fig. 7(d)). Some spherical Si particles sat at the bottoms of the dimples, as marked by the arrows in Fig. 7(d).3.4 In-situ observation of deformation and crackFigure 8 shows a typical microstructure duringZi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930924Fig. 5 EBSD image quality maps of as-cast specimen (a) and specimens solutionized at 490 °C (b−e), and evolution of average grain area of α-Al matrix versus solution time at 490 °C (f)Fig. 6 Mechanical properties of as-cast, as-quenched and T6 samples after solution treatment at 490 °C: (a) Yield strength (YS); (b) Ultimate tensile strength (UTS); (c) Elongationin-situ tensile deformation. All the view fields were at the crack tip, where severe plastic deformation occurred. In the die-cast sample (see Figs. 8(a) and (b)), the α-Al grains exhibited clear and uniform slip bands, indicatinga uniform plastic deformation. The slip bands stopped in the eutectic region at the α-Al grain boundary. The eutectic silicon was too stiff to deform continuously and harmoniously with the α-Al grains, thus micro-cracksZi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 925Fig. 7 Fracture surfaces of as-cast sample (a) and T6-treated samples at different solution time (b −d)Fig. 8 Deformation at main crack during in-situ tensile test for as-cast (a, b) and T6-treated (c, d) samples ((b) is magnified view of red square in (a), and (d) is magnified view of red square in (c))generated. Both interface decohesion and brittle cracks could be observed (see Fig. 8(b)), i.e., as deformation proceeded, the crack expanded rapidly along the grain boundaries, as shown in Fig. 9(a).In the T6-state sample (see Figs. 8(c) and (d)), thepolished surface of specimen was significantly disturbed at the crack tip. The α-Al grains still remained the capability of slipping after precipitation reaction because the low content of Mg limited the amount of strengthening precipitates. The slip bands wereZi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 926Fig. 9Grain boundaries after deformation for as-cast (a) and T6-treated (b) samplesdisordered, coarse and prone to bifurcate. In Fig. 8(d), the Si particles scattered in the α-Al matrix, indicating the presence of the eutectic before solution treatment. The intricate fiber-like Si network collapsed, and the slip bands could propagate deeply into the eutectic region. Without large aspect ratios and coarse interfaces, most Si particles detached from the α-Al matrix at the interface due to deformation mismatch (see Fig. 8(d)). The plastic deformation smoothly transmitted from grain to grain, or from one dendrite arm to another without inter- granular cracking (see Fig. 9(b)). Meanwhile, small cracks generated at the interface, and became larger when external load was amplified. These minor cracks could act as the potential precursors of the main crack, thus the main crack expanded in zigzags and branched occasionally.4 Discussion4.1 Optimization of solution treatment temperatureConventional practice suggests that ~540 °C is a recommended solution temperature for Al−Si−Mg alloy [4,5,23], in order to shorten the solution time. However, the risk of gas pore expansion and blistering was not taken into consideration because there are no high pressure gas pores inside the traditional castings. As shown in Fig. 2(f), blistering happened on the surface of die-cast tensile test bar only after solution treatment at 540 °C for 30 min.The main concern of solution treatment of die castings is to finish solutionizing before serious pore expansion and blistering. One logical adjustment is lowering the solution temperature, at which the gas pressure in pores remains relatively low and the alloy still has competent strength to delay the pore expansion and blistering. In this study, most solution treatments were carried out at 490 °C, which was ~75 °C lower than the solidus line and ~67 °C lower than the eutectic reaction temperature (see Fig. 10). The experimental result indicated that distinct pore expansion happened very late at 490 °C (Figs. 2(b) and (c)). Besides, 490 °C was estimated to be just slightly higher than the solvus line to ensure the complete dissolution of β-Mg2Si phases, according to Fig. 10 [24,25]. Solution treatment experiment demonstrated that it took a short time to finish the dissolution of β-Mg2Si phases and the spheroidization of the eutectic silicon (see Fig. 4). Substantial improvement on mechanical properties could be achieved after aging process (see Fig. 6). The tensile test also demonstrated that the mechanical properties of 520 °C-solutionized T6-state bars have no distinct improvement over those of the 490 °C-solutionized ones.Fig. 10 Computed vertical sections of phase diagram of Al−Mg−Si alloy with 90 wt.% Al [24,25]The experiment mentioned above demonstrated that the solution treatment of die castings could be finished at a low temperature in a short time. That was the result of the fine microstructure of die castings (see Fig. 3). According tolaw, the evaluated diffusion distances, i.e.,were about 11.3 and 9.4 μm respectively, for Mg and Si, where D was the diffusion coefficient in α-Al matrix at solution temperature of 763 K [26], and t=900 s was the solution time. These distances exceeded the average dendrite arm radius, i.e.,Zi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 927about 8 μm (see Figs. 4(a) and 5(a)), indicating that it was close to the completion of homogenization across the dendrite arms when the solution time was 15 min.4.2 Silicon evolution during solution treatmentAltering the morphology of the eutectic silicon is one of the most important aims of solution treatment. Modified with Sr, the flaky Si turned into fiber-like morphology, which had been extensively studied [27,28]. The observation of deeply etched sample showed that the eutectic Si fibers were connected (see Fig. 3). However, the microstructure was unstable due to the vast and complicated interface between α-Al grains and Si fibers. The modified Si was easier to fragment, as shown in Ref. [29].The microstructure evolution is driven by the minimization of the total free energy of the system. This was achieved by reducing the interface area between Si phase and α-Al matrix [30]. As a result of the Rayleigh − Taylor surface instability, unstable perturbations increased in amplitude before the connected area finally pinched off, and the fibrous silicon fragmented into discrete droplets [31,32]. The second stage was spheroidization, during which the droplets tended to decrease the interfacial curvature. Capillarity-induced diffusion gave rise to the apparent flux of solute from the region with high curvature to that with low curvature [33] as a result of Gibbs −Thomson effect. In that way, the specific interface area was reduced. Further reduction in interfacial energy of the system was accomplished with larger particles growing and smaller particles shrinking, i.e., the Ostwald ripening [34]. The growth of the silicon particles followed the traditional LSW (Lifshitz −Slyozow −Wagner) theory [35]:ktR R =-303(1)where t is the time, R is the averageradius of the particles, 0R is the initial average radius at t =0, and k is a constant. The evolution of the average radius of the particles fitted very well with Eq. (1) after 30-min solution treatment (see Fig. 11, the square of correlation coefficient of this linear regression is 0.993). The evolution of the silicon particles was a combination of fragmentation, spheroidization and Oswald ripening. As the Oswald ripening proceeded, the distribution of the particle radius was adjusted asymptotically following the LSW theory.4.3 Effect of microstructure evolution on deformationand crackingThe morphology of the eutectic Si particles had a great influence on the mechanical properties of the alloy. As a brittle phase, the Si could either strengthen the matrix or promote crack nucleation.Fig. 11 Cube of equivalent eutectic Si radius vs solution timeFor die-cast specimen, the eutectic was divorced (see Fig. 3) as a consequence of extremely high cooling rate [1]. The configuration signified that the silicon − fiber-reinforced Al-matrix composite filled the inter-dendritic arm space and the inter-granular space of the α-Al grains. The hard, brittle and semi-continuous (see Fig. 3(b)) silicon fibers comprised the skeleton to sustain the whole material and provided a great strengthening effect. The complicated interface between α-Al grains and Si fibers ensured the sturdy cohesion. When external force exerted on the specimen, the fiber-like Si could bear a lot of load due to its higher elastic modulus and inflexibility for yield. So, the die-cast specimen exhibited a relatively high strength. On the other hand, the Si particles isolated the plastic α-Al grains from each other (see Fig. 3(a)), and the plastic deformation in one α-Al grain could not be compatible with that in its neighbors (see Figs. 8(a) and (b)), i.e., the potential ductility was significantly suppressed. As the external force increased, the plastic incompatibility between silicon fibers and α-Al grains would ultimately lead to either silicon cracking or interface decohesion (see Fig. 8(b)), leading to an elongation of only ~7.5% (see Fig. 6(c)).The skeleton began to decompose after 5-min solution treatment at 490 °C via necking and fragmentation of the fiber-like Si (Fig. 4(a)). The contribution of solute Mg atoms to yield strength is estimated to be less than 13 MPa, with the equation in Ref. [36]. The strengthening effect of the Si particles with small aspect ratio and smooth surface was much lower than that of the fibrous Si. Accordingly, both YS and UTS would decrease whereas the elongation increased (see Fig. 6). The fracture mode also changed during the 5-min solution treatment (see Fig. 7(b)). However, 5-min solution treatment was not enough to dissolve all the β-Mg 2Si phases (see Fig. 4(a)), so the following artificial age could not make full use of Mg atoms in this alloy to strengthen the α-Al matrix.Zi-hao YUAN, et al/Trans. Nonferrous Met. Soc. China 29(2019) 919−930 928After fragmentation and spheroidization during 30-min solution treatment (see Fig. 4(c) and Fig. 7(c)), the α-Al grains joined together to form a continuous matrix. Discontinuous Si droplets could not hinder all the moving dislocations across the borders, and neighboring α-Al grains cooperated with each other efficiently during plastic deformation (see Figs. 8(c) and (d)). After the dissolution of β-Mg2Si phase, both Mg and Si atoms precipitated as meta-stable phases in the α-Al grains during artificial age [12,15,22]. The corresponding T6-state strength reached the maximum value. Prolonged solution time could not provide higher supersaturation after water quenching. Thus, the strengthening effect of precipitation kept constant. However, the α-Al grains coarsened (see Fig. 5(f)), which weakened the alloy strength according to Hall−Patch equation. Thus, the combined effect led to the decrease in mechanical properties (see Fig. 6).When plastic deformation occurred, the dislocations in α-Al grains moved forward to the grain boundaries, and stopped when encountering the Si particles. Further deformation and dislocation accumulation caused stress concentrations at the interface between α-Al grains and Si particles [21,37]. When stress concentrations exceeded some critical values, crack would initiate either in the silicon particles or at the interface (see Fig. 8(d)). The actual crack mode was a very local event as a result of many factors [37]. According to the in-situ observation (see Fig. 8(d)), the latter cracking mode became dominant for the heat-treated specimen. The Si particles bound loosely with the α-Al matrix. And because the micro-cracks could grow and expand along the interfaces, the Si particles were removed from the α-Al matrix. Further deformation would lead to the join of the micro-cracks as well as the nucleation of new micro-cracks until the final fracture, leaving Si particles at the bottom of the dimple (see Fig. 7(d)).After spheroidization, shorter solution treatment would be more beneficial to the ductility of the material due to smaller particles, shorter inter-particle spacing and larger population of particles. Firstly, despite their lack of the capacity for plastic deformation, those small spherical particles were more flexible to minimize the plastic incompatibility with matrix through movement and rotation (see Fig. 8(d)). Thus, the nominal stress required for particle cracking increased as particle size decreased [37]. Secondly, shorter inter-particle spacing implied shorter mean free path, less dislocation accumulation [38] at interface and lower stress concentration before the initial cracking. Regardless of the cracking mode, the size of the micro-crack reflected the magnitude of deformation incompatibility between the particles and the matrix. The stress concentration factor was proportional to the square root of the crack size, thus smaller inter-particle spacing would lead to smaller micro-cracks and lower stress concentration. Thirdly, severe plastic deformation would occur in front of the tip of cracks. Keeping constant volume fraction of particles during coarsening process, larger population of particles would lead to more micro-cracks. The deformed regions in front of micro-cracks would absorb more energy by work hardening. In this respect, large population of micro-cracks dispersed the plastic deformation, preventing severe strain localization and delaying the expansion of large cracks.5 Conclusions(1) Contrary to the conventional practice, the solution treatment of the die-cast AlSi10MnMg alloy should be carried out under a relatively low temperature in a shorter time. Substantial strengthening effect could be achieved in the following artificial aging process. Prolonging solution time and elevating solution temperature would be meaningless or even harmful due to microstructure coarsening, surface blistering and/or gravitational deflection.(2) The Si particles changed rapidly during solution treatment. After the fragment and spheroidization, the coarsening process of Si particles followed the LSW equation.(3) The morphology and the size of Si particles affected the mechanical properties and the fracture mode through affecting the interaction of α-Al grains. References[1]GUO Zhi-peng, XIONG Shou-mei, CHO S H, CHOI J K. Study onheat transfer behavior at metal/die interface in aluminum alloy die casting process. I: Experimental study and determination of the interfacial heat transfer coefficient [J]. Acta Metallurgica Sinica, 2007, 43: 1149−1154. (in Chinese)[2]JIA Liang-rong, XIONG Shou-mei, LIU Bai-cheng. Study onnumerical simulation of mold filling and heat transfer in die casting process [J]. Journal of Materials Science & Technology, 2000, 16: 269−272.[3]OZHOGA-MASLOVSKAJA O, GARIBOLDI E, LEMKE J N.Conditions for blister formation during thermal cycles of Al−Si−Cu−Fe alloys for high pressure die-casting [J]. Materials & Design, 2016, 92: 151−159.[4]YANG Chang-lin, LI Yuan-bing, DANG Bo, LÜ He-bin, LIU Feng.Effects of cooling rate on solution heat treatment of as-cast A356 alloy [J]. Transactions of Nonferrous Metals Society of China, 2015, 25: 3189−3196.[5]DANG Bo, LIU Cong-cong, LIU Feng, LIU Ying-zhuo, LIYuan-bing. Effect of as-solidified microstructure on subsequent solution-treatment process for A356 Al alloy [J]. Transactions of Nonferrous Metals Society of China, 2016, 26: 634−642.[6]CAO Han-xue, HAO Meng-yao, SHEN Chao, LIANG Peng. Theinfluence of different vacuum degree on the porosity and mechanical properties of aluminum die casting [J]. V acuum, 2017, 146: 278−281.[7]LI Xiao-bo, XIONG Shou-mei, GUO Zhi-peng. Improved。
微电子技术及其发展(Microelectronic technology and its development)
微电子技术及其发展（Microelectronic technology and itsdevelopment）This article is contributed by woaixi27Pdf documents may experience poor browsing on the WAP side. It is recommended that you first select TXT, or download the source file to the local view.Microelectronic technology and its developmentXiu Ping sun(Department of physics, Changchun Institute of light and mechanical engineering, Jilin 130022)Microelectronics is on solid (mainly semiconductor) miniaturized circuit materials composed of Branch Electronics subsystem and system, is a main research and application of motion of electrons or ions in a solid material and use it to realize the signal processing science. Micro electronics is to realize the circuit and system integration the purpose, circuit and system which implements the also known as integrated circuits and integrated system is miniaturization. Application of microelectronics is the microelectronics technology, it is the key of information technology. The spatial scale of microelectronics technology is usually in micron and nanometer units. At present, the level of development of microelectronics technology and the scale of the industry has become an important symbol of the economic strength of a country. A, position and role of microelectronics 1. microelectronics technology is theinformation technology base The basic information is a common form of objective things and movement characteristics of the state, is the third largest resource after material and energy, is one of the three pillars of human material and spiritual civilization development. It will realize the letter differences can be distinguished and natural earthquake, and also the occurrence time given a nuclear explosion. Equivalent position and so on. In addition, the use of seismic methods can also predict the volcano eruption, the research on reservoir induced earthquake can provide security for large reservoir monitoring on mine earthquake is an important means to protect the safety of mines.Key information socialization is a variety of computer and communication, but it is based on microelectronics technology. In the field of information technology, microelectronic technology can realize the information acquisition, transmission and storage, processing and output, information technology has become the cornerstone of.1946 in February was born the first named numerical integrator and calculator electronic computer in the United States the Moore academy, when the machine is composed by the 18000 tubes, weighing 30 tons, covers an area of 150m2, power of 140kW, but the running speed of this huge monster is only 5000 times per second, storage only 1000, stable operation time of only 7 minutes. Imagine that the computer can enter the office, enterprise and family? At that time, so some people think that the world is only 4 such computers is enough, but now all over the world The computer does not include computer, there are millions of Taiwan, caused the great changes of the technology base is 2.. Microelectronics microelectronics technology is an importantstrategic role in the electronic information industry is a sign of the prosperity of the country on the national economy of the first earthquake as a normal natural phenomenon, there are moments when possible, we actively do the prediction before the earthquake, usually need to learn some basic knowledge of earthquake, even if the earthquake would happen, do not panic, take scientific attitude and measures to actively respond to the earthquake disaster as far as possible to reduce to a minimum.Brief introduction of the authorLi Xiang, male, born in 1976 in Shandong Qihe. Master, graduated from the Physics Department of East China Normal University biological physics. Now Shanghai education press science editor. Published several articles. Wang Qing, female, was born in Shanghai in 1976. Graduated from Elementary Education College of Shanghai Normal University. Now the.1996 middle school mathematics teachers in Shanghai city was rated as excellent graduates of Shanghai normal system.2001 years as Putuo District excellent teacher.6Modern physics knowledgeIn the food chain, the development of modern economy data show that the gross national product (G) growth of 100 per dollar, 10 yuan is about 300 NP electronic industry output support, which includes integrated circuit products 1 yuan. And according to the data measured, the integrated circuitcontribution to the national economy is far higher than other categories of products, if the quality of the steel units of G contribution of 1, while the car is 5, TV is NP 30, the computer was 1000, and the contribution of integrated circuit is as high as 2000. so a Japanese economists who control the VLSI technology will control the world industry, the British is that if a country does not grasp the microelectronic technology, it will immediately enter the ranks of developed countries. It is expected around 2007, the microelectronics industry will be more than the value of steel Iron industry. 3.. The micro electronics industry promotes the transformation of other industries. The microelectronics industry has developed rapidly, except for the great contribution of the microelectronic industry itself to the national economy,The permeability and strong. Almost all of the traditional industries as long as combined with microelectronics technology, intelligent transformation of microelectronic chips, will make traditional industry theprocessofrejuvenation. Such as boiler feed water pump in power plant fan, fan power plant energy consumption accounted for 72% of all, and all sectors of the country's wind machine, the total power consumption of the pump power generation accounted for 36%, only for the fan, water pump frequency control and other electronic technology transformation, the annual generating capacity can be saving 65 billion 900 million kwh, equivalent to 3 Gezhouba Dam power plant; communication transmission transformation using the microelectronic technology, the electric locomotive can save 20% ~ 40%. The internal combustion engine fuel can be 12% ~ 14%. Microelectronics technology not only in energy saving,material saving and so on can make the upgrading of traditional industries, but also can make the traditional product structure, performance Changes in areas such as the revolutionary quality. For example, when using the unified control of microcomputer motor device to replace the traditional uranium - worm gear, worm gear, the car will no longer be purely mechanical products, electronic vehicles will cause the car revolution; other digital mechanical processing equipment, digital communication technology is based on Microelectronics technology as the foundation. At present the microelectronic technology has penetrated into the national economy and national defense construction, and all aspects of family life. Two, the general rules of microelectronics technology development of microelectronics technology both from affecting the development speed and the production of human society and life, can be said to be unprecedented in the history of science and technology, the other is any industry can not match. But there are also some regularity of development of microelectronic technology. 1. Moore MOS integrated circuit has become The core of the microelectronics industry, and the memory and microprocessor are the two typical products of the most representative MOS in integrated circuits, their development level usually marks the development level of the microelectronics technology. Since 60s, integrated circuit14 Volume 3 issue (total 81 issues)The development has been followed in 1965 Intel co-founder Moore (G ordon E. Moore) development of the integrated circuit industry predicted that the degree of integrated circuit every 3 years, an increase of 4 times, the feature size was reduced2 times every3 years, this is the famous Moore's law. 2. cases of narrow geometric scaling law of law in 1974 by the Nader board (Dennard) is proposed. The basic idea is to guide his constant in the internal electric field of MOS device under the condition of the longitudinal, scaling device lateral dimension, in order to increase the transconductance and reduce the load capacitance, thereby improving the performance of the integrated circuits. At the same time, the power voltage and device to shrink the size of the same amount the internal electric field device. To maintain constant scaling law called constant field rules, referred to as CE law. If the structure size of devices, After the power supply voltage, threshold voltage according to the law of CE reduced K times, the gate delay time is reduced to K times, mark power integrated circuit performance while reducing delay product K 3 times. At the same time, the channel length and the width of the narrow area of the transistor K 2 times, so the number of transistors on the same chip area into a set the integrated density will increase K 2 times. The scaling law is the basis for the rapid development of VLSI, decades of research and development of integrated circuit technology and device physics are around this point. It is because in the device scaling technology continues to progress and success. The integrated circuit is a brilliant achievements today. However, the law of constant electric field is also a big problem, in order to overcome the existing problems in the law of CE, has been proposed The constant voltage scaling law (CV law), which keep the supply voltage and the threshold voltage is constant, the device scaling law of.CV generally applies only to the channel length is greater than 1 m mu of other parameters, it is not suitable to the device channel length is short. Three, some trends of the 1.21 centurymicroelectronics the development of the technology will continue to silicon CMOS circuit as the mainstream in the development of microelectronic technology's goal is to constantly improve the performance and price performance ratio of the system integration, so it requires to improve the integration of the chip, which is the source of strength shrinking feature size of semiconductor devices. With MOS technology as an example, the channel length reduction can improve the integration the circuit speed; at the same time reduce the channel length and width also reduces the size of the device, improves the integration degree, which can be integrated more transistors in the same size on the chip And can even more complicated structure of electronic system more perfect performance is integrated in a chip. With the increase of the integration level, system speed and reliability is greatly improved, the drop of the price of the.21 century,At least half of twenty-first Century will continue to silicon microelectronics technology as the mainstream. Although great progress has been made in the study of compounds and other new materials of microelectronics, but far does not have the alternative to silicon based process conditions. According to the development of science and technology7The law of development, a new technology from birth to become a mainstream technology generally requires 20 years. In addition, the world trillions of dollars of equipment and technology investment of 30, has led to the formation ofSilicon technology has very powerful industrial capacity. Many famousmicroelectronics experts predict, micro electronics industry will enter the automotive industry in 2030, the aviation industry that mature Chaoyang industry. Even into the microelectronics industry like automobile, aviation and other mature industries, it will still maintain a rapid development trend, as the car aviation industry has developed nearly 100 years still has great potential for development, there is no doubt that the microelectronics industry silicon based at least in the next few years will maintain rapid development trend at present. 2. integrated system is the focus of development of integrated system of microelectronics technology (IS) twenty-first Century Is from the perspective of the whole system, the processing mechanism, model algorithm, chip structure, design circuits of different levels until the device is combined in a single (or a few) complete the function of the system on a chip, integrated system design must be designed from the beginning of the top-down system behavior. Many studies show, and integrated circuit (IC) compared the composition of the system, because the IS design can be integrated and overall consideration of the entire system, can achieve the technical index of higher performance in the same technology conditions. For example, if u use IS method and 0135 m process design system on chip, complexity and processing system for the same rate can the same performance system, achieve the equivalent by 0125 mu m or 0118 m manufacturing process of IC; also There is, compared with the conventional design method of IC chip, using IS design method to complete the number of transistors also needed functions can reduce 1 orders of magnitude. The integrated system rely on the 2 is a very broad spectrum of background, the possibility of a person or several people to conquer the world is very small. The design of system level more includethe function, behavior, algorithm, architecture and even thinking, conception of system background. Change of microelectronic technology from IC to IS is not only a conceptual breakthrough, but also an inevitable result of the development of information technology, it will lead to another based on micro electronics technology information technology revolution.21 century will be the rapid development of IS technology the combination of strong vitality. During the period of the birth of a new growth point of the technology of microelectronics technology 3. microelectronics and other disciplines is that it can reduce The mass production with high reliability and high precision of the micro structure of the module. This technique combined with other disciplines, will produce a series of new discipline and major economic growth point. Micro electro mechanical system (MEMS) is the microelectronics technology to broaden and extend it, microelectronics and precision machining technology mutual integration, realize the system integration of mechanical.MEMS microelectronics and opens up a new field of technology and industry. They can not only reduce the mechanical and electrical system8The cost, but also not many large size electromechanical system to complete the task. It is because of MEMS devices and systems with small size, light weight, low power consumption, low cost, high reliability, excellent performance and powerful advantages of traditional sensor can not match the MEMS in aviation, aerospace, automotive, biomedical, environmental monitoring, military and almost all areas of people are exposedto have a very broad application prospects. The combination of microelectronics and biotechnology close to DNA chip as the representative of the biological engineering chip will be another hot field of microelectronics in twenty-first Century and a new economic growth point. It is based on the biological science as the basis, the use of the characteristics and function of biological organisms, tissues or cells, design with desired traits of new species or varieties, and engineering technology. Combined with the production, it is a product of the combination of science and life science and technology, with high added value, less resource consumption and a series of features, has received widespread attention. The most representative biological chip is DNA chip, the chip can be detected or discovered genetic changes in short this time, no doubt on the genetics, disease diagnosis, treatment and prevention, genetic engineering has extremely important function of.DNA chip can be used not only for genetic research, biomedical etc.,Moreover, with the development of DNA chip, microelectronic biological information system will be formed, and this technology will be widely used in agriculture, industry and environmental protection in all aspects of human lifeBrief introduction of the authorSun Xiuping, born in Jilin in 1965 1988. Graduated from the Physics Department of Jilin University in 2001 with a master's degree in microelectronics and solid state electronics. Now Changchun Institute of Optics and fine mechanics lecturer, engaged in the research work of university physics teaching andoptoelectronic technology. Modern physics knowledge One。
The evolution of microstructure and mechanical pro
International Journal of Minerals, Metallurgy and Materials Volume 25, Number 9, September 2018, Page 1080https:///10.1007/s12613-018-1659-7Corresponding author: Hai-tao Zhang E-mail:haitao_zhang@© University of Science and Technology Beijing and Springer-Verlag GmbH Germany, part of Springer Nature 2018The evolution of microstructure and mechanical properties duringhigh-speed direct-chill casting in different Al–Mg2Si in situ compositesDong-tao Wang, Hai-tao Zhang, Lei Li, Hai-lin Wu, Ke Qin, and Jian-zhong CuiKey Laboratory of Electromagnetic Processing of Materials, Ministry of Education, Northeastern University, Shenyang 110819, China(Received: 1 November 2017; revised: 9 May 2018; accepted: 10 May 2018)Abstract: The effect of high-speed direct-chill (DC) casting on the microstructure and mechanical properties of Al–Mg2Si in situ composites and AA6061 alloy was investigated. The microstructural evolution of the Al–Mg2Si composites and AA6061 alloy was examined by optical microscopy, field-emission scanning electron microscopy (FE-SEM) and transmission electron microscopy (TEM). The results revealed that an increase of the casting speed substantially refined the primary Mg2Si particles (from 28 to 12 μm), the spacing of eutectic Mg2Si (from 3 to 0.5 μm), and the grains of AA6061 alloy (from 102 to 22 μm). The morphology of the eutectic Mg2Si transformed from lamellar to rod-like and fibrous with increasing casting speed. The tensile tests showed that the yield strength, tensile strength, and elongation improved at higher casting speeds because of refinement of the Mg2Si phase and the grains in the Al–Mg2Si composites and the AA6061 alloy. High-speed DC casting is demonstrated to be an effective method to improve the mechanical properties of Al–Mg2Si composites and AA6061 alloy billets.Keywords: Al–Mg2Si in situ composite; casting speed; grain size; primary Mg2Si; mechanical property1. IntroductionAl–Mg2Si in situ composites have potential applications in the aviation and automotive industries [1–6]. These com-posites offer numerous attractive advantages, including a high elasticity modulus, low density, good wear resistance, and a low thermal expansion coefficient, because of the bene-ficial properties of the Mg2Si intermetallic compound [7–13]. In as-cast Al–Mg2Si composites, the primary Mg2Si particles are coarse and exhibit an irregular morphology [14–19], which is detrimental to the composites’ mechanical proper-ties, thereby limiting their range of applications. Therefore, improving the microstructure and mechanical properties remains a critical issue in the further development of Al–Mg2Si composites. Moreover, Al–Mg–Si alloys take a large proportion of the total aluminum alloys, which are used as structural materials. As a conventional Al–Mg–Si alloy, AA6061 alloy exhibits favorable forming properties, high strength, high corrosion resistance, and good welding performance [20–22]. Therefore, Al–Mg–Si alloy has been widely used in industrial and construction applications. However, the coarse intermetallic phase and low solid solu-bility limit the mechanical properties of AA6061 billet be-cause of the low cooling rate in conventional DC casting.As an important method for producing aluminum alloy billets, DC casting can be used to produce aluminum alloy billets in large quantities. DC casting is important for indus-trial applications of aluminum alloys. However, the conven-tional DC casting process has many disadvantages, includ-ing a low cooling rate, coarse precipitated phase, low solid solubility, and slow melt flow; these shortcomings result in a coarse microstructure and poor mechanical properties of the cast aluminum alloy billets.In the DC casting process, the casting speed can affect the cooling rate, geometry of the liquid sump, and the melt flow. Increasing the casting speed can refine the grain and intermetallic phase, increase the solid solubility, decrease the thickness of the segregation layer on the billet surface, and improve the surface quality of the billet [23–24]. These are very advantageous for the industrial production and ap-D.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting (1081)plication of aluminum alloys. Generally, researchers have focused on the effect of casting speed (100–200 mm·min–1) on the microstructures of aluminum alloys [25–28]. Howev-er, the low casting speed is not sufficient to substantially in-crease the cooling rate, resulting in a limited refinement ef-fect on the grain and intermetallic phase in aluminum alloys.In the present work, we developed a high-speed (300 mm·min–1) DC casting process to improve the microstruc-ture and mechanical properties of Al–Mg2Si in situ compo-sites and AA6061 alloy. Moreover, in order to satisfy the strong cooling demand in a high-speed DC casting process, we improved the design of the cooling water system. The aim of this paper is to investigate the effects of high-speed DC casting on the microstructure and mechanical properties of Al–Mg2Si composites and AA6061 alloy.2. ExperimentalFig. 1 shows a schematic of the high-speed DC casting experiment. First, the starting block was positioned in the copper mold. The melt was poured into the copper mold, which formed a solid shell immediately in the primary cooling region. The starting block was then steadily with-drawn from the copper mold. Finally, the billet surface was directly cooled by water jets (secondary cooling) to achieve the DC casting process. In the conventional DC casting process, cooling water is jetted onto the billet surface by only single-row nozzles, which is insufficient to provide strong cooling during the high-speed DC casting process and results in the melt breaking out. We designed an im-proved secondary cooling system with the nozzles of three rows (Positions 1-3 in Fig. 1). This multiple cooling water system improved both the availability of cooling water and heat transfer and provided sufficient cooling during the high-speed DC casting process.Fig. 1. Schematic of the high-speed direct chill casting process.The chemical compositions of different alloys are listed in Table 1. Commercial pure Al, pure Mg, and Al–20wt%Si master alloy were used to prepare the Al–16%Mg2Si and Al–8%Mg2Si composites and the AA6061 alloy in a 100-kW resistance furnace. When the pure Al and Al-20wt%Si were melted, pure Mg was added to the melt at 650–670°C. The melt was subsequently heated to 800°C and maintained at this temperature for 20 min. The melt was degassed with hexachloroethane dry tablets (0.5wt% of the molten alloy) at 780°C for 6 min, following slag removal. After being held for 15 min, the melt was stirred manually for 3 min to en-sure thorough mixing. Finally, the melt was poured into a copper mold to start the DC casting process. The casting speed was varied in the range 50–300 mm·min–1. Billets with a diameter of φ106 mm were successfully prepared via the high-speed DC casting process.Table 1. Chemical compositions of the alloys wt% Alloy Mg SiFeCuCr Zn MnAl–16%Mg2Si10.11 5.890.07 0.03 0.020.030.06 Al–8%Mg2Si 5.05 2.950.07 0.03 0.020.030.06AA6061 1.020.610.520.310.250.250.15Specimens were cut from the same positions of the soli-dified billets for microstructural observation. The micro-structures were observed on a Leica optical microscope after being etched with 1vol% HF solution. The morphologies of the primary and eutectic Mg2Si crystals were examined on a Zeiss ULTRA PLUS field-emission scanning electron mi-croscope after the samples were deeply etched with 20vol% NaOH solution for 20 s at room temperature. The ImagePro Plus software was used to analyze the size and the area frac-tion of primary Mg2Si particles, the area fraction and the spacing of eutectic Mg2Si, and the grain size and area frac-tion of the intermetallic phase. The eutectic Mg2Si phases were observed on a FEI TECNAI G220 transmission elec-tron microscope operated at 200 kV.The tensile test samples were machined according to standard ASTM B557M. The tensile test was conducted on a SHIMADZU AG-X100 kN tension machine at a cross-head speed of 1 mm·min–1. Each tensile result was the average value of eight tension specimens. The fracture surfaces of the tensile test specimens were observed by field-emission scanning electron microscopy (FE-SEM).3. Results and discussion3.1. MicrostructuresFigs. 2(a)–2(d) shows the microstructural evolution of1082 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018the Al–16%Mg 2Si composite at different casting speeds. The increased casting speed effectively refines the primary Mg 2Si; the morphology of the primary Mg 2Si transforms from an irregular to a regular polygon shape. As evident in Figs. 2(e)–2(h), the increase of casting speed substantial-ly refined the eutectic Mg 2Si in the Al–8%Mg 2Si compo-site. In the Al–Mg 2Si composites, the eutectic Mg 2Si ex-hibits a lamellar morphology and a large eutectic spacing when cast at low speed (Figs. 2(i) and 2(k)). The increase of the casting speed refines the eutectic Mg 2Si and sub-stantially decreases the eutectic spacing (Figs. 2(j) and2(l)).Fig. 2. Microstructures of Al–Mg–Si alloys at different DC casting speeds: Al–16%Mg 2Si cast at (a) 50 mm·min –1, (b) 100 mm·min –1, (c) 200 mm·min –1, and (d) 300 mm·min –1; eutectic morphology of Al–16%Mg 2Si cast at (i) 50 mm·min –1 and (j) 300 mm·min –1; Al–8%Mg 2Si cast at (e) 50 mm·min –1, (f) 100 mm·min –1, (g) 200 mm·min –1, and (h) 300 mm·min –1; eutectic morphology of Al–8%Mg 2Si cast at (k) 50 mm·min –1 and (l) 300 mm·min –1.D.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting …1083Figs. 3(a)–3(h) shows the microstructural evolution of the AA6061 alloy cast at different speeds. The Mg 2Si phase was refined and transformed from bulk and strip-like shapes to dot-like shapes with increasing casting speed; the reticular α-AlFeSi phase was also substantially refined at high casting speeds. Figs. 3(b) and 3(g) shows the evolution of the grain structures of the AA6061 alloy specimens. The grain size ofthe AA6061 alloy decreases with the increase of casting speed.Fig. 3. Microstructures of AA6061 alloy at different casting speeds: (a) optical image, (b) grains, and (c) SEM image of AA6061 al-loy cast at 50 mm·min –1; optical images of the AA6061 alloy cast at (d) 100 mm·min –1 and (e) 200 mm·min –1; (f) optical image, (g) grains, and (h) SEM image of AA6061 alloy at 300 mm·min –1.Figs. 4(a)–4(c) shows the area fraction and the size of the primary Mg 2Si at the different casting speeds. With increas-ing casting speed, the area fraction of the primary Mg 2Si decreases from 6.9% (50 mm·min –1) to 4% (300 mm·min –1)1084 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018and the size of the primary Mg 2Si decreases from 28 (50 mm·min –1) to 11 μm (300 mm·min –1). Moreover, the spacing of the eutectic Mg 2Si decreases from 2.7 (50 mm·min –1) to 0.4 μm (300 mm·min –1) in Fig. 4(d). These results indicate that the high-speed DC casting adequately refines the mi-crostructure of the Al–16%Mg 2Si composite. Figs. 4(a) and 4(d) shows the area fraction of the eutectic Mg 2Si and the spacing of eutectic Mg 2Si in the Al–8%Mg 2Si samples cast at different speeds. Notably, the eutectic spacing decreasesfrom 3.5 (100 mm·min –1) to 0.4 μm (300 mm·min –1). The area fraction of the eutectic Mg 2Si increases from 7.3% to 13.2%. As shown in Figs. 4(a) and 4(b), the grain size of the AA6061 alloy decreased from 105 (50 mm·min –1) to 26 μm (300 mm·min –1) and the area fraction of the intermetallic phases decreased from 5.2% (50 mm·min –1) to 2.8% (300 mm·min –1) with increasing casting speed. These results in-dicate that the high casting speed improves the solid solubil-ity of the different alloy elements in the α-Al solid solution.Fig. 4. Area fraction, grain size and eutectic spacing of the different Al–Mg–Si alloys: (a) area fraction; (b) grain size of the AA6061 alloy; (c) the primary Mg 2Si size; (d) the eutectic spacing as functions of the casting speed.Fig. 5 shows TEM images of the eutectic Mg 2Si of Al–Mg 2Si composites cast at high DC casting speed. Com-bined with the morphology observations of the eutectic Mg 2Si in Figs. 2(j) and 5(a), these results show that the eu-tectic Mg 2Si in the Al–16%Mg 2Si composite exhibits a rod-like and fibrous shape, with a diameter of 200 nm. By contrast, the eutectic Mg 2Si in the Al–8%Mg 2Si composite exhibits a long and lath-shaped morphology, with a width from 100 to 200 nm and length from 2 to 4 μm, as shown in Fig. 5(b). The eutectic Mg 2Si in the different Al–Mg 2Si composites exhibits different morphologies and sizes at the highest casting speed.The growth rate of Mg 2Si crystals differs substantially in different directions because of the faceted growth style. The growth of Mg 2Si crystals is slower than that of α-Al. Theα-Al can surround the eutectic Mg 2Si in the growth process. Therefore, the continual growth of Mg 2Si crystals depends on re-nucleation. The increased cooling rate promotes nucleation and growth of the eutectic and thus results in the observed fine eutectic structure with rod-like and fibrous shapes. 3.2. Mechanical propertiesFig. 6(a) shows the yield strength of the different alloys cast at different speeds. The yield strengths of the different alloys all improve with increasing casting speed. Notably, the Al–8%Mg 2Si composite exhibits higher yield strength than the Al–16%Mg 2Si composite and the AA6061 alloy. High casting speeds effectively increase the cooling rateD.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting (1085)during the DC casting process, resulting in a finer Mg 2Si phase in the Al–Mg 2Si composites and finer grains in the AA6061 alloy. The refinement of grains and the Mg 2Si phase causes an increase in the number of grain boundaries, which can act as obstacles to dislocation motion [29–30]. The increase in yield strength due to grain boundaries, ΔσGB ,is described by the Hall–Petch equation [30]:12GB y k d σ-∆= (1) where d is the average grain diameter and k y is the Hall–Petch slope. Decreasing grain size and refining the Mg 2Si phase improve the yield strength of the different al-loys via the augmentation of grain boundaries.Fig. 5. TEM images of the eutectic Mg 2Si of the Al–Mg 2Si in situ composites cast at high DC casting speed (300 mm·min –1): (a) Al–16%Mg 2Si; (b) Al–8%Mg 2Si.Figs. 6(b) and 6(c) shows the tensile strength and the elongation at the different casting speeds. Both the tensile strength and the elongation improve with increasing casting speed in the different alloys. The fine and regular Mg 2Si par-ticles strengthen the load-bearing capacity and suppress crack propagation along the particles, which results in enhance-ment of the strength and the ductility of the Al–16%Mg 2Si composite. Moreover, the refinement of grains and the re-duction of intermetallic phases results in increases of thetensile strength and the elongation in the AA6061 alloy.The Al–16%Mg 2Si composite cast at the highest casting speed (300 mm·min –1) exhibits lower elongation (4.33%) than the AA6061 alloy (18.85%) because of the brittle Mg 2Si particles. The Al–8%Mg 2Si composite exhibits theFig. 6. Tensile properties of different Al–Mg–Si alloys: (a) yield strength; (b) tensile strength; (c) elongation.1086 Int. J. Miner. Metall. Mater ., Vol. 25, No. 9, Sep. 2018highest yield and tensile strength; it shows the lowest elon-gation (0.82%) because of its high area fraction and long lath shape of eutectic Mg 2Si (Fig. 5(b)). Fig. 7 shows the engineering stress–strain curves of the different alloys. The Al–8%Mg 2Si composite does not exhibit the evident plastic deformation stage during the tensile process, consistent with its low elongation. The engineering stress–strain curves of the Al–16%Mg 2Si composite show some extent of plastic deformation; thus, its elongation increases in samples cast at high speeds. The AA6061 alloy undergoes the obvious plas-tic deformation stage and greater elongation with increasing casting speed. During the high-speed DC casting process,the high cooling rate is the decisive factor for improving themicrostructure and mechanical performance.Dislocation theory can explain the improvement of ten-sile strength [29–30]. The process of plastic deformation can yield a large number of moving dislocations. In the Al–Mg 2Si composites, the grain boundary of the Mg 2Si phase hinders the dislocation glide and leads to dislocation accumulation. Substantial dislocation accumulation will produce the driving force for dislocation glide. With in-creasing Mg 2Si crystal size, an increase in dislocation ac-cumulation implies a higher driving force of dislocation glide. The accumulated dislocations are easier to glide, re-sulting in a diminished strengthening effect. Therefore, finer Mg 2Si phases result in less dislocation accumulation and a weaker driving force of dislocation glide; it blocks disloca-tion glide during the deformation process and thereby im-proves the tensile strength. Moreover, solution strengthening is also a factor responsible for tensile-strength improvement in the AA6061 alloy because of the increased solid solubili-ty of the different alloying elements. 3.3. FractographyFigs. 8(a)–8(f) shows the fracture surfaces of the different alloys cast at different speeds. Fig. 8(a) shows a typical fracture surface of the Al–16%Mg 2Si composite cast at low speed. The fracture surface includes the complete Mg 2Si particles, which implies that the fracture occurs at the inter-face between the particles and the matrix. In conventional DC casting, coarse and irregular Mg 2Si particles result in an increase in both the stress concentration and crack initiation. The cracks often initiate at susceptible and weak points (i.e., coarse and irregular Mg 2Si particles) along the interface between the matrix and the particle. The particles are diffi-cult to bear higher local active stress in comparison with their intrinsic yield strength. Therefore, the coarse Mg 2Si particles break the continuity of the Al matrix and decreaseFig. 7. Engineering stress–strain curves for specimens cast at different speeds (as cast): (a)Al–16%Mg 2Si; (b) Al–8%Mg 2Si; (c) AA6061.D.T. Wang et al., The evolution of microstructure and mechanical properties during high-speed direct-chill casting (1087)the ductility of the composite. Moreover, the lamellar and coarse eutectic structure encourages the propagation of cracks and thus further restricts the tensile properties of the Al–16%Mg 2Si composite.Fig. 8. Fracture surfaces of specimens cast at different DC casting speeds: Al–16%Mg 2Si cast at (a) 100 mm·min –1 and (b) 300 mm·min –1; Al–8%Mg 2Si cast at (c) 100 mm·min –1 and (d) 300 mm·min –1; and AA6061 cast at (e) 100 mm·min –1 and (f) 300 mm·min –1.The refinement of Mg 2Si particles enhances both the continuity of the Al matrix and the load-bearing ability of the Mg 2Si particles. When the interface-bearing stress is higher than the intrinsic yield strength of Mg 2Si particles, internal cracks occur at the Mg 2Si particles, as shown in Fig. 8(b). Moreover, the fine dimples were observed in the Al matrix, which enhances the ductility of the composite. In this case, the fracture mechanism is both brittle fracture and small extent of ductile fracture. Fine, regular, and homoge-neous Mg 2Si phases can restrain crack initiation to extend along the Mg 2Si particles and strengthen the cohesion of the α-Al matrix.Fig. 8(c) shows the fracture surface of an Al–8%Mg 2Si composite cast at low speed; it exhibits clear cleavage cha-racteristics derived from its intrinsic brittleness. This re-markable brittle fracture results in this sample exhibiting the lowest elongation among the investigated specimens. With increasing casting speed, the fracture surface did not exhibit an obvious cleavage plane in Fig. 8(d); the elongation slightly increased. The increase of the casting speed had no evident effect on the transformation of the fracture characteristic. Fig. 8(e) shows a typical fracture surface of the AA6061 alloy. The excellent ductility of the alloy coincides with the formation of substantial dimples on the formed separation surface. With increasing casting speed, the elongation fur-ther increases because of an increase in the number of dim-ples and the refinement of the dimples, as shown in Fig. 8(f).1088 Int. J. Miner. Metall. Mater., Vol. 25, No. 9, Sep. 20184. ConclusionsThe effect of high-speed DC casting on microstructures and mechanical properties of different Al–Mg–Si alloys was studied. The following conclusions were drawn from the results of the present investigation:(1) A high DC casting speed substantially refines the primary Mg2Si particles and the eutectic Mg2Si structure of the Al–Mg2Si in situ composites. It also effectively decreas-es the grain size and refines the intermetallic phases of AA6061 alloy. With increasing casting speed, the primary Mg2Si transforms from an irregular to a polygonal mor-phology; the eutectic Mg2Si changes from lamellar to rod-like and fibrous morphologies. The morphology of eu-tectic Mg2Si shows the difference in the different Al–Mg2Si composites when the billets were cast at the highest casting speed.(2) The microstructural improvements, including the re-finement of theMg2Si phase, the morphology transformation of the Mg2Si phase, and the decrease of grain size, effec-tively strengthen the mechanical properties of the Al–Mg2Si composites and the AA6061 alloy. At the highest casting speed, the AA6061 alloy exhibits the ductile fracture cha-racteristic and the Al–Mg2Si composites shows more brittle fracture characteristic, which is in accordance with the dif-ference of the elongation.AcknowledgementsThis work was financially supported by the Science and Technology Program of Guangzhou, China (No. 2015B090926013), Postdoctoral Science Foundation of China (No. 2015M581348), Postdoctoral Science Founda-tion of Northeastern University (No. 20150302), and the Doctoral Foundation of Chinese Ministry of Education (No. 20130042130001).References[1] J. Zhang, Z. Fan, Y. Wang, and B. Zhou, Microstructural re-finement in Al-Mg2Si in situ composites, J. Mater. Sci. Lett.,18(1999), No. 10, p. 783.[2] C. Li, Y.Y. Wu, H. Li, and X.F. Liu, Morphological evolu-tion and growth mechanism of primary Mg2Si phase inAl-Mg2Si alloys, Acta Mater., 59(2011), No. 3, p. 1058.[3] B.H. Yu, D. Chen, Q.B. Tang, C.L. Wang, and D.H. 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微机电系统工程专业英语词汇
微机电系统工程专业英语词汇微机电工程材料该课程介绍了微机电系统中常用的材料的性质、选择和应用，包括金属、半导体、陶瓷、聚合物等。
该课程涉及的专业英语词汇如下：中文英文微机电系统Microelectromechanical Systems (MEMS)材料Material性质Property选择Selection应用Application金属Metal半导体Semiconductor陶瓷Ceramic聚合物Polymer晶体结构Crystal structure晶格常数Lattice constant晶向Crystal orientation晶面Crystal plane点阵Lattice基元Primitive cell单位晶胞Unit cell晶界Grain boundary缺陷Defect掺杂Doping禁带宽度Band gap width载流子Carrier导电性Conductivity阻值Resistance应力Stress应变Strain弹性模量Elastic modulus泊松比Poisson's ratio热膨胀系数Thermal expansion coefficient热导率Thermal conductivity热容量Heat capacity熔点Melting point硬度Hardness韧性Toughness耐腐蚀性Corrosion resistance中文英文表面粗糙度Surface roughness表面处理Surface treatment氧化层Oxide layer沉积方法Deposition method微机电器件与系统该课程介绍了微机电系统中常见的器件和系统的原理、结构和功能，包括传感器、执行器、开关、滤波器、振荡器等。
该课程涉及的专业英语词汇如下：中文英文微机电器件与系统Microelectromechanical Devices and Systems (MEDS)原理Principle结构Structure功能Function传感器Sensor执行器Actuator开关Switch滤波器Filter振荡器Oscillator振动Vibration静电力Electrostatic force电容Capacitance电感Inductance电阻Resistance压电效应Piezoelectric effect热电效应Thermoelectric effect磁电效应Magneto-electric effect光电效应Photoelectric effect声波Acoustic wave表面声波Surface acoustic wave (SAW)布拉格反射器Bragg reflector共振腔Resonant cavity质量灵敏度Mass sensitivity频率响应Frequency response功率消耗Power consumption信噪比Signal-to-noise ratio (SNR)线性度Linearity稳定性Stability可靠性Reliability微机械学该课程介绍了微机械系统的基本理论和方法，包括微尺度下的力学、流体力学、热力学、电磁学等。
Microstructure and mechanical properties of twinned
Microstructure and mechanical properties of twinned Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy processed by mechanical alloying and spark plasmasinteringSicong Fang,Weiping Chen,Zhiqiang Fu ⇑School of Mechanical and Automotive Engineering,South China University of Technology,Guangzhou,Guangdong 510640,Chinaa r t i c l e i n f o Article history:Received 9July 2013Accepted 29August 2013Available online 7September 2013Keywords:High entropy alloys Mechanical alloying Spark plasma sintering Nanoscale twins Microstructurea b s t r a c tMost of multi-component high entropy alloys (HEAs)only consist of metallic elements.In the present paper,by introducing nonmetallic carbon element,non-equiatomic Al 0.5CrFeNiCo 0.3C 0.2HEA has been successfully prepared by mechanical alloying (MA)and spark plasma sintering (SPS)process.Alloying behavior,microstructure,phase evolution and mechanical properties of the alloy were investigated systematically.During the MA process,a supersaturated solid solution with both face-center cubic (FCC)and body-center cubic (BCC)structures was formed within 38h of milling.However,a major FCC phase,a BCC phase,Cr 23C 6carbide and an ordered BCC phase were observed after SPS.The FCC phase is enriched in Fe–Ni,the BCC phase is enriched in Ni–Al and the ordered BCC phase is especially enriched in Al,respectively.In addition,nanoscale deformation twins obviously presented only in partial FCC phase after SPS.The compressive strength and Vickers hardness of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy are 2131MPa and 617±25HV,respectively.Ó2013Elsevier Ltd.All rights reserved.1.IntroductionTraditional alloys are typically composed of one principal ele-ment that occupies at least 50at.%in the composition,with minor additions of other elements to obtain definite microstructure and properties,such as Fe-,Al-,Cu-and Mg-based alloys [1].However,high entropy alloys (HEAs)that proposed by Yeh et al.in 2004have broken this conventional paradigm [2,3].This kind of alloys are de-fined as alloys which consist of at least five principal elements,and the concentration of each constituent element ranges from 5to 35at.%.Previous researches show that HEAs can be processed to form simple solid solution structures instead of intermetallics and other complicated compounds [4–6].This phenomenon is commonly attributed to the high configurational entropy in the solid solution state of HEAs [2,7].Furthermore,HEAs have also exhibited interesting properties such as high hardness and high strength [8,9],good thermal stability [10],outstanding wear and oxidation resistance [3,11],which offer great potential for engi-neering applications.The HEA systems explored in the past decade show that metal-lic elements are the most commonly used,e.g.Al,Cr,Fe,Co,Ni,Cu,Ti,etc.[12–14].It is known that the proper addition of nonmetallic elements like C in some traditional alloys is favorable to theirstructures and properties [15,16].However,to the best of our knowledge,HEAs with addition of C element have been rarely investigated and reported.In contrast,AlCrFeNiCo HEA system pre-pared by arc-melting has been extensively studied in existing liter-atures [5,17,18].Wang et al.[19]have investigated microstructure and mechanical properties of Al x CrFeNiCo (06x 62)HEAs,finding that the as-cast Al x CrFeNiCo alloys can possess FCC and/or BCC structure(s)depending on the aluminum content.Increasing con-centration of Al in this alloy system can lead to the formation of BCC phase,which possess high strength and high hardness while inferior plasticity.Apparently,the concentration of Al element should be in an appropriate range to achieve optimum properties.Hence,considering all the factors discussed above,Al 0.5CrFeNiC-o 0.3C 0.2HEA was designed.Moreover,the most widely studied processing route for HEAs is arc-melting (casting),and only a few reports deal with mechanical alloying (MA)[20–23].As a widely used solid state processing route,MA can easily fabricate nanocrystalline materials with good homogeneity from elemental powders,thus MA can be an ideal way to prepare HEAs [20,24,25].In addition,SPS can rapidly con-solidate alloy powders to high density by applying pressure and passing electric pulse current within short soaking time [26].HEAs synthesized by MA and SPS have been reported to possess good densification characteristics,as well as high strength and high hardness [27–29].Hence,MA and SPS were combined to prepare Al 0.5CrFeNiCo 0.3C 0.2HEA.Alloying behavior and phase evolution0261-3069/$-see front matter Ó2013Elsevier Ltd.All rights reserved./10.1016/j.matdes.2013.08.099Corresponding author.Tel.:+862087113832;fax:+862087112111.E-mail addresses:kopyhit@ ,fu.zhiqiang@ (Z.Fu).during milling and consolidation processesied.Microstructure and mechanical properties ofwere investigated.2.Experiment proceduresHEA with a nominal composition of Al0.5pressed in molar ratio)was prepared by drythen by wet milling for4h in ethanol.Al,Cr,Fe,elemental powders with purity higher than99.9size of645l m(325mesh)were mechanicallywas carried out in a high energy planetaryPlanetary Ball Mill)at300rpm with a ball toof10:1under argon atmosphere.Highvials and tungsten carbide balls(10mm inthe milling media.In order to confirm the alloymilling,powder samples were taken out after27,38,42h respectively.The42h ball milledsubsequently sintered by Dr.Sinter Model SPS-825Spark PlasmaSintering System(Sumitomo Coal Mining Co.Ltd.,Japan)at 1273K for8min at a pressure of30MPa under vacuum(residual cell pressure<8Pa).The samples were heated to873K within 4min,while from873K to1173K and from1173K to1273K, heating rates of75K minÀ1and50K minÀ1were used, respectively.The milled powders and the bulk alloy after sintering were ana-lyzed by a Bruker D8ADVANCE X-ray diffractometer(XRD)with a Cu K a radiation.The microstructure of the alloy was revealed by etching in aqua regia and observed using scanning electron micros-copy(SEM,Zeiss Supra40,Carl Zeiss NTS GmbH,Germany).Thin-foil specimens were prepared by mechanical thinning followed by ion milling at room temperature and were analyzed by a transmis-sion electron microscopy(200kV TEM,JEM-2100,JEOL,Japan) with selected area electron diffraction(SAED)analysis.According to GB/T7314-2005[30],the room-temperature compressive prop-erties of the cylindrical samples(U3mmÂ4.5mm in size)were measured with an Instron5500testing system at a strain rate of 1Â10À3sÀ1.Three compression tests were performed to obtain average value.Hardness measurement was conducted using a Dig-ital Micro Hardness Tester HVS-1000Vickers hardness instrument under a load of300gf.The reported hardness value is an average of at least10measurements.3.Results and discussion3.1.XRD analysis and microstructure during MAFig.1shows the XRD patterns of Al0.5CrFeNiCo0.3C0.2HEA pow-ders prepared under different milling durations.It can be seen that diffraction patterns of all alloy elements are included in the initial blending powder.As the milling time increases,drastic decrement of diffraction intensity is observed after6h of milling.The peaks of Al,Co and C elements dissolve most rapidly.According to Chen et al.[31],the alloying sequence for the Cu0.5NiAlCoCrFeTiMo alloy system correlates best with the melting point of the component elements.Thus rapid dissolution of both Al and Co could be asso-ciated with their lower melting point than other elements.In con-trast,the early disappearance of diffraction peak of C may result from its smallest atomic fraction in the alloy.Many diffraction peaks can hardly be seen when the milling time reaches up to 27h.The complete disappearance of all the elemental peaks and the formation of solid solution are founded within38h of milling. Predominant peaks corresponding to an face-centered cubic(FCC) and a body-centered cubic(BCC)crystal structures are respectively observed.Subsequently,the powder was subjected to wet ball milling in ethanol for4h with the aim to obtainfine metallic pow-der for being conductive to sintering.As the milling time extends to42h,the diffraction peaks exhibit no change except for a minor broadening and a significant increment of diffraction intensity.The crystal size and lattice strain of the BCC and FCC phases with different milling time have been calculated from the X-ray peak broadening using Scherrer’s formula after deducting the instrumental contribution.The calculated results are listed in Table1.The crystalline sizes of both BCC and FCC phases after 42h of milling are slightly refined compared with38h of milling, while the lattice strain of these two phases increase as the milling time prolongs.The crystallite refinement can be attributed to the circulation of crushment and agglomeration during the MA process.Reasons for the increment of lattice strain include size mismatch effect between the constituent elements,increasing grain boundary fraction and high dislocation density imparted by MA[21].The crystallite refinement and high lattice strain may account for the above-mentioned intensity increment and peak broadening in the diffraction.The SEM images of the Al0.5CrFeNiCo0.3C0.2HEA powders of dif-ferent milling time are shown in Fig.2.The primitive powder shows a granular size of less than40l m.In the early period of MA(as shown in Fig.2(b)),the particles cold weld together to form even larger particles than that of primitive powder.Subsequently, when the milling time reaches15h(Fig.2(c)),most of the cold welded agglomerations are crushed down to smaller particles. The27h milled alloy powder reveals an average particle size of approximately5l m as shown in Fig.2(d),and the particles cold weld again when the milling time reaches38h(Fig.2(e)).This cir-culation of crushing and cold welding induced by the ball mill gradually reduces the crystalline size and facilitates the diffusion and alloying among different elements.The38h milled alloy pow-der is subsequently wet ball milled for4h with alcohol as the mill-ing media.It can be seen from Fig.2(f)that the elliptoid particles are fractured,exhibiting a lamellar structure.The particle size for thefinal powder is much smaller than that of38h milled powder.1.XRD patterns of Al0.5CrFeNiCo0.3C0.2HEA powders with different milling time.Table1The crystalline size and lattice strain of Al0.5CrFeNiCo0.3C0.2HEA during MA.Milling time(h)Crystalline size(nm)Lattice strain(%)BCC FCC BCC FCC 3812130.710.734211120.760.77974S.3.2.Phase evolution and microstructure after SPSFig.3illustrates the XRD pattern of Al0.5CrFeNiCo 0.3C 0.2HEA detailed analysis of the XRD pattern suggests a BCC phase,a Cr 23C 6carbide phase and an visible.Except the ordered BCC phase,the other three phases are calculated to be (BCC),10.652Å(Cr 23C 6),respectively.After the FCC phase reduces from 0.77%after MA phase from 0.76%to 0.27%,which confirmsthe nearly annihilation of defects introduced during MA after sin-tering.As hereinbefore mentioned,the main phases of Al 0.5CrFe-NiCo 0.3C 0.2HEA powder after 42h of milling are the BCC and FCC solid-solution phases,demonstrating that densification at 1273K for 8min has resulted in phase evolution.The MA process could in-duce large strain and defects which might lead to the extension of solubility,thus the milled powders are generally in a non-equilib-rium state.It is certain that a reordering process happened during SPS,leading to the metastable supersaturated solid solutions of MA to more stable phases.This reordering can be correlated with the above-mentioned annihilation of defects introduced by severe plastic deformation during MA.Furthermore,the huge pulsed elec-tric current during the SPS process can also facilitate the phase evolution.Fig.4shows the SEM micrographs of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA consolidated by SPS.Few porosities can be seen in low magni-fied image (Fig.4(a)).Actually,relative density of the sintered Al 0.5-CrFeNiCo 0.3C 0.2HEA sample,which is calculated with respect to theoretical density,reaches up to 99.6%.Two distinctive areas are visible in Fig.4(a),viz.,irregular bulk areas and irregular pot hole areas.The irregular bulk areas are most likely to be the main phase because they account for a higher volume fraction.While the pot hole areas can be consist of phases which are removed by corrosive agent.High magnified image (Fig.4(b))shows that some nanoscale black spots are dispersed in irregular pot hole areas.Most of these spots might be numerous nanoscale phases,and a small minority of spots might be ultrafine porosities.In summary,bulk Al 0.5CrFe-NiCo 0.3C 0.2HEA might consist of at least three types of phases.The TEM bright field image and corresponding selected area electron diffraction (SAED)patterns of Al 0.5CrFeNiCo 0.3C 0.2HEAAl 0.5CrFeNiCo 0.3C 0.2HEA powders with different milling time.(a)0h,(b)6h,(c)15h,(d)27h,(e)38h and (f)Fig.3.XRD pattern of Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.after SPS are shown in Fig.5.In order to confirm the phase compo-sition of Al 0.5CrFeNiCo 0.3C 0.2HEA,EDS/TEM method was used to measure chemical compositions (in at.%)of the phases in Fig.5(a)(regions marked A–Y).Results of the chemical composi-tion analysis are listed in Table 2.It is noticed that all the regions can be classified as four different phases.Crystal structures of these phases are measured by their corresponding SAED patterns as shown from Fig.5(b–e).As a result,these four phases are measured to be an FCC phase,a BCC phase,Cr 23C 6carbide and an ordered BCC phase,respectively.It is worth mentioning that diffraction pattern of region C exhibits a FCC twin structure of the FCC phase.Clear microstructures and corresponding SAED patterns of twins will be illustrated hereinafter.As shown in Table 2,the FCC phase isdepleted in Al,Cr compared with the nominal CrFeNiCo 0.3C 0.2HEA.The BCC phase is Ni-rich,Fe-depleted and C-depleted,while the or-greatly rich in Al.The lattice parameters of measured by SAED are 3.700Å(FCC),Å(Cr 23C 6)and 4.320Å(Ordered BCC),Cr 23C 6phase also presents Fe,Ni,Co and Al elements,especially Fe shows a high concentration (12.7±0.4at.%).In conclusion,the bulk Al 0.5CrFeNiCo 0.3C 0.2HEA exhibits a Fe–Ni-rich FCC phase,a Ni ÀAl-rich BCC phase,a Cr 23C 6carbide phase and a Al-rich ordered BCC phase,which is in accordance with the XRD result (Fig.3).It is noticeable that the grain size of these phases shows a wide distri-bution ranging from several hundred nanometers to 1l m.How-ever,the Cr 23C 6and the ordered BCC phases are much finer than the FCC and the BCC phases.It is a reasonable explanation that the Cr 23C 6and the ordered BCC phases are formed in a phase evo-lution and reordering process during SPS.The formation of these four phases is complicated.According to the Gibbs phase rules,the number of equilibrium phases (p )is p =n +1for the alloy that contains n elements.Since phase formation not in equilibrium conditions,the number p >n +1.However,the Al 0.5CrFeNiCo 0.3C 0.2four types of phases,showing much fewer Yeh et al.[2]has proposed that it is possible solutions when HEAs contain five or more 4.SEM micrographs of Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.(a)low-magnified image,and (b)high-magnified image.corresponding SAED patterns of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.(a)bright field image,(b)SAED pattern of FCC [011]1]zone axis (Region D),(d)SAED pattern of Cr 23C 6[112]zone axis (Region B),and (e)SAED pattern of ordered BCC [012]Table 2Chemical composition (in at.%)analysis results of the phases by EDS/TEM.RegionsPhases Cr Fe Ni Al Co C A,C,F,G,I,J,K,M,N,O,P,T,V,W,X,Y FCC 21.8±0.931.8±1.229±1.1 4.4±0.68.5±0.8 4.5±0.5D,E,Q,R BCC 19.7±1.218.6±2.032.8±1.217.5±1.87.2±0.5 4.2±0.4B,L Cr 23C 663.5±2.112.7±0.4 2.9±0.4 1.8±0.3 2.6±0.316.5±1.6H,S,UOrdered BCC 1.6±0.5 2.4±0.42±0.486.8±1.7 1.2±0.46±1.0Nominal composition–25252512.57.55are randomly distributed in the crystal lattice,though the alloy exhibits high mixing entropy.It reveals that high mixing entropy is insufficient to dominate the formation of phases in HEAs sys-tems.Zhang et al.[32]related the formation of simple solid solu-tion to the mixing enthalpy (D H mix )and atomic size difference (d ).Subsequently,Yang and Zhang [33]proposed a solid-solution formation rule for multi-component HEAs based on the calculation of most of reported HEAs.According to Yang et al.,two parameters can be used to estimate the phase formation behavior of HEAs:O is defined as a parameter of the entropy of mixing timing the average melting temperature of the elements over the enthalpy of mixing,d is defined as the mean square deviation of the atomic size of ele-ments.It is proposed that HEAs form simple crystal structures when O P 1.1and d 66.6%.These two parameters are defined by Eqs.(1)and (2),respectively.X ¼T m D S mix j D H mix jð1Þd ¼ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiX n i ¼1c i ð1Àr i = r Þ2q ð2Þwhere c i is the atomic percentage of the i th component, r ¼Pn i ¼1c i r i is the average atomic radius and r i is the atomic radius of the i th component.T m ,D S mix and D H mix are calculated as follows:T m ¼Xn i ¼1c i ðT m Þið3ÞHere,(T m )i is the melting point of the i th component of alloy.D S mix¼ÀR Xn i ¼1ðc i ln c i Þð4Þwhere c i is the mole percent of component,P ni ¼1c i¼1,and R(=8.314JK À1mol À1)is gas constant.D H mix ¼X n i ¼1;i –jX ij c i c j ð5Þwhere X ij ¼4D H mix ijis the regular solution interaction parameter between the i th and j th elements,c i or c j is the atomic percentage of the i th or j th component,and D H mix ijis the enthalpy of mixing of binary liquid alloys.Table 3presents the value of mixing enthal-py of atom-pairs of Al 0.5CrFeNiCo 0.3C 0.2HEA.The calculated results of O and d for the alloy are 1.41and 10.01%,respectively.It is obvi-ous that O well matches while d breaks the solid-solution formation rules for the multi-component high entropy alloys proposed by Yang et al.This phenomenon indicates that the formation of Cr 23C 6carbide and ordered BCC phase is reasonable in this alloy sys-tem and it well conforms to Yang’s research.Firstly,it is obvious that the value of O shows the relative predominance of T m D S mix and D H mix .Since O >1is obtained,T m D S mix should be the predomi-nant part of the free energy.It is known that mixing entropy indi-cates the tendency for the formation of random solution while enthalpy of mixing in any system indicates the tendency for order-ing or clustering,thus the formation of solid-solution phases ought to be much easier than the formation of intermetallic compoundsand other ordered phases in the Al 0.5CrFeNiCo 0.3C 0.2HEA.Obviously,the formation of FCC solid-solution phase as the main phase of Al 0.5-CrFeNiCo 0.3C 0.2HEA is mainly correlated with the alloy’s high mix-ing entropy.Secondly,the relatively large value of d indicates that the atomic size difference between components is too large for this alloy system to form entire solid-solution phases.Because the large atomic size mismatch could lead to serious lattice distortion and subsequently increase the corresponding strain energy which could lower the stability of solid-solution.Actually,the interstitial solubil-ity of carbon element in alloy is quite limited,so that carbon ele-ment has a strong tendency to exist in the form of carbonization or graphite in alloy.Moreover,the diffusion of atoms in the matrix could be suppressed due to large atomic size difference.Thus it facilitates the atomic segregation,even results in the formation of amorphous structures [35].As listed in Table 3,atomic radiuses of C (0.77Å)and Al (1.43Å)are significantly different from the other elements.This could be a reasonable explanation for the formation of the Cr 23C 6carbide and the Al-rich ordered BCC phase.In addition,it can also be found that the mixing enthalpy of most of the atomic pairs is highly negative (shown in Table 3).Al and C are the main contributor to the negative enthalpy of mixing in this alloy,indicat-ing their strong tendency for ordering or clustering.It is interesting to note that the mixing enthalpy of C and Cr (À61kJ/mol)is the most negative.This could be another factor which contributes to the formation of the Cr 23C 6carbide in the alloy.In addition,as above-mentioned,Fe has a high concentration (12.7±0.4at.%)in Cr 23C 6phase,revealing that Fe has high solubility in this type Cr 23C 6phase.According to Table 3,the mixing enthalpy of C and Fe (À50kJ/mol)stays a highly negative level,leading to Fe atoms preferring C sites.Mechanically alloyed powders of Al 0.5CrFeNiCo 0.3-C 0.2HEA exhibit simple solid solution structure within 38h of mill-ing,which can be attributed to the formation of supersaturated solid solution.The solid solubility extension can be ascribed to the high mixing entropy effect as well as the non-equilibrium state of the MA process [36].Interestingly,twinned FCC phase is also ob-served.Fig.6shows TEM images and corresponding SAED patterns of the FCC phase with nanoscale twins.It is worth pointing out that nanoscale twins are found only in the FCC phase,which has been confirmed by EDS/TEM and the corresponding diffraction patterns.The lamella thickness of the nanoscale twins shown in Fig.6(a)is less than 60nm.The corresponding SAED pattern (the matrix axisis [011]M and the twin axis is ½0 1 1 T)which is shown in the inset of the upper right of Fig.6(a)indicates the nanoscale twins belong to the FCC phase.Actually,existing researches in deformation twin-ning of nanocrystalline materials are mainly focus on FCC crystal structure metals,which is probably attributed to their favorable capability of deformation twinning [37].Fig.6(b)illustrates a noticeable twin structure of the FCC phase surrounding by grey phases.These grey phases are found to be the Cr 23C 6carbide through EDS/TEM analysis,revealing that the formation of Cr 23C 6might have effect on deformation twinning.Since Cr 23C 6shows a hard brittle texture [38],it could be presumed that partial FCC phase between the Cr 23C 6carbide is not readily deformed and con-solidated.Under the isostatic pressure of 30MPa during the SPS process,certain parallel crystal faces of partial crystals moving opposite to each other along a direction with a certain value of dis-placement distance.Thus twinning in the FCC phase may occur dur-ing the phase evolution and densification with the aim to be more stable and attaining complete densification [29].Moreover,a cer-tain twin system could be activated by a critical resolved shear stress which can be achieved during crystal deformation and phase evolution process.It is worth pointing out that similar nanoscale twins in CoNiFeAl 0.6Ti 0.4and CrCoNiFeAl 0.6Ti 0.4HEAs prepared by MA and SPS have been observed in previous studies of our research group [28,29].Table 3The chemical mixing enthalpy D H mix ij ;kJ =mol of binary equiatomic alloys calculated by Miedema’s approach [34].Element (atomic sizes,Å)C Ni Cr Co Fe Al C (0.77)–À39À61À42À50À36Ni (1.24)––À70À2À22Cr (1.25)–––À4À1À10Co (1.26)––––À1À19Fe (1.26)–––––À11Al (1.43)––––––S.Fang et al./Materials and Design 54(2014)973–9799773.3.Mechanical propertiesFig.7shows the room-temperature compressive stress–strain curve of bulk Al 0.5CrFeNiCo 0.3C 0.2HEA after SPS.The compressive strength and compression ratio of the alloy are 2131MPa and 3.0%,respectively.The average Vickers hardness of Al 0.5CrFeNiC-o 0.3C 0.2HEA after SPS has been measured to be 617±25HV.How-ever,plastic deformation behavior is not characterized in theindicating a limited compression plasticity of mechanical properties of Al 0.5CrFeNiCo 0.3C 0.2HEA HEAs of AlCrFeNiCo HEA system are listed The listed HEAs are prepared by casting 0.5CrFeNiCo 0.3C 0.2HEA studied in this paper.hardness of these as-cast HEAs has beenreported.Obviously,Al 0.5CrFeNiCo 0.3C 0.2alloy exhibits the highest hardness of the HEAs listed in Table 4.The high strength and high hardness of Al 0.5CrFeNiCo 0.3C 0.2HEA is possibly due to the formation of the ordered BCC phase and the Cr 23C 6carbide,the nanocrystalline structure,as well as solid solu-tion strengthening mechanism of Al atoms.It is worth pointing out that nanoscale twins might also have a considerable effect on strengthening the alloy.The strengthening mechanism of nano-scale twins is associated with the complicated interaction between dislocations and twin boundaries.It is believed that twin bound-aries are effective in blocking dislocation motion,especially when the thickness of twin/matrix lamellae decreases down to the nano-scale,and a Hall–Petch-type relationship exists for twin boundary strengthening [39,40].The limited compression ratio may owing to the brittle Cr 23C 6carbide,the ordered BCC phase,solid solution strengthening of Al atoms and ultrafine porosities.4.ConclusionsAl 0.5CrFeNiCo 0.3C 0.2high entropy alloy with nanocrystalline has been successfully synthesized by MA and SPS.A supersaturated so-lid solution with both FCC and BCC structures is evidently observed after MA.After SPS,Cr 23C 6carbide and an ordered BCC phase are newly formed.The TEM analysis results confirm that the alloy con-sists of one FCC phase,one BCC phase,Cr 23C 6carbide and one or-dered BCC pared with the nominal composition of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy,the FCC phase is enriched in Fe–Ni,the BCC phase is enriched in Ni–Al,while the ordered BCC phase is especially enriched in Al.The grain size of these phases shows a wide distribution ranging from several hundred nanometers to 1l m.Nanoscale deformation twins present only in partial FCC phase.The compressive strength,compression ratio and Vickers hardness of Al 0.5CrFeNiCo 0.3C 0.2high entropy alloy are 2131MPa,3.0%and 617±25HV,respectively.AcknowledgementsThe authors wish to acknowledge the financial support by National Natural Science Foundation of China (Grant No.51271080.)and Project supported by Guangdong Provincial Natural Science Foundation of China (Grant No.S2011010002227.).References[1]Zhao ZD,Chen Q,Chao HY,Huang SH.Microstructural evolution and tensilemechanical properties of thixoforged ZK60-Y magnesium alloys produced by two different routes.Mater Des 2010;31(4):1906–16.[2]Yeh JW,Chen SK,Lin SJ,Gan JY,Chin TS,Shun TT,et al.Nanostructured high-entropy alloys with multiple principal elements:novel alloy design concepts and outcomes.Adv Eng Mater 2004;6(5):299–303.and SAED pattern of twinned FCC phase.(a)a twinned FCC grain with 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Microstructure evolution and mechanical properties of
Microstructure evolution and mechanical properties of1 000 MPa cold rolled dual-phase steelZHAO Zheng-zhi(赵征志), JIN Guang-can(金光灿), NIU Feng(牛枫), TANG Di(唐荻), ZHAO Ai-min(赵爱民) Engineering Research Institute, University of Science and Technology Beijing, Beijing 100083, ChinaReceived 10 August 2009; accepted 15 September 2009Abstract: The microstructure evolution of 1 000 MPa cold rolled dual-phase (DP) steel at the initial heating stages of the continuous annealing process was analyzed. The effects of different overaging temperatures on the microstructures and mechanical properties of 1 000 MPa cold rolled DP steel were investigated using a Gleeble−3500 thermal/mechanical simulator. The experimental results show that ferrite recovery and recrystallization, pearlite dissolution and austenite nucleation and growth take place in the annealing process of ultra-high strength cold rolled DP steel. When being annealed at 800 ℃ for 80 s, the tensile strength and total elongation of DP steel can reach 1 150 MPa and 13%, respectively. The microstructure of DP steel mainly consists of a mixture of ferrite and martensite. The steel exhibits low yield strength and continuous yielding which is commonly attributed to mobile dislocations introduced during cooling process from the intercritical annealing temperature.Key words: cold rolled dual-phase steel; microstructure evolution; recrystallization; mechanical property; overaging temperature1 IntroductionAdvanced high-strength steels (AHSS) have been used in the automotive industry as a solution for the weight reduction, safety performance improvement and cost saving. Among them, the dual-phase (DP) steels, whose microstructure mainly consists of ferrite and martensite, are an excellent choice for applications where low yield strength, high tensile strength, continuous yielding, and good uniform elongation are required [1−4].The continuous annealing process to produce cold rolled DP steels typically has the following stages: heating to the intercritical temperature region, soaking in order to allow the nucleation and growth of austenite, slow cooling to the quench temperature, rapid cooling to transform the austenite into martensite, overaging, and air cooling. The amount and morphology of the constituents formed depend on such annealing parameters. The effects of the retained austenite, ferrite, and martensite morphologies on the mechanical behavior of DP steels have been intensively investigated[5−9]. As we all known, overaging treatment is an important process during the production of dual-phase steel. It can reduce the hardness of martensite and improve the comprehensive mechanical properties of DP steel [10−14].The purpose of the present research was to study the microstructure evolution of cold rolled DP steel at the initial heating stages of the continuous annealing process using a Gleeble simulator. At the same time, the effects of overaging temperature on the mechanical properties of DP steel were also studied. The microstructures of specimens simulated on a Gleeble simulator, were analyzed using scanning electron microscopy (SEM) and transmission electron microscopy (TEM).2 ExperimentalThe chemical compositions of the experimental steel (mass fraction, %) were: 0.14−0.17C, 0.40−0.60Si, 1.70−1.90Mn, 0.02−0.04Nb, 0.40−0.60Cr, ≤0.010P, ≤0.010S, 0.02−0.06Al and balance Fe. Firstly, experimental steels were smelted in a 50 kg vacuum induction furnace. After smelting, experimental steels were forged into 35 mm×100 mm×100 mm cubic samples. The forged slabs were reheated to 1 200 ℃and soaked for 1 h. The hot rolled thickness was 3.5 mm after 6 passes rolling. The finish rolling temperature was about 880 ℃. The coiling temperature was 620 ℃. After being pickled in hydrochloric acid, the hot rolledFoundation item: Project(2006BAE03A06) supported by the National Key Technology R&D Program during the 11th Five-Year Plan Period Corresponding author: ZHAO Zheng-zhi; Tel: +86-10-62332617; E-mail: zhaozhzhi@ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s564bands were cold rolled to the final thickness of 1.0 mm, and the reduction was about 70%. Finally, the cold rolled sheets were cut into the samples for the simulation of continuous annealing experiment.The microstructure evolution at the initial steps of the continuous annealing process was studied using a Gleeble 1500 simulator. The steel was heated at 10 ℃/sto the different heating temperatures (550, 630, 670, 710, 730, 750 and 780 ℃) and held for 20 s followed by water-quenching. The effects of different overaging temperatures on the microstructures and mechanical properties of DP steel were investigated using a Gleeble 3500 simulator. The processing schedules and parameters used are shown in Fig.1. The soaking temperature of intercritical region was set at 800 ℃, soaking time is 80 s; after a slow cooling, the samples were rapidly cooled to 240, 280, 320 and 360 ℃, respectively and soaked for 300 s; at last, the samples were air cooled to the room temperature.Fig.1 Continuous annealing process of DP steelAfter heat treatment, the steel sheet would be cut into standard tensile specimens (length 200 mm, gauge length 50 mm). The tensile test was performed with CMT4105-type tensile test machine to test mechanical properties. The longitudinal cold rolling plane sections of samples after annealing were prepared and etched with 4% natal. The microstructure was analyzed by scanning electron microscopy (SEM). Some samples were analyzed using transmission electron microscopy (TEM).3 Results and discussion3.1 Mechanical properties and microstructures ofsamples after hot-rolling and continuousannealingTable 1 shows the tensile test data for the two samples after hot-rolling and continuous annealing in terms of yield strength, ultimate tensile strength and total elongation. When the annealing temperature is 800 ℃and soaking time is 60 s, the tensile strength reaches 1 110 MPa and the total elongation reaches 12%. Compared with the hot-rolled samples, the yield strength and total elongation of sample after annealing are similar, but the tensile strength increases by about 450 MPa. The yield ratio decreases obviously. The engineering uniaxial tensile stress—strain curve of the sample after continuous annealing is characterized by very uniform plastic flow until necking. There is no physical yield point and yield point extension, that is, the steel exhibits continuous yielding which is commonly attributed to mobile dislocations introduced during cooling from the intercritical annealing temperature. Many dislocation sources come into action at low strain and plastic flow begins simultaneously through the specimen, thereby suppressing discontinuous yielding[15].Table 1 Mechanical properties of samples after hot rolling and annealingConditionYieldstrength/MPaTensilestrength/MPaYieldratio*Totalelongation/% Hot rolling555 665 0.83 16 Annealing540 1110 0.49 12* Yield ratio is defined as the ratio of yield strength to tensile strength.The microstructures of the hot-rolled and cold-rolled samples are shown in Fig.2. It can be observed that hot rolled steel features a band microstructure, i.e. pearlite band in a ferrite grain matrix. The ferrite grain size is measured to be 5.0−9.0 µm. After cold rolling, the microstructure consists of elongated grains of ferrite and deformed colonies of pearlite (Fig.2(b)). After cold-rolling, there is an increase in the stored energy of the steel due to the high dislocation density and this provides the driving pressure for the ferrite recrystallization during annealing process. The total ferrite grain boundary area increases and the cementite laminar structure in pearlite is broken down. The latter has been shown to promote spheroidization of cementite during subsequent annealing process.The SEM micrograph of the sample after annealing is given in Fig.3(a). The microstructure of DP steel consists of a mixture of ferrite, martensite, martensite/austenite constituent. There is also some bainite in the microstructure. The martensite islands are homogeneously distributed in ferrite matrix. The DP steel has finer grain size and the size of ferrite grain and martensite island are about 1.0−2.0 µm. Some martensite islands have a bright white circle around the edge, and the center of martensite is of irregular black structure.ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s565Fig.2 Microstructures of steel after hot rolling (a) and cold rolling (b)Fig.3 SEM images (a) and TEM micrograph (b) of steel after continuous annealingThe main reason is the manganese partitioning will occur during the continuous annealing process. During the heating process, a high-Mn side lap forms around austenite, which makes the hardenability of austenite island edge higher than that of the center. So, it makes high-Mn side lap form around martensite in the cooling process. The volume fraction of martensite is about 40%, which is the main reason for DP steel with a higher strength. After the continuous annealing process, band structure is significantly improved, which plays an important role in improving the performance of DP steel.The fine structures of martensite and ferrite are shown in Fig.3(b) by the TEM observation. The lath martensite is fine, and is relatively clean; at the same time, a very high density of dislocations can be observed in the ferrite grain adjacent to martensite. These dislocations are generated in order to accommodate transformation induced strain built between martensite transformed by quenching and retained ferrite. In addition, they are known to be mobile and play an important role on rapid, extensive strain hardening of DP steel from the onset of its plastic deformation.3.2 Microstructure evolution at initial steps ofcontinuous annealing processThe microstructure evolution at the initial stages of the continuous annealing process is very important for producing the ultra-high strength DP steel. During the annealing process of high strength DP steel, ferrite recovery and recrystallization, pearlite dissolution and austenite nucleation and growth will occur. When the sample is heated to 550 , the℃microstructure has no visible change as compared with the cold rolled sample. The ferrite grain is stretched along the rolling direction significantly; lamellar pearlite is stretched along the rolling direction too. At the same time, there are some carbide particles in the ferrite matrix, as shown in Fig.4(a). At this temperature, the recrystallization nucleus was not found in the structure. So, at this stage the sample is still at the recovery stage. When the heating temperature is 630 , the℃recrystallization nucleus begins to appear in the microstructure. The nucleus of crystal appears mainly nearby the large deformation ferrite (Fig.4(b)). The recrystallization nucleus is fine and equiaxed. Large deformation storage power is present in the large deformation region. So, recrystallization nucleus forms in this region firstly. With the heating temperature increasing, the recrystallization nucleus begins to grow. Therefore, the size of recrystallization is uneven at this stage, as shown in Fig.4(c). When the heating temperature is 670 ℃, the deformation structure still exists in the microstructure. With the temperature increasing, the deformed ferrite grains are replaced by recrystallization ferrite grains. When the heating temperature is 710 , the d℃eformation structure has already vanished, which is replaced by theZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s566equiaxed recrystallization grain. So, the process of recrystallization completes basically. In the ferrite recrystallization process, the pearlite transforms to granular from lamellar gradually.When the heating temperature is 730 ,℃it begins to enter the two-phase region; and the ferrite and spheroidised carbides begin to transform to austenite. A small amount of austenite nucleates in the original pearlite region, as shown in Fig.4(e). Austenite nucleates mainly in the ferrite and pearlite grain boundary; and a part of austenite also nucleates in the carbide particles of ferrite. After austenite nucleation, it begins to grow rapidly. At this stage, the pearlite dissolves rapidly. When the temperature reaches 750 , the austenite℃transformation occurs obviously. The bright white particle which distributes in the ferrite matrix is the martensite island. The martensite transforms from austenite during the rapid cooling process. At the same time, a small amount of martensite particles can also be observed in ferrite; and there are still some non-dissolved carbide particles in the ferrite matrix. The initial austenite growing-up is mainly controlled by the carbon Fig.4Microstructure evolutions duringcontinuous heating process: (a) 550 ℃; (b)630 ℃; (c) 670 ℃; (d) 710 ℃; (e) 730 ℃; (f)750 ℃; (g) 780 ℃ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s567diffusion in the austenite, and the diffusion path is along the pearlite/austenite interface. When the annealing temperature is 780 , the austenite volume increase℃s, and the number of carbide particles is reduced gradually. There is only a very small amount of carbide particles distributing in ferrite matrix.3.3 Effect of overaging temperature onmicrostructure and mechanical properties ofDP steelThe overaging is a temper treatment to harden martensite in the dual-phase steel, reduce the hardness of martensite and improve the comprehensive mechanical properties[16]. Fig.5 shows the effect of overaging temperature on the mechanical properties of dual-phase steel. All the samples are intercritically annealed at 800℃ with different overaging temperatures. As can be seen from Fig.5, the highest tensile strength is achieved in the sample overaged at 280 ℃. The yield strength is 560 MPa, the tensile strength is 1 150 MPa, and the total elongation reaches 13%. The good combination of high strength and toughness properties is obtained. And then, with the increase of overaging temperature, the yield strength and tensile strength of samples decrease, while the total elongation increases. When the overaging temperature reaches 360 ℃, the tensile strength decreases, the yield strength does not change significantly. The mechanical properties of sample cannot meet the necessary requirements of CR980DP. At the same time, the stress—strain curve of the steel shows discontinuous yielding behaviour and develops yield plateaus.Fig.6 shows the SEM microstructures with different overaging temperatures. It can be seen that the microstructure mainly consists of dark grey ferrite grains and white martensite. When the overaging temperature is 360 ℃, the martensite boundary is fuzzier than that of sample overaged at 320 ℃, and there are more carbides, which is due to the effects of tempering on the martensite, such as the volume contraction of martensite during the tempering, the changes of the martensite strength and additional carbon clustering or precipitation near the ferrite and martensite interfaces.Fig.5 Effects of different overaging temperatures on mechanical propertiesFig.6 SEM images of microstructures of DP steel overaged at different temperatures: (a) 240 ℃; (b) 280 ℃; (c) 320 ℃; (d) 360 ℃ZHAO Zheng-zhi, et al/Trans. Nonferrous Met. Soc. China 19(2009) s563−s568 s5684 Conclusions1) When the DP steel is annealed at 800 ℃ for 80 s and overaged at 280 ℃, the tensile strength and total elongation of ultra-high strength dual-phase steel can reach 1 150 MPa and 13%, respectively.2) The microstructure of DP steel consists of a mixture of ferrite, martensite, martensite/austenite constituent. There are also some bainites in the microstructure. The martensite islands are homogeneously distributed in ferrite matrix. The ferrite and martensite island grain size are about 1.0−2.0 µm. When the overaging temperature reaches 360 ℃, the tensile strength decreases, the yield strength does not change significantly. The mechanical properties of sample cannot meet the necessary requirements of CR980DP. At the same time, the steel shows discontinuous yielding behaviour and develops yield plateaus.References[1]KANG Yong-lin. Quality control and formability of the mordernMotor plate [M]. 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Microstructure and mechanical properties of
Microstructure and mechanical properties ofZrB 2–SiC–ZrO 2f ceramicLin Jia,a ZhangXinghong,a ,⇑Wang Zhi ,b ,⇑and Han Wenbo aaNational Key Laboratory of Science and Technology on Advanced Composites in Special Environments,Harbin Institute of Technology,Harbin 150001,PR ChinabSchool of Aeronautics and Astronautics,Faculty of Vehicle Engineering and Mechanics,State Key Laboratory of Structural Analysis for Industrial Equipment,Dalian University of Technology,Dalian 116024,PR ChinaReceived 10January 2011;accepted 12January 2011Available online 15January 2011ZrB 2–SiC–ZrO 2f ceramic was fabricated by hot-pressing at 1850°C for 1h under a uniaxial load of 30MPa in vacuum.The ZrB 2–SiC–ZrO 2f ceramic thus produced showed excellent fracture toughness due to the addition of ZrO 2fiber.The increase in toughness was attributed mainly to the addition of the ZrO 2fiber,which could enhance fiber pull-out,crack bridging and crack branching.In addition,the stress-induced transformation toughening was also considered to be main reason for the improvement in toughness.Ó2011Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Ceramics;Microstructure;Mechanical properties;CompositesZirconium diboride (ZrB 2)is one of the families ofmaterials known as ultra-high-temperature ceramics [1].Because of strong covalent bonding and low self-diffu-sion,high temperatures and external pressures are re-quired to densify monolithic ZrB 2[2].In previous studies,nominally stoichiometric ZrB 2without additives has only been densified by hot-pressing at 2000°C or higher with pressures of 20–30MPa,or at reduced tem-peratures (1790–1840°C)with much higher pressures (800–1500MPa)[3].Recent studies have shown that the addition of SiC particles improves the densification of ZrB 2by maintaining a fine grain size and a uniform distri-bution of the reinforcing phase,and enhances oxidation resistance by promoting the formation of silicate-based glasses that inhibit oxidation at temperatures between 800and 1700°C [4].Unfortunately,unsatisfactory frac-ture toughness is still the obstacle preventing ZrB 2–SiC ceramic from being widely used,especially for applica-tions in thermal shock conditions with high heat transfer and/or rapid environmental temperature changes,such as furnace elements,plasma arc electrodes,hypersonic air-craft,reusable launch vehicles,or rocket engines and ther-mal protection structures for leading edge parts onhypersonic reentry space vehicles [1–3,5].One method for improving thermal shock resistance is to tailor the structure on multiple length scales to produce architec-tures that are engineered to enhance thermal shock resis-tance while maintaining load-bearing capability—zirconia (ZrO 2)particles [6],graphite flakes [7],carbon fi-bers [8]or SiC whiskers [9]have been used as toughening materials.In addition,among these materials ZrO 2is of particular interest as it undergoes at least three crystallo-graphic transformations (monoclinic,tetragonal and cu-bic)when it cools from high temperature to room temperature,shown as the following transformation [10]:222C 2370C 2370C950C1170ZrO c ZrO t ZrO m −°⎯⎯←⎯→⎯°−°⎯⎯←⎯→⎯°−It has been shown that yttria-stabilized tetragonal zir-conia polycrystals (Y-TZP)possess superior fracture toughness.The high toughness of the Y-TZP monoliths arises from the volume expansion (4–5%)because of the stress-induced martensitic phase transformation of tetragonal to monoclinic symmetry in the stress field of propagating cracks,known as transformation tough-ening [11,12].Furthermore,fiber is found to be effective in strengthening and toughening ceramic materials,and1359-6462/$-see front matter Ó2011Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.scriptamat.2011.01.019⇑Correspondingauthor.Tel./fax:+8645186402382;e-mail:jiajia10182003@Scripta Materialia 64(2011)872–875/locate/scriptamatthe toughening is mainly the result of three mechanisms:fiber bridging,fiber pullout and crack deflection [13,14].It was expected that the toughness can further be improved when both fiber-toughening and phase transformation are incorporated into the ceramic ma-trix.Nevertheless,there are few reports in the open liter-ature on the cooperative toughening of fibers and transformation toughening of ZrO 2fiber-toughened ZrB 2–SiC ceramics.In this study,in order to improve further the fracture toughness of ZrB 2–SiC ceramic so that this material can be used more widely in severe environments,ZrB 2–SiC ceramic toughened with ZrO 2fiber was fabricated by hot pressing.The microstructure and the mechanical properties of the ZrB 2–SiC–ZrO 2f ceramic were investi-gated and are discussed in detail.The purpose of this paper is to take the lead in reporting ZrB 2–SiC ceramic toughened by ZrO 2fiber,and this study clearly showed that the addition of ZrO 2fiber into ZrB 2–SiC ceramic is a promising way to improve the fracture mercially available ZrB 2powder (2l m,>99.5%,Northwest Institute for Non-ferrous Metal Research,PR China),SiC (1l m,>99.5%,Weifang Kaihua Mi-cro-powder Co.Ltd.,PR China)and ZrO 2fiber (mean diameter 5–8l m,mean length 200l m,>99%,Shan-dong Huolong Ceramic Fiber Co.Ltd.,PR China)were used as raw material.The powder mixtures of ZrB 2plus 20vol.%SiC plus 15vol.%ZrO 2fiber (ZrB 2–SiC–ZrO 2f )were ball-mixed for 20h in a polytetrafluoroethylene bottle using ZrO 2balls and ethanol as the grinding med-ia.After mixing,the slurry was dried in a rotary evapo-rator.The resulting powder mixtures were hot-pressed at 1850°C for 1h under a uniaxial load of 30MPa in vacuum.The microstructural features of the hot-press-ing composite were observed by scanning electron microscopy (SEM;FEI Sirion,Holland)with simulta-neous chemical analysis by energy-dispersive spectros-copy (EDS;EDAX Inc.).The phase composition was determined by X-ray diffraction (XRD;Rigaku,Dmax-rb,Cu K a =1.5418A˚).The bulkspecimens was measured by the Archimedes Flexural strength (r )was tested by three-point on 3mm Â4mm Â36mm bars,using a 30and a crosshead speed of 0.5mm min À1.Each was ground and polished with diamond slurries a 1l m finish.The edges of all the specimens fered to minimize the effect of stress to machining flaws.Hardness (Hv 0.5)was by Vickers’indentation with a 4.9N load 10s on polished sections.Fracture was evaluated by a single-edge notched beam a 16mm span and a crosshead speed of 0.05using 2mm Â4mm Â22mm test bars,on the used for the flexural strength measurements.All and fracture bars were cut with the tensile pendicular to the hot-pressing direction.A number of six specimens were tested for each tal condition.The hot-pressed ZrB 2–SiC–ZrO 2f ceramic to produce bars for mechanical property test sured bulk densities of 5.44g cm Àing a ture calculation,and assuming that the true were 6.09g cm À3for ZrB 2, 3.21g cm À3for 6.11g cm À3for ZrO 2[1,9],the theoretical density of the ZrB 2–SiC–ZrO 2f ceramic was calculated to be 5.52g cm À3.Based on this true density,the relative den-sity of the ZrB 2–SiC–ZrO 2f ceramic was as high as 98.6%.An XRD spectrum obtained from the fractured and polished surface of the ZrB 2–SiC–ZrO 2f ceramic is shown in Figure 1.Phase analysis indicates that the pre-dominant phases of this hot-pressed ZrB 2–SiC–ZrO 2f ceramic are ZrB 2,SiC and t-ZrO 2as well as a trace of ZrC on the polished surface of the ZrB 2–SiC–ZrO 2f ceramic.The formation of ZrC was attributed to the reaction of SiC with ZrO 2:2ZrO 2ðs Þþ3SiC ðs Þ¼2ZrC ðs Þþ3SiO ðg ÞþCO ðg Þð1ÞThe gaseous products,i.e.CO and SiO,were readily removed by the high vacuum (pressure 5Pa),which is thermodynamically favorable to reaction (1).Further-more,the very fine ZrC particles formed in situ on the surface of the ZrO 2and SiC particles were highly sinter-able,which could also provide a higher driving force for sintering as densification is driven by minimization of surface free energy,which was also thought to be main reason responsible for the high densification [3].In addi-tion,the effect of the ZrC phase on the relative density was not considered because the exact content of this phase was not calculated,and moreover the density of the ZrC phase is close to the densities of the ZrB 2and ZrO 2phases [1,3].As seen in Figure 1,the diffraction peak of m-ZrO 2phase was observed on the fracture sur-face of the ZrB 2–SiC–ZrO 2f ceramic.It is known that when subjected to external load,stress concentration in the hot-pressed ZrB 2–SiC–ZrO 2f ceramic will result in the phase transformation from t-ZrO 2to m-ZrO 2with volume change [10].According to the formula of Toraya et al.[15],the volume fraction of the m-ZrO 2(Vm )was calculated by measuring the intensities of the (111)and (11 1)reflec-tions of the monoclinic phase and the (111)peak of the tetragonal phase:L.Jia et al./Scripta Materialia 64(2011)872–875873where Xm is the integrated intensity ratio,and Im and It are the peak intensities of the m-ZrO 2and t-ZrO 2,respectively.Furthermore,the obtained Vm was individ-ually normalized to the volume fraction of ZrO 2(V ZrO 2)in each composite as follows:Vmtot ¼Vm ÂV ZrO 2100ð4ÞTherefore,the result of V mtot on the fracture surface minus that on the polished surface equals the fraction transformed from t-ZrO 2to m-ZrO 2during fracture (i.e.t-ZrO 2transformability).The obvious volume expansion upon phase transformation from t-ZrO 2to m-ZrO 2was calculated,and found to favor an increase in the fracture toughness of the ZrB 2–SiC–ZrO 2f ceramic.SEM images of the polished surface of the ZrB 2–SiC–ZrO 2f ceramic are presented in Figure 2.It was con-firmed by EDS analysis (not shown here)that the small dark phase was SiC which as dispersed uniformly in the lighter ZrB 2matrix;the rod-like phase was ZrO 2fiber.A uniform distribution of short ZrO 2fibers in ZrB 2matrix can be seen in Figure 2B.It was expected that the slight reaction of SiC grain with ZrO 2fiber occurred due to the lower hot pressing temperature and the lower content of SiC grains.Mechanical scratches on ZrO 2fibers andsize was estimated by measuring at least 120grains,and found to be 4.5and 3.4l m for ZrB 2and SiC,pared with raw particles,the growth of ZrB 2grains was inhibited by SiC grain because of the reaction of SiC grains with trace oxide impurities on the ZrB 2particle surfaces [4].It can be seen from the insert in Figure 3A that a perfect interface between ZrO 2fiber and other phases was observed,which also indicated that there was no obvious reaction between SiC grains and ZrO 2fibers.The perfect interface of ZrO 2fiber,SiC and other phases enhanced the mechan-ical properties of the ZrB 2–SiC–ZrO 2f ceramic [3].Sig-nificant pits and fiber roots occurred in the fracture surface of the ZrB 2–SiC–ZrO 2f ceramic,which indicated that the ZrO 2fibers were pulled out during the fracture process,as shown in Figure 3B.In order to further investigate the effect of ZrO 2fibers on the crack propagation models,typical crack propa-gation paths were derived using Vickers’indentation method as shown in Figure 4.The radial crack at the edge of Vickers’indentation clearly revealed that the crack propagation models caused by addition of ZrO 2fiber mainly included crack branching and crack bridg-ing.It is believed that such reaction depletes the energy of crack propagation during fracture and leads to the improvement of the fracture toughness [17].874L.Jia et al./Scripta Materialia 64(2011)872–87524±0.9GPa for ZrB2–SiC ceramics but statistically equivalent to19.3±0.4GPa of the similar ZrB2–SiC–ZrO2p ceramic[4,6].It has been recognized that the hardness of a material is generally decreased by the addition of weak second phases,such as carbon/graph-ite,h-BN and pores[18].Compared with the ZrB2–SiC ceramics[4],the reduction in the hardness of the ZrB2–SiC–ZrO2f ceramic was ascribed to its lower rela-tive density.The measured fracture toughness of the ZrB2–SiC–ZrO2f ceramic ranged from 6.3to7.6 MPaÁm1/2(average 6.8±0.6MPaÁm1/2),which was obviously higher than reported results for monolithic ZrB2(2.3–3.5MPaÁm1/2)and ZrB2–SiC composites (4.0–5.3MPaÁm1/2)[3,4].Furthermore,the fracture toughness of the ZrB2–SiC–ZrO2f ceramic was obvi-ously higher than6.0±0.2MPaÁm1/2of similar ZrB2–SiC–ZrO2p ceramic reported in the literature[6].Com-bined with XRD and SEM analysis,the improvement in strength and toughness was attributed mainly to the phase transformation toughening,fiber pull-out,crack bridging and branching,because these interaction effects absorb fracture energy.In conclusion,the ZrB2–SiC–ZrO2f ceramic was hot-pressed at1850°C for1h under a uniaxial load of 30MPa in vacuum.The relative density of the ZrB2–SiC–ZrO2f ceramic was calculated to be98.6%and the XRD spectra indicated ZrO2phase transformation from tetragonal to monoclinic symmetry.Theflexural strength of the ZrB2–SiC–ZrO2f ceramic was1085±118MPa, which is higher than788±78MPa of the similar ZrB2–SiC–ZrO2p ceramic.The hardness of the ZrB2–SiC–ZrO2f ceramic was found to be as high as18.4±1.3GPa,which is slightly lower than the value of about24±0.9GPa found for ZrB2–SiC ceramics but statistically equivalent to19.3±0.4GPa of the similar ZrB2–SiC–ZrO2p cera-mic.It was recognized that the hardness of a material was in general decreased by the addition of weak second phases,such as carbon/graphite,h-BN and -pared with the ZrB2–SiC ceramics,the reduction in the hardness of the ZrB2–SiC–ZrO2f ceramic was mainly as-cribed to its lower relative density.The measured fracture toughness of the ZrB2–SiC–ZrO2f ceramic ranged from 6.3to7.6MPaÁm1/2(average6.8±0.6MPaÁm1/2),which increased by approximately134%compared to the re-ported results of monolithic ZrB2(2.3–3.5MPaÁm1/2), 46%for ZrB2–SiC composites(4.0–5.3MPaÁm1/2), and13%for the similar ZrB2–SiC–ZrO2p ceramic (6.0±0.2MPaÁm1/2).The observed toughening mecha-nisms were attributed tofiber pull-out,crack bridging, crack branching and phase transformation toughening. This study clearly showed that the addition of ZrO2fiber to ZrB2–SiC ceramic is a promising way to improve the fracture toughness of this material.This work was supported by the NSFC (51072042,10725207),the Science Fund for Outstanding Youths of Heilongjiang Province and China Postdoc-toral Science Foundation Funded Project(2010048 1220).[1]X.H.Zhang,Q.Qu,J.C.Han,W.B.Han,C.Q.Hong,Scripta Mater.59(2008)753.[2]X.H.Zhang,W.J.Li,C.Q.Hong,W.B.Han,J.C.Han,Scripta Mater.59(2008)1214.[3]W.G.Fahrenholtz,G.E.Hilmas,I.G.Talmy,J.A.Zayko-ski,J.Am.Ceram.Soc.90(5)(2007)1347.[4]A.L.Chamberlain,W.G.Fahrenholtz,G.E.Hilmas,D.T.Ellerby,J.Am.Ceram.Soc.87(6)(2004)1170.[5]D.W.Ni,G.J.Zhang,Y.M.Kan,Y.Sakka,ScriptaMater.60(2009)615.[6]X.H.Zhang,W.J.Li,C.Q.Hong,W.B.Han,J.C.Han,Mater.Lett.62(2008)2404.[7]X.H.Zhang,Z.Wang,X.Sun,W.B.Han,C.Q.Hong,Mater.Lett.62(2008)4360.[8]F.Y.Yang,X.H.Zhang,J.C.Han,S.Y.Du,J.AlloysCompd.472(2009)395.[9]X.H.Zhang,L.Xu,S.Y.Du,W.B.Han,J.C.Han,C.Y.Liu,Scripta Mater.59(2008)55.[10]R.C.Garvie,R.H.J.Hannink,R.T.Pascoe,CeramicSteel,Nature,London,1975.[11]C.L.Yang,H.I.Hsiang,C.C.Chen,Ceram.Int.31(2005)297.[12]G.A.Gogotsi,V.I.Galenko,S.P.Mudrik,B.I.Ozersky,Ceram.Int.36(2010)345.[13]B.Budiansky,Y.Q.L.Cui,Mech.Mater.21(1995)139.[14]J.P.Singh,D.Singh,M.Sutaria,Composites:Part A30(1999)445.[15]H.Toraya,M.Yoshimura,S.Somiya,J.Am.Ceram.Soc.67(1984)119.[16]M.Singh,R.Asthana,Mater.Sci.Eng.A460–461(2007)153.[17]K.T.Fabert,A.G.Evans,Acta Metall.31(1983)565.[18]X.J.Zhou,G.J.Zhang,Y.G.Li,Y.M.Kan,P.L.Wang,Mater.Lett.61(2007)960.。
挤压速度对6082铝合金圆管显微组织和机械性能的影响
机械加工与制造M achining and manufacturing挤压速度对6082铝合金圆管显微组织和机械性能的影响向 耀1，2，梁豪辉1，2，周晶哲1，2，李一峰2，3，秦 简2，3，张 海2，3（１．广东澳美铝业有限公司，广东 佛山 ５２８１３７；２．苏大澳美轻金属研究院，广东 佛山 ５２８１３７；３．苏州大学高性能金属结构材料研究院，江苏 苏州 ２１５００６）摘 要：由于良好的机械性能和轻量化效果，铝合金挤压型材在运输行业中有着广泛的应用。
而提高挤压速率是降低生产成本提高产能的有效方式。
但在该过程难以避免产生粗晶。
本文主要介绍挤压速度对６０８２圆管微观组的影响，分析了变形条件对软化机制的影响，并对机械性能和进行了分析总结。
关键词：６０８２铝合金；挤压速率；机械性能中图分类号：ＴＧ１４６．２１ 文献标识码：Ａ 文章编号：１００２－５０６５（２０２０）１０－００２８－２Effect of extrusion speed on Microstructure and mechanicaI properties of 6082 aIuminum aIIoy tube XIANG Yao1,2, LIANG Hao-hui1,2, ZHOU Jing-zhe1,2, LI Yi-feng2,3, QIN Jian2,3, ZHANG Hai2,3（１．Ｇｕａｎｇｄｏｎｇ Ａｏｍｅｉ Ａｌｕｍｉｎｕｍ Ｃｏ．， Ｌｔｄ，Ｆｏｓｈａｎ ５２８１３７，Ｃｈｉｎａ；２．Ａｏｍｅｉ Ｉｎｓｔｉｔｕｔｅ ｏｆ ｌｉｇｈｔ ｍｅｔａｌｓ，Ｆｏｓｈａｎ ５２８１３７，Ｃｈｉｎａ；３．Ｒｅｓｅａｒｃｈ Ｉｎｓｔｉｔｕｔｅ ｏｆ ｈｉｇｈ ｐｅｒｆｏｒｍａｎｃｅ ｍｅｔａｌ ｓｔｒｕｃｔｕｒａｌ ｍａｔｅｒｉａｌｓ， Ｓｕｚｈｏｕ Ｕｎｉｖｅｒｓｉｔｙ，Ｓｕｚｈｏｕ ２１５００６，Ｃｈｉｎａ）Abstract: Ｂｅｃａｕｓｅ ｏｆ ｉｔｓ ｇｏｏｄ ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ ａｎｄ ｌｉｇｈｔｗｅｉｇｈｔ ｅｆｆｅｃｔ， ａｌｕｍｉｎｕｍ ａｌｌｏｙ ｅｘｔｒｕｄｅｄ ｐｒｏｆｉｌｅｓ ａｒｅ ｗｉｄｅｌｙ ｕｓｅｄ ｉｎ ｔｒａｎｓｐｏｒｔａｔｉｏｎ ｉｎｄｕｓｔｒｙ． Ｉｎｃｒｅａｓｉｎｇ ｅｘｔｒｕｓｉｏｎ ｒａｔｅ ｉｓ ａｎ ｅｆｆｅｃｔｉｖｅ ｗａｙ ｔｏ ｒｅｄｕｃｅ ｐｒｏｄｕｃｔｉｏｎ ｃｏｓｔ ａｎｄ ｉｎｃｒｅａｓｅ ｐｒｏｄｕｃｔｉｏｎ ｃａｐａｃｉｔｙ． Ｈｏｗｅｖｅｒ， ｉｔ ｉｓ ｄｉｆｆｉｃｕｌｔ ｔｏ ａｖｏｉｄ ｃｏａｒｓｅ ｇｒａｉｎｓ ｉｎ ｔｈｉｓ ｐｒｏｃｅｓｓ． Ｔｈｉｓ ｐａｐｅｒ ｍａｉｎｌｙ ｉｎｔｒｏｄｕｃｅｓ ｔｈｅ ｉｎｆｌｕｅｎｃｅ ｏｆ ｅｘｔｒｕｓｉｏｎ ｓｐｅｅｄ ｏｎ ｔｈｅ ｍｉｃｒｏｓｔｒｕｃｔｕｒｅ ｏｆ ６０８２ ｔｕｂｅ， ａｎａｌｙｚｅｓ ｔｈｅ ｉｎｆｌｕｅｎｃｅ ｏｆ ｄｅｆｏｒｍａｔｉｏｎ ｃｏｎｄｉｔｉｏｎｓ ｏｎ ｓｏｆｔｅｎｉｎｇ ｍｅｃｈａｎｉｓｍ， ａｎｄ ｓｕｍｍａｒｉｚｅｓ ｔｈｅ ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ．Keywords: ６０８２ ａｌｕｍｉｎｕｍ ａｌｌｏｙ； ｅｘｔｒｕｓｉｏｎ ｒａｔｅ； ｍｅｃｈａｎｉｃａｌ ｐｒｏｐｅｒｔｉｅｓ挤压是铝的三大主要的生产加工方式之一，挤压速度的快慢直接和生产效率直接相关，提高挤压速度是工程技术人员不懈的追求。
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Advanced Performance Materials4,95–103(1997)c 1997Kluwer Academic Publishers.Manufactured in The Netherlands. Microstructure Development and Mechanical Properties of Ni Matrix/Carbide CompositesJ.LIRA-OLIV ARES,A.R.DI GIAMPAOLO AND M.VELEZGrupo de Superficies e Interfaces,Depto Ciencia de los Materiales,Universidad Sim´o n Bol´ıvar,Apartado89,000 Caracas1080-A,VenezuelaI.C.GRIGORESCUDepto Tecnolog´ıa de Materiales,Intevep,S.A.,Apartado76,343,Caracas1070-A,VenezuelaAbstract.Nickel alloy matrix/dispersed carbide(VC,NbC,WC)composites were prepared by eitherflame-spray,liquid phase sintering,or solid state sintering.A commercial Ni-B-Si alloy was mixed with15%vol. of carbide particles and used to prepare composite coatings byflame-spray,bulk composite materials by solid state sintering(below1045◦C)or liquid phase sintering(above1050◦C).Phase characterization was perfor-med by X-ray diffraction,optical microscopy,scanning electron microscopy and X-rayfluorescence in energy and wavelength dispersive spectroscopy modes.Similar microstructural features were produced by thermal spray and liquid phase sintering:a Ni-rich matrix and a boron-rich intergranular phase.Sintered samples showed higher wear resistance than the coatings obtained byflame spraying.In both cases the wear mechanism is dominated by the plasticflow of the Ni-rich ductile matrix and the fracture of a boron-rich intergranular phase,the latter serving as a retainer.Carbide removal was observed for solid state sintered samples.Keywords:carbides,composites,metal matrix composites,niobium carbide,vanadium carbide,wearIntroductionApplications requiring high-wear resistance often make use of composite materials consist-ing offine hard ceramic particles dispersed in a transition metal matrix(Ni,Co,Cr),which serves as a tough binder.Ceramic materials include WC,TiC,B4C,TiB2,VC,TaC,NbC, and Al2O3,i.e.[1].The mechanical stability of these systems,which in turn determine tool performance,depends on the characteristics of the metal/carbide interface and the nature of the intergranular phases.Vanadium carbide is a less common hardening agent despite its high hardness and wear resistance[2].It shows good tribological performance when dispersed in an iron alloy matrix or when deposited as a continuous coating[3,4].The use of NbC has been limited to Ni matrices.The wetting behavior at the Ni/NbC interface is relatively poor compared to WC or TiC when employed as reinforcements[5].However,it has been observed that NbC can impart excellent wear characteristics to some Ni-Cr-B-Si alloys,when present as dispersed particles in proportions between20to50vol.%[6].In order to systematically develop composite materials,a combination of experimen-tal and theoretical studies has been undertaken.This work describes the methods used to develop Ni alloy/MC(M=V or Nb)composite materials for applications in the oil96LIRA-OLIV ARES ET AL. industry.Their behavior is compared with a conventional composite material system like Ni alloy/WC.These traditional composites based on NiBSi alloys(with eventual addition of Cr,Mo and/or Fe)are commonly used in the form of thick coatings deposited either by thermal spraying or welding.These coatings are quite successfully controlling abra-sion and erosion in the oil industry.Some typical applications are pistons for subsoil pumps coated by fused thermal spraying,welded hard-bands on drilling tubes and lin-ings for the restoration of steel components in contact with slurry in refinery equipment. The expectedfinal use for the composite systems developed in this work are as coat-ings for steel substrates and sintered preforms for fabrication of inserts in wear-resistant pieces.Materials and methodsThe precursors used in this study were commercial powders,as follows:—Vanadium carbide,99%purity,9µm mean particle size,V8C7as major phase(Good Fellow);—Niobium carbide,99%purity,5.5µm mean agglomerates size,formed by irregularly shaped particles,of size less than1µm(Cerac Corporated);—Tungsten carbide,99.5%purity,45µm mean agglomerates size and particle size of about1µm(Cerac Corporated);—Nickel alloy,spherical particles of a NiBSi based alloy(wt%composition:2.0Si,1.5B,0.2Fe,0.01Mn,0.03Cr,balance Ni),45µm mean size,consisting of Ni as the majorphase and Ni3B as a minor phase(Chromtec alloy from Eutectic+Castolin).Carbide powders were milled in a SPEX-800ball milling,with WC milling charge. During milling NbC and WC agglomerates were fractured to about1µm mean particle size,while VC particles size remained unchanged.A mixture of15%vol.of carbide and 85%metal matrix powders was prepared in an electric mortar,by blending for30min.in toluene as a lubricant.Thefirst process used in this study wasflame-spraying of the NiBSi alloy and the NiBSi/MC mixture with an oxyacetyleneflame torch.The coatings were deposited on SAE1045steel substrates.After deposition,the coatings were submitted to local fusion at about1050–1100◦C using theflame torch.The second process consisted of treating bulk samples of the NiBSi alloy and NiBSi/ carbides mixture under conditions of liquid phase or solid state sintering.The powder was uniaxially pressed at113MPa where green density of the pellets reached a constant value. The samples were heat-treated in an argon atmosphere between900and1100◦C.Sintered samples were metallographically prepared by diamond polishing and chemical etching in accordance with the ASTM E407standard in a50:50vol.nitric acid/acetic acid solution. The microstructural morphology was observed by SEM with EDS and WDS detectors to determine the element distributions.Vickers microhardness measurements were obtained with a Shimadzu tester,model4046 with a0.49N load applied during15s(ASTM E384).The reported values are averageMICROSTRUCTURE DEVELOPMENT97 results offive measurements.Density tests were carried out according to the ASTM-D2320 method and six measurements were taken for each sample.Friction and wear tests were carried out on samples metallographically polished and ultra-sonically cleaned in acetone.These tests were performed on the Falex LFW-1equipment in a block-on-ring configuration,using standard SAE4620steel rings(diameter of49.2mm) with a sliding speed200rpm,a normal load of22.4N and test durations of70min.Previous work[7]showed that the steady state was reached during this time.The friction force was continuously recorded during tests in order to determine the average friction coefficient in the steady state region.The wear volume was calculated from the measured values of the width of the wear scar,assuming a cylindrical shape[8].Results and discussionSeveral temperatures and sintering times were tested in order to determine the conditions for solid and liquid state sintering.The SEM analysis revealed that the liquid phase in the NiBSi alloy was produced between1045and1050◦C.Figure1shows the microstructure of the NiBSi alloy treated under an argon atmosphere and conditions of(a)solid state sintering at1000◦C during30min.and(b)liquid phase sintering at1100◦C during5min.Optimal parameters for liquid phase sintering of the nickel alloy were found to be1055◦C and5min. These conditions correspond to the maximum density(see Table1)without deformation.At higher temperatures or higher sintering times,considerable grain growth and macroscopic deformation of the samples was observed.The liquid phase sintered samples showed a Ni-B rich intergranular amorphous phase,as was indicated by WDS[9,10]and XRD analysis[10].This intergranular phase,which was hard and brittle(see Table1),was not formed below1045◦C;instead a hard intragranular precipitate phase,as in the precursor alloy powders was observed(figure1(a)).Figure1(c)shows the microstructure offlame-sprayed NiBSi samples.They showed coalescence of NiBSi grains and an intergranular phase like the liquid phase sintered NiBSi Table1.Average microhardness and relative density of the NiBSi alloy and NiBSi/MC composites sintered at various conditions.Sintering conditions Microhardness(HV50)RelativeTemperature Time Ni Intergranular MC density Material(◦C)(min.)grains phase grains(%)NiBSi alloy95030———87.910455318——90.1105552651030—92.4 NiBSi/15%vol.VC95030431—263481.710405———89.910505———92.410501203411234252793.3 NiBSi/15%vol.NbC105052931134179494.8NiBSi/15%vol.WC10505———92.998LIRA-OLIV ARES ET AL.Figure1.SEM micrographs of the NiBSi alloy powder:(a)sintered at1000◦C for30min.(solid state sintering), (b)sintered at1100◦C for5min.(liquid phase sintering),(c)deposited byflame-spray process.(Continued on next page)MICROSTRUCTURE DEVELOPMENT99Figure1.(Continued.)alloy.It was also observed by SEM that in the NiBSi/VC composite,the carbide particles were somewhat less uniformly distributed and with smaller dimensions in comparison with the sintered samples.This is probably due to the high oxidation susceptibility of this carbide at temperatures above400◦C[11].Figure2shows the microstructure of NiBSi/VC samples,heat treated under conditions of (a)solid state sintering and(b)liquid phase posite samples produced by solid state sintering showed the same microstructure observed in the nickel alloy,with the carbide particle located between the nickel grains(comparefigures1(a)and2(a)).The original grit carbide particles appear as rounded particle,suggesting that the vanadium carbide dissolves in the metal matrix.Infigure2(b),three different zones were observed by SEM.The matrix was formed by Ni-rich grains.The second zone is an intergranular network,formed by solution and reprecipitation of the original intragranular precipitates.As detected from the WDS analysis,the B concentration in this network was substantially higher than in the Ni-rich grains.The VC carbide particles were embedded in the intergranular Ni-B-rich network.Figure3shows the microstructure of NiBSi/WC composites obtained by solid state sintering;similar microstructure was observed for NiBSi/NbC pared with the microstructure of the NiBSi/VC system,these composites show a more homogeneous distribution of the carbide particles,due to the smaller particle size of the carbide precursors. Qualitative diffusion experiments carried out in previous work[10]in the NiBSi/NbC system showed that during heating,B migrates towards the NbC particles,while silicon is apparently uniformly distributed among the different phases.100LIRA-OLIV ARES ET AL.Figure2.A representative microstructure of NiBSi/15%vol.VC sintered at:(a)950◦C for30min.(solid state sintering),and(b)1100◦C for5min.(liquid phase sintering).The boron profile obtained by WDS line scan shows maximum concentration in the intergranular phase.MICROSTRUCTURE DEVELOPMENT101Figure3.SEM micrograph of NiBSi/15%vol.VC WC composite obtained by solid state sintering.Table1shows the average microhardness and density measurements of the NiBSi alloy and NiBSi/MC composites.The higher microhardness of the Ni-alloy grains was obtained for solid state sintering conditions(431HV).For liquid phase sintered samples the micro-hardness decreases,since the intragranular precipitates were segregated;the intergranular network observed in this case showed a microhardness of about1100HV.Standard devia-tions between5and9%were found for these measurements.The relative density increased with the sintering temperature and reached maximum values of about93%at1050◦C.How-ever this density is quite low for industrial application and other fabrication processes such as hot isostatic pressing are being explored.Table2shows the results of the measured friction coefficient and wear loss.Low variation was observed in the mean friction coefficient in the steady-state stage.Wear differences between sintered samples are also low;however some influence of the sintering temperature and the carbide concentration was found.The tipe of carbide also determine the wear volume:NiBSi/WC and NiBSi/VC composites showed similar values,although,a higher wear rate was observed for NiBSi/NbC.Flame-sprayed samples showed significant wear loss compared with sintered samples.This behavior could be explained by the fact that the B-rich network is more densely distributed in thermally coated samples,increasing the probability of exposing a larger fraction of this brittle phase to high stresses.As observed by SEM in previous work[7],upon catastrophic failure of the brittle intergranular network, the ductile Ni phase is free toflow under pressure,increasing the wear rate.The presence of VC particles increased wear rate in the case offlame sprayed samples.For solid state sintering,lower wear volume and carbide removal was observed.102LIRA-OLIV ARES ET AL. Table2.Influence of sintering conditions on wear behavior.Sintering AverageFabrication temperature friction Wear volume Material procedure(◦C)coefficient(mm3)NiBSi Alloy Thermal spray—0.680.22Liquid phase sinter.11000.760.12 NiBSi/15%vol.VC Thermal spray—0.660.35Solid state sinter.10400.940.09Liquid phase sinter.10500.850.11Liquid phase11000.700.14 NiBSi/30%vol.VC Liquid phase sinter.10500.870.08NiBSi/15%vol.NbC Liquid phase sinter.10500.800.16NiBSi/15%vol.WC Liquid Phase Sinter.10500.910.08Note:The sintering time was5min.for all samples.ConclusionsFlame spraying and liquid phase sintering were used to produce NiBSi/MC composites. These materials were characterized by the presence of a Ni-rich matrix with a brittle B-rich intergranular network.The wear and friction behavior is rather similar for all materials since it appears to be controlled by the fracture of the B-rich net and the plasticflow of the ductile Ni-rich matrix.For solid-state sintering,intergranular precipitates,similar to the precursor powders and microstructures,free of the intergranular deleterious phase were obtained.This microstruc-ture seems to decrease the wear behavior.However solid state sintered composites showed lower density,so the fabrication process must be improved in order to obtain denser material with low concentration of the intergranular phase.The stability of the NiBSi/VC and NiBSi/NbC composite materials was shown to depend on the stability of the amorphous intergranular net,more than on the adhesion of the carbides to the metal matrix,as was the case for NiBSi/WC[8].AcknowledgmentsThe authors would like to thank Fernandez,R.,Hernandez,A.,and Valecillos,F.for their helpful assistance in the experiments.Finantial aid from Intevep,S.A.and the USB “Decanato de Investigaciones”is appreciated.References1.Tiegs,T.N.and McDonald,R.R.:Ductile Ni3Al alloys as bonding agents for ceramic materials in cuttingtools.U.S.Patent5,015,290,May14,1991.MICROSTRUCTURE DEVELOPMENT103 2.Exner,H.E.:Physical and chemical nature of cemented carbides.International Metals Review,4:149–173,1979.3.Eyre,T.S.:Wear characteristics of metals.Tribology International,9:1–10,1976.4.Arai,T.:Carbide coating process by use of molten borax bath in Japan.Heat Treating,1:15–22,1979.5.Knotek,O.,Lohage,P.,and Reimann,H.:Nickel-based wear resistant coatings by vacuum melting.ThinSolid Films,108:449–458,1983.6.Knotek,O.,Lohage,P.,and Reimann,H.:Reactions between Ni-Cr-B-Si matrices and carbide additions incoatings during fusion treatment.Thin Solid Films,83:361–367,1981.7.Grigorescu,I.C.,Lira Olivares,J.,Di Giampaolo,A.R.,Lavelle,B.,Contreras,H.,and Ruiz,H.:Fric-tion coefficient variation and wear mechanism of Ni-B-Si alloy reinforced with VC.Proceedings of the6th International Congress on Tribology,Eurotrib93Budapest,3:225–230,1993.8.Lira-Olivares,J.and Grigorescu,I.C.:Friction and wear behavior of thermally sprayed nichrome-WC coat-ings.Surface and Coating Technology,33:183–190,1987.9.Grigorescu,I.C.,Di Rauso,C.,Drira-Halouani,R.,Lavelle,B.,Di Giampaolo,A.R.,and Lira,J.:Phasecharacterization in Ni alloy-hard carbide composites for fused coatings.Surface and Coating Technology, 76–77:494–498,1995.10.Di Giampaolo,A.R.,Lira-Olivares,J.,Valecillos,F.,and Velez,M.:Microstructural and chemical character-ization of nickel alloy-NbC composite.Surface Modification Technologies VII,T.S.Sudarshan,K.Ishizaki, M.Takata&Y.K.Kamata(eds.),The Institute of Materials,383–398,1994.11.Santafe,C.and Borgianni,C.:Study of the oxidation kinetics of vanadium carbide.Oxidation of Metals,9:415–425,1975.。