Phase composition of alumina–mullite–zirconia refractory materials

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齐鲁工业大学专业英语(材料)

第七单元
切割、研磨、抛光shearing grinding and polishing ，釉的热膨胀系数the coefficient of thermal expansion of the glaze ，熔块和着色氧化物frit and coloring oxide ，素烧和釉烧bisque -firing and glost -firing ，细模的磨料finely ground abrasive ，crack pit and flaw 裂纹凹陷缺陷，grinding glazing and decoration 研磨，上釉，装饰，to prevent warping or slumping during firing 避免烧结时变形和倒窑
第三单元
point defects点缺陷，chemical formulate 化学方程式，commence with 从。。开始，charged interstitial 带电间隙，host lattice 主晶格，no -mental 非金属，effectively neutral charge 有效中性电荷，vacancy pair 空位对，缺陷化学defect chemistry ，电子空穴electron hole ，亚原子粒子subatomic particles ，有效电荷effictive charge ，一价阴离子univalent anion ，二价阳离子divalent cation
9球磨粉末和单轴热压Ball -milled powders and hot uniaxial pressing
10陶瓷粉末固态化处理过程的传质动力The driving force for mass transport by solid -state process for ceramic powders

phase_diagram

• Phases:
The physically and chemically distinct material regions that result (e.g., a and b).
AluminumCopper Alloy
Adapted from Fig. 9.0,
Callister 3e.
3
• Changing T can change # of phases: path A to B. • Changing Co can change # of phases: path B to D.
e.g., if T = 100C, solubility limit = 80wt% sugar.
2
COMPONENTS AND PHASES
• Components:
The elements or compounds which are mixed initially
(e.g., Al and Cu)
--a composition (e.g., wt%Cu - wt%Ni), and --a temperature (T) How many phases do we get? What is the composition of each phase? How much of each phase do we get?
• 2wt%Sn < Co < 18.3wt%Sn • Result:
--a polycrystal with fine b crystals.
Pb-Sn system
Adapted from Fig. 9.10,
Callister 6e.
17
MICROSTRUCTURES IN EUTECTIC SYSTEMS-III

铝的晶格常数-体弹模量及弹性常数分子模拟

铝的晶格常数-体弹模量及弹性常数分子模拟CCCalculation of material lattice constant and bulk modulus 2012 Calculation of materiallatticeconstantand bulk modulusSummary ：Aluminum is one of the world's most used metals, the calculated aluminum lattice constant and bulk modulus can be used to improve the performance of the aluminum consequently make better use of aluminum. In virtue of molecular dynamics simulation software ，we can solve the lattice constant . By the derivative of the lattice con s tant, the bulk modulus can be obtained. The elastic constants of a material display the elasticity and we can use the software material studio to simulate and get them. The simulation results match the experimental values. Key words:Aluminum, lattice constant, bulk modulus, elastic constant , simulation.Introduction:In materials science, in order to facilitate analysis about the way in which the crystal particles are arranged, the basic unit can be removed from the crystal lattice as a representative (usually the smallest parallel hexahedron) as a compositionunit of dot matrix, called a cell (i.e. solid State Physics "original cell" concept); lattice constant (or so-called lattice constant) refers to the side length of the unit cell, in other words, the side length of each parallel hexahedral cells. Lattice constant is an important basic parameters of crystal structure. Figure A is the basic form of the lattice constant.Figure ALattice constant is a basic structural parameter，which has a direct relationship with the bondings between the atoms ，of the crystal substance. It reflects the changes in the internal composition of the crystal of the lattice constant, force state changes, etc.The bulk modulus (K or B) of a substance measures thesubstance's resistance touniform compression. It isdefined as the ratio of theinfinitesimal pressure increaseto the resulting relative decreaseof the volume. Its base unit is thepascal. Figure B describes theeffect of bulk modulus.Figure BThe bulk modulus canbe formally defined by theequation:where is pressure, isvolume, and denotes thederivative of pressure withrespect to volume.Our research object isaluminum，whose atomicnumber is 13 and relative massis 27.The reserves of aluminumranks only second to ferrumcompared with other metallicelements. Aluminum andaluminum alloy are consideredthe most economic andapplicable in many applicationfields as a consequence of theirexcellent properties. What’smore，increased usage ofaluminum will result fromdesigners' increased familiaritywith the metal and solution tomanufacturing problems thatlimit some applications.The crystal structure of aluminum is face-centered cubic. The experimental value of lattice constant and bulk modulus are 0.0.40491nm and 79.2Gpa. Computing theory and methods：Our simulation is on the basis that aluminum is offace-centered cubic crystal structure. We can get the exact value of lattice constant in virtue of molecular dynamics simulation software. Then by the derivation of lattice constant for energy E，we obtain bulk modulus.To start with ，compile a script for the use of operation and simulation in lammps. We set periodicboundary conditions in the script and create an analog box，whose x ,y ,z coordinate values are all confined to [0,3].Run the script in lammps, calculating the potential energy，kinetic energy as well as the nearest neighbor atoms for each atom. Finally put out the potential energy function of aluminum.Extract the datas under the linux system produced by lammps to continue the computation by means of matlab，from which we can get the lattice constant through several times of matching.Figure CFigure C——the curve shows the relationship between cohesive energy and lattice constant，which is what we get in the process of computing in matlab，points out the lattice constant corresponding with the least cohesive energy. The horizontal ordinate of the rock bottom stand for the lattice constant of aluminum which can be clearly located as 0.40500nm.Since we have obtained the lattice constant，we simulated the visualization of aluminu’s crystal structure. Figure D is what we get through the visualization.Figure DThe bulk modulus is defined as：As for cubic cell，the formula can be transformed into the following pattern：The bulk modulus can be calculated with the formula above combined with the latticeconstant. Finally，the bulk modulus is 78.1Gpa.Besides，to enrich our research, we have calculated the elastic constants of aluminum. Because of the symmetry of face-centered cubic，aluminum only has three elastic constants. To reach the target ，we have to establish a cell of aluminum in the software material studio in the first place and then transform it into a primitive cell. Figure E and figure F are the cell and the primitive cell we have established during the simulation .Figure EFigure FIt comes to the CASTEP step after the geometry optimization. We managed to makeE cut=350eV andK points=16*16*16 ，both of which are extremely importantand also sensitive to the calculation of elastic constant，after several trials. Figure G is the primitive cell that has been through the geometry optimization.Figure GThe calculated results are as follows:C11=106.2GpaC12=60.5GpaC44=28.7Gpa Correspondingly ，the experimental range of the three elastic constants are listed below：C11=108~112GpaC12=61.3~66GpaC44=27.9~28.5GpaAs we can see，although a little outside of the value range of the experimental standards, our results are right within the error range.Conclusion：Our group has calculated the lattice constant and bulk modulus of aluminum，both of which coincide with the experimental value，by means of lammps and matlab. Moreover ，we have found out that bulk modulus has a close relationshipwith temperature. As lattice constant haven’t made any change under small change of temperature while the energy of the material have changed，so we concluded that temperature change can influence bulk modulus as a consequence of the change of cohesive energy change resulting from temperature change.There are still problems in our research as you can see that the three elastic constants are a little out of the value range. But this group is the one closest to the experimental value. Our group have concluded that the errors result from the script which can affect the accuracy of the simulation.Besides，the most valuable thing we have learned is that we must seek the solutions and never give up in face of difficulties.References：[1]材料科学基础（胡赓祥、蔡珣、戎咏华上海交通大学出版社）[2]Ayton, Gary; Smondyrev, Alexander M; Bardenhagen, Scott G; McMurtry, Patrick; Voth, Gregory A.“Calculating the bulk modulus for a lipid bilayer with none quilibrium molecular dynamics simulation". Biophysical Society. 2002.[3]Cohen, Marvin (1985).CCCalculation of material lattice constant and bulk modulus 2012 "Calculation of bulk modulus ofdiamond and zinc-blende solids". Phys. Rev. B 32: 7988–7991. [4] Watson, I G; Lee, P D; Dashwood, R J; Young, P . Simulation of the Mechanical Properties of an Aluminum Matrix Composite using X-ray Microtomography: Physical Metallurgical and Materials Science. Springer Science & Business Media. 2006. [6] Ashley, Steven. Aluminum vehicle breaks new ground. Engineering--Mechanical Engineering. Feb 1994. [7] Sanders, Robert E, Jr; Farnsworth, David M. Trends in Aluminum Materials Usage forElectronics. Metallurgy. Oct2011.附文：lammps脚本units metalboundary p p patom_style atomicvariable i loop 20variable x equal 4.0+0.01*$ilattice fcc $xregion box block 0 3 0 3 0 3 create_box 1 boxcreate_atoms 1 boxpair_style adppair_coeff * * AlCu.adp Almass 1 27neighbor 1 binneigh_modify every 1 delay 5 check yes variable p equal pe/108variable r equal 108/($x*3)^3 timestep 0.005thermo 10compute 3 all pe/atomcompute 4 all ke/atomcompute 5 all coord/atom 3dump 1 all custom 100 dump.atom id xs ys zs c_3 c_4 c_5dump_modify 1format"%d %16.9g %16.9g %16.9g%16.9g %16.9g %g"min_style sdminimize 1.0e-12 1.0e-12 1000 1000print "@@@@ (latticeparameter,rho,energy per atom):$x$r $p"clearnext ijump in.al至于代码的含义，可参考其它资料或查询lammps官网manual。

多重乳液法制备毫米级氧化铝空心球

Key words: alumina hollow spheres; oil1-in-water-in-oil2 emulsion; composite droplets; flow focusing micro-channel The suitability of target pellets requires high quality hollow micro-spheres with high sphericity and diameters that reach to millimeter-scale in the ICF energy technology[1]. The alumina hollow spheres have great potential application prospects using as targets because of the excellent physical and chemical performance[2]. In the current study, the ceramic hollow spheres have been prepared in various materials with wide sizes[3-8]. However, the size of prepared hollow spheres are mostly in nanometer and micrometer scale, the related literature reports are infrequent which the size of spheres reach to millimeter-scale. The preparation methods are mainly divided into chemical and physical methods[9-11]. The chemical reaction hardly reach to millimeter-scale[12], the advantage of physical template method is that the sizes of hollow spheres are easily controlled by changing sizes of template and reach to millimeter-scale[2]. Inspired by the micro-encapsulation technology, alumina hollow spheres with wall thickness of 30–100 µm and diameter of 600–2500 µm were prepared by Wang, et al using oil-in-water emulsion droplets as precursor by the method of soft template[13]. Liu, et al prepared the submicron Cu2O hollow spheres composed of small Cu2O nanoparticles of 22 nm in size via using a multiple emulsion (O/W/O) system as the template[14]. Yodthong, et al prepared hollow chitosan microspheres with 100 µm diameter by an oil1-in-water-in-oil2 emulsion solvent diffusion method though mixing and stirring[11]. Consequently, using double emulsion micro-encapsulation technology to synthesize ceramic hollow spheres is feasible. The emulsion micro-encapsulation technology is based on the complex system multiple emulsion termed “emulsion of emulsion” which have received a great deal of attention due to their improved stability and facilitated control of their properties and play critical roles in various fields including foods, cosmetics, pharmaceuticals and chemicals[15-16]. Compared with the simple emulsions consisting of only one phase, each dispersed globule in the double emulsions forms composite droplets with core-shell structure made of two immiscible liquids, and multiple liquid phase compartments separated from the water phase by a layer of oil phase compartment[17]. Based on the combination way of two phase fluid, the major types of composite droplets formed by multiple emulsions are the water-oil-water (W/O/W) and oil-water-oil (O/W/O) double emulsions consisting of oil (water) droplets dispersed within larger water (oil) droplets, which are them-

相变界面稳定性概念

Kinetics of isothermal phase transformations above and belowthe peritectic temperature:Phase-ﬁeld simulationsG.Boussinot *,E.A.Brener,D.E.TemkinInstitut fu ¨r Festko ¨rperforschung,Forschungszentrum Ju ¨lich,D-52425Ju ¨lich,Germany Received 5May 2009;received in revised form 9October 2009;accepted 13November 2009AbstractWe present phase-ﬁeld simulations of isothermal phase transformations in the peritectic system below and above the peritectic tem-perature.The physical processes involved are of diﬀerent natures,involving either a triple junction or a liquid-ﬁlm-migration (LFM)mechanism.Below the peritectic temperature,one of the solid phases steadily grows along the other.Above the peritectic temperature the phase transformation proceeds via the LFM mechanism.To the best of our knowledge,this mechanism has not been discussed in the literature as a generic process of phase transitions in peritectic systems.In addition to the LFM process,we also simulate melting along the solid–solid interface.Finally,we make a simpliﬁed linear stability analysis of the liquid ﬁlm,supporting our simulation results.Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Peritectic phase diagram;Phase-ﬁeld model;Triple junction;Liquid-ﬁlm-migration1.IntroductionIndustrial alloys,such as steels,often exhibit peritectic phase diagrams.At the peritectic temperature T P ,two solid phases and the liquid phase are in a three-phase equilib-rium.In comparison with eutectic alloys (in which an equi-librium also exists between three phases),phase transitions in peritectic systems have been much less investigated.A typical experiment concerning alloys involves direc-tional solidiﬁcation,in which the sample is pulled in a tem-perature gradient.The properties of the solidiﬁed alloy strongly depend on the microstructure left behind the solid-iﬁcation front.In peritectic systems,the directional solidi-ﬁcation exhibits a large variety of microstructures.For example,one can observe coupled growth (simultaneous growth of alternative lamellae of the two solid phases par-allel to the growth direction),discrete bands (alternative layers of the diﬀerent solid phases perpendicular to the growth direction),island formation (one solid phaseembedded in a matrix of the other solid phase)or oscilla-tory regimes.A review of these solidiﬁcation microstruc-tures can be found in Refs.[1,2].The variety of diﬀerent scenarios is due to the peritectic phase diagram,in which the two solid phases play a diﬀer-ent role.The peritectic temperature T P corresponds to an equilibrium between the peritectic phase ðc Þ,the primary phase ðd Þand the liquid phase ðL Þ—see Fig.1.Below the peritectic temperature ðT 0<T P Þ,and suﬃ-ciently above c c ,the system is in the two-phase region ðc þL Þof the phase diagram (see Fig.1),and the solid–liquid ðc þL Þequilibrium mixture is stable (with the liquid concentration c L c ðT 0Þand the solid concentration c c L ðT 0Þ).On the other hand,the system is also in the ðd þL Þtwo-phase region,but the ðd þL Þequilibrium mixture is meta-stable (with the liquid concentration c L d ðT 0Þand the solid concentration c d L ðT 0Þ).Above the peritectic temperature ðT 0>T P Þthe ðc þL Þmixture is metastable,while the ðd þL Þmixture is stable.Hence,depending on the temper-ature T 0,the phase transformation from a metastable equi-librium to a stable equilibrium takes the form:ðd þL Þ!ðc þL Þfor T 0<T P and ðc þL Þ!ðd þL Þfor1359-6454/$36.00Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.actamat.2009.11.017*Corresponding author.E-mail address:guilaume.boussinot@ (G.Boussinot)./locate/actamatAvailable online at Acta Materialia 58(2010)1750–1760T0>T P.The magnitude of the driving force for these phase transformations can be represented by the dimen-sionless supersaturation D¼ðc L dÀc L cÞ=ðc cÀc dÞ.We note that D<0ðD>0Þfor T0<T PðT0>T PÞ.Because diﬀusion in solids is very slow,one can observe aðdþcÞmixture for long time when the sample is below T P,even if this is outside the solid–solid two-phase region.The large variety of directional solidiﬁcation microstruc-tures in peritectic systems makes a general theory quite a challenging task.Indeed,since the system evolves with the temperature gradient,solidiﬁcation can occur above or below the peritectic temperature.The existence of the temperature gradient excludes the presence of any solid phase far ahead of the solidiﬁcation front,where the tem-perature is high,and the presence of any liquid phase far behind the front,where the temperature is low.Because of the change in the stability of the two solid–liquid equilib-ria when one crosses T P,the directional solidiﬁcation pro-cess is of a complex nature.An isothermal process,which is more simple,can be understood as the limit of vanishing gradient in a directional process.In this paper we investigate isothermal transformations below and above the peritectic temperature.The composi-tion of the sample is varied throughout the wholeðdþLÞtwo-phase region(hypo-and hyperperitectic compositions) for temperatures below T P,and is close to the peritectic composition c c for temperatures above T P.Our purpose is toﬁnd generic steady-state moving patterns.For this purpose we use the phase-ﬁeld numerical method.The code used is the one developed in Ref.[3]for eutectic directional growth.Thus we had to adapt it to the isothermal process. The conversion to the peritectic system is straightforward since it only consists of changing the phase diagram.The phase-ﬁeld method has proven its eﬃciency in reproducing the solidiﬁcation microstructures of peritectic systems.In particular,it qualitatively reproduces the vari-ety of microstructures during directional solidiﬁcation in peritectic systems[4,5].The phase-ﬁeld method is based on continuousﬁelds which locally represent the physical state of the system.They obey some evolution equations, and these equations are solved at each point of the compu-tational grid.When describing a classical solidiﬁcation process of a pure material,one uses a temperatureﬁeld and a phaseﬁeld.The value of the phaseﬁeld indicates whether the system is solid or liquid.At the solid–liquid interface,the phaseﬁeld varies smoothly(on aﬁnite thick-ness)from the value for the liquid to the value for the solid. Taking the limit of vanishing interface thickness,one must recover the sharp interface description of the problem.The sharp interface description essentially consists of solving the diﬀusion equation inside domains with boundary con-ditions at the interfaces.This is the real physical approach when one considers a separation of length scales between interface thickness(usually of atomic scale)and micro-structure.As in eutectic systems,the transformations in peritectic systems involve two solid phases(c and d)in addition to the liquid ually,more than one phase ﬁeld is deﬁned in order to distinguish more than two phases.The triple junction is the point where the three phases meet.In the sharp interface approach,one usually assumes thermodynamic equilibrium at the triple junction,leading to Young’s law.The latter states that the sum of the forces due to surface tensions acting on the triple junction van-ishes.Once surface tensions are known,the orientations of two interfaces with respect to the third one are deter-mined.Unlike the sharp interface approach,the triple junc-tion is not a special point in the phase-ﬁeld model since one solves the same evolution equations of theﬁelds at each grid point.However,in the limit,where the thickness of the interfaces vanishes,the phase-ﬁeld model description of the triple junction should satisfy Young’s law,which is a non-trivial task[3].The possible process in peritectic systems which involves the motion of the triple junction is schematically represented in Fig.2a in the spirit of Ref.[6].Another process can take place without the triple junction(Fig.2b).Liquid-ﬁlm-migration(LFM)is where a solidiﬁcation front and a melting front move together separated by a liquid layer.The kinetics are then controlled by the diﬀusion in the liquidﬁlm.In the following we present our phase-ﬁeld simulations of isothermal transformations in a model peritectic system above and below the peritectic temperature.In Section2, we present brieﬂy the phase-ﬁeld model(for more details, see Ref.[3]).In Section3,we present our simulation results in the case of T0<T P.The process,which involves a triple junction,occurs via a steady-state motion where the stable c phase grows along the metastable d phase.ThistransitionG.Boussinot et al./Acta Materialia58(2010)1750–17601751is known as the“peritectic reaction”.The conﬁguration of the interfaces near the triple junction was suggested by Kurz and Fischer[6]and numerically reproduced by Nes-tler and Wheeler[7].We investigate the steady-state veloc-ity dependence of peritectic reaction on average concentration and temperature.We would also like to mention the work of Das et al.[8,9]on the peritectic trans-formation d!c occurring through solid diﬀusion as the next stage of the microstructure evolution after the peritec-tic reaction.In Section4,we present our numerical results above the peritectic temperatureðT0>T PÞ.The phase transforma-tion proceeds via a LFM mechanism.To the best of our knowledge,this mechanism has not been discussed in the literature as a generic process of phase transitions in peri-tectic systems.We investigate the temperature dependence of the steady-state velocity of the LFM.Moreover,in the case of T0>T P,we numerically investigate the melting of peritectic alloys along a c=d interface,and study the tem-perature dependence of the steady-state velocity.In Section 5,we brieﬂy discuss the stability of the LFM process using a simpliﬁed linear analysis,and then give an explanation for the qualitative diﬀerence of physical processes involved in the cases of T0<T P and T0>T P.Finally,we make a brief comment on directional solidiﬁcation.2.ModelThe free-boundary problem corresponding to the phase transformation in the peritectic system consists of the determination of the pattern produced by the diﬀerent phases(i.e.the location of the interfaces between phases) and the concentrationﬁeld.The concentrationﬁeld c obeys the diﬀusion equation in the liquid with diﬀusion coeﬃcient D:@c@t¼D$2c;ð1Þand we neglect diﬀusion in the solid phases.This may not be appropriate for alloys with high T P.However,we have veriﬁed that the introduction of a solid bulk diﬀusion coef-ﬁcient D s%D=10does not alter qualitatively the results of the present paper.Furthermore,since we investigate gener-ic patterns for small deviations from peritectic equilibrium, we used a temperature-independent diﬀusion coeﬃcient D.The mass balance equation at an interface between phase i and phase j reads:V nðc iÀc jÞ¼ÀD i@c=@n ji þD j@c=@n jj;ð2Þwhere V n is the normal velocity,and n is the vector normal to the interface.c iðjÞis the concentration of phase i(j)at the interface,D iðjÞis the diﬀusion coeﬃcient in phase i(j),andthe normal gradient@c=@n jiðjÞis evaluated on the i(j)sideof the interface.If one neglects diﬀusion in cðdÞsolid phaseðD s=D!0Þ, Eq.(2)reads:V nðc L cðdÞÀc cðdÞLÞ¼ÀD@c=@n jL ;ð3Þat a cðdÞ=L solidiﬁcation front where theﬂux of materialfrom the solid into the liquid can be neglected.For acðdÞ=L melting front,the situation is diﬀerent.The concen-tration of the cðdÞsolid phase at the interface c cðdÞL is diﬀer-ent from the bulk concentration in the cðdÞmetastablephase c1.In the limit D s=D!0(i.e.the diﬀusion lengthin the solid is smaller than the other length scales of theproblem),theﬂux of material from the cðdÞphase intothe liquid phase becomes independent of D s(see e.g.Ref.[10]for more details)and cannot be neglected:ÀD s@c=@n jcðdÞ¼V nðc cðdÞLÀc1Þ:For the cðdÞ=L melting front,Eq.(2)thus reads:V nðc L cðdÞÀc1Þ¼ÀD@c=@n jL:ð4ÞThe liquid concentration at the solid–liquid interface isc¼c L cðdÞÀðc cÀc dÞd cðdÞj:ð5ÞThis corresponds to the local equilibrium at a given tem-perature and takes into account the Gibbs–Thomson cor-rection due to the local curvature j of the interface,which is assumed to be positive for a convex solid–liquidinterface.The capillary lengths are given by:d cðdÞ¼r cðdÞðc PÀc cðdÞÞðc cÀc dÞf00Lðc PÞ;ð6Þwhere r cðdÞis the surface energy of the L=cðdÞinterface,f LðcÞis the free energy density of the liquid phase,andf00LðcÞis its second derivative with respect to c.We use the phase-ﬁeld method to solve this problem.The model was developed in Ref.[3]to study eutectic direc-tional growth.Two diﬀerent solid phases and one liquidphase should be described.Three phaseﬁelds are thendeﬁned:piðrÞ;i¼c;d;L.The value of p iðrÞis1whenthe system is in phase i at the point r,and0otherwise.Thenpiis diﬀerent from1and0only within a transition regionof width W,i.e.at the interfaces between diﬀerent phases.The interface thickness W is arbitrary,and should not belinked to any physical length scale within the spirit of thephase-ﬁeld method.Here,the surface tensions are assumedisotropic for the sake of simplicity.We have an additionalconstraint:pcðrÞþp dðrÞþp LðrÞ¼1;ð7Þwhich allows the evolution equations to be solved only fortwo phaseﬁelds.In addition to these two phaseﬁelds,onealso introduces a concentrationﬁeld.We refer to Eqs.(4.1)and(4.2)in Ref.[3]for the time evolution of the phaseﬁelds and the concentrationﬁeld.In the diﬀusion equation,the diﬀusion coeﬃcient is chosen proportional to the liquidphaseﬁeld pLðrÞ.This allows Eq.(1)to be recovered in thelimit of vanishing interface width W,but is not the onlychoice.The theoretical two-phase equilibrium proﬁle for thephaseﬁelds is the well-known hyperbolic tangent with acharacteristic length scale W.This theoretical proﬁle isreached numerically by increasing W.If W is too small,1752G.Boussinot et al./Acta Materialia58(2010)1750–1760the small number of grid points within an interface leads to numerical problems.Since W has also to be much smaller than the characteristic length scales of the microstructure, one has toﬁnd a compromise in order to avoid a prohibi-tive CPU calculation time.In the model for directional solidiﬁcation,the calcula-tion proceeds through a regular shifting procedure of the simulation box in order to reproduce the pulling of the sample with constant velocity in the temperature gradient. The simulation box then represents the system in the frame in which the temperature is constant in time.In our case, the velocity is not imposed,and we should calculate it. Thus,we also use a shifting procedure,but one that is not regular.We locate a point of interest(e.g.the tip of a solidiﬁcation front),and when this point has advanced by a certain small amount of grid points n,we shift the simu-lation box.This point of interest then stays approximately at the same position in the simulation box during the calcu-lation.If n0is the number of time steps needed to go through the n grid points,we calculate the velocity by V¼ðn D xÞ=ðn0D tÞ,where D x is the grid step and D t is the time step.When a steady state is reached,the interfaces and the concentrationﬁeld are time independent in the sim-ulation box(in the frame moving with the steady-state velocity).They then move with the steady-state velocity in the laboratory frame.At the lateral sides of the simula-tion box,we impose zero normal derivatives for the phase ﬁelds and the concentrationﬁeld.The physical parameters used in our calculations are the ones given in Table1of Ref.[11].They were used to model an Fe–Ni alloy.In Fig.3,we present the corresponding phase diagram.The horizontal axis is the dimensionless concentration~c¼ðcÀc PÞ=ðc cÀc dÞ,and the vertical axis represents the dimensionless supersaturation D¼ðc L dÀc L cÞ=ðc cÀc dÞ.The dimensionless equilibrium concentra-tions at the peritectic temperature are~c c¼À1:16and ~c d¼À2:16(and0for the liquid phase).The liquidus and solidus lines are considered parallel for each two-phase region.The liquidus lines are given by~c L cðDÞ¼À1:19Dand~c L dðDÞ¼À0:19D.In Fig.3theðcþdÞtwo-phase region is represented by vertical lines.A realistic descrip-tion is not relevant,since no diﬀusion is allowed in the bulk solid phases and theðcþdÞequilibrium does not enter the physical model described by Eqs.(1),(3)and(4).In the phase-ﬁeld model,the capillary lengths given by Eq.(6)are not chosen independently(see Ref.[3]for more details).However,the average d0¼ðd cþd dÞ=2is chosen. Hence,we calculated the average of the capillary lengths given in Ref.[11]and took it as an input for our simula-tions.The interface energies determine the contact angles at the triple junction through Young’s law.In the phase-ﬁeld model used here,these energies are considered equal for all interfaces.Thus,the interfaces at a triple junction divide the circle into three equal angles.Finally,the diﬀu-sion coeﬃcient D gives the time scale of the physical pro-cess,and again,the value is taken from Ref.[11].To verify that the phase-ﬁeld model is appropriate to study peritectic systems,we reproduced directional solidiﬁ-cation processes for which some results of coupled growth obtained from a boundary-integral-method are given in Refs.[11,12].In Fig.4we present the steady-state patterns of coupled growth obtained from phase-ﬁeld simulations which correspond to one period of the lamellar structure. The lamellae spacing(horizontal width of the presentedﬁg-ure)corresponds to k¼205d0for Fig.4a and k¼102:5d0 for Fig.4b.In both cases,the dimensionless concentration in the liquid far ahead of the solidiﬁcation front is ~c¼À1:52.The thermal lengths,as deﬁned in Ref.[11], areðl Tc¼224d0;l Td¼1391:5d0Þfor Fig.4a,andðl Tc¼37:5d0;l Td¼232d0Þfor Fig.4b.The interface width is in both cases,and for the rest of the present paper, W¼1:6d0.For Fig.4a the physical parameters are the same as in Ref.[11](Fig.9with G¼3K mmÀ1).For Fig.4b the physical parameters are the same as in Ref.[12](Fig.4.6).Weﬁnd good agreement in the steady-state pattern obtained by the phase-ﬁeld method and by the boundary-integral method.Moreover,this kind of micro-structure has been often reported in experiments in the Fe–Ni system[13],in the Ni–Al system[14],and more recently in the Cu–Sn system[15].However,a slight qualitative diﬀerence is present between our phase-ﬁeld simulations and the boundary-integral results.In Fig.4a and b one can observe,at the tri-ple junction,the angle adopted by the c=d interface with respect to the direction of velocity.In comparison,in Fig.9in Ref.[11]and Fig.4.6in Ref.[12],the solid–solid interface remains straight(aligned with the velocity direc-tion)until the triple junction.In our phase-ﬁeld calcula-tions the angle is a feature of the diﬀuseness of the interfaces.On the c=d side of the triple junction the liquidphaseﬁeld pLdecays smoothly,and so does the bulkdiﬀusion coeﬃcient(which is proportional to pL).Some diﬀusion then occurs at the c=d interface in the vicinity of the triplejunction.G.Boussinot et al./Acta Materialia58(2010)1750–176017533.Transformation at T 0<T PIn the case T 0<T P the steady state obtained in our sim-ulations corresponds to the so-called peritectic reaction [16],where the stable c phase solidiﬁes along the metastable d phase.This transformation involves a triple junction.A schematic illustration of the triple junction during this transformation is given in Ref.[6].The d phase melts,with an impurity ﬂux towards the melting front,and the c phase solidiﬁes,rejecting impurities.Far ahead of the triple junc-tion ðþ1Þ,the system consists of a ðd þL Þequilibrium,imposed as a boundary condition in our simulations.The d phase fraction is related to the average concentration by the lever rule.The most advanced point of c phase is used as the reference point for our shifting procedure.The width of the simulations is equal to k ¼102:5d 0,and the interface width is still W ¼1:6d 0.We have used several initial conﬁgurations to investigate the robustness of the obtained steady state,and one of these is displayed in Fig.5.It consists of a ðd þL Þequilibrium on the right side of the box (according to the lever rule)and a ðc þL Þequi-librium on the left side of the box.In between a liquid ﬁlm of thickness k =4is left.The concentration in the liquid phase is c L c on the left side of the box (light grey)and c L d everywhere else (dark grey).Whatever the initial thicknessof the liquid ﬁlm or phase arrangement on the left side of the box,the evolution of the simulation proceeds through a rapid shrinking of the liquid layer,followed by the growth of the c phase along the d phase.In Fig.6we show the steady-state pattern obtained for diﬀerent average concentrations.The supersaturation is D ¼À0:07.The c phase (shown in white)grows to the right along the melting d phase (shown in black).The product of the transformation is a ðc þd þL Þmixture,except for the case of ~c ¼À1:9where the c front touches the channel wall and a ðc þd Þmixture is produced.In Fig.7we plot the steady-state dimensionless velocity Vd 0=D with respect to the dimensionless average concentration ~c .For small j ~c j the velocity shows weak variation,and decreases when j ~c j increases.For j ~c j ¼1:9where the c front touches the chan-nel wall,the velocity is signiﬁcantly higher.In Fig.6we see that when the liquid phase fraction is large ð~c ¼À0:5;À0:7;À0:9;À1:2Þ,the c solidiﬁcation front exhibits almost the same parabolic shape for any ~c .This shows that the inﬂuence of the channel boundary is negli-gible,and the c =L front can be interpreted as an Ivantsovparabola [17].We have checked that the observed PE´clet numbers qualitatively obey the Ivantsov relation.The dif-fusion ﬁeld in the vicinity of the triple junction is almost independent of the average concentration,and the steady-state velocity is then weakly dependent on ~c (see Fig.7).When j ~c j increases,the diﬀusion ﬁeld begins to be inﬂu-enced by the channel boundary and the c =L front becomes ﬂat behind the triple junction.The steady-state velocity then decreases.For the case ~c ¼À1:9,the velocity of the transformation ðd þL Þ!ðc þd Þis signiﬁcantly higher than the one obtained for the transformation ðd þL Þ!ðc þd þL Þ.This suggests the presence of this solution in a certain range of average concentration,which could overlap with the one of the ðd þL Þ!ðc þd þL Þtransformations.Note that the comparison between the patterns obtained from phase-ﬁeld simulations in Fig.6and theschematicFig.4.Pattern of the steady-state coupled growth during directional solidiﬁcation simulated with the phase-ﬁeld model.These simulations should be compared with Refs.[11,12](seetext).Fig.5.Initial conﬁguration of the simulation for ~c ¼À0:7.The left (right)side of the box contains a c ðd Þphase fraction according to the lever rule.The concentration in the liquid (L )is c L c ðc L d Þon the left (right)side of the box.1754G.Boussinot et al./Acta Materialia 58(2010)1750–1760pattern in Fig.2a in the vicinity of the triple junction is irrelevant since the physical length scales are of the order of the interface width of the phase-ﬁeld model.However,we veriﬁed that the phase-ﬁeld simulations reproduce the typical pattern depicted in Fig.2a when the interface width is small compared to the length scales at the triple junction.Note also that,when we assume a small diﬀusion coeﬃ-cient in the solid phases,the c =d solid–solid interface becomes slightly inclined,which suggests that no d phaseexists at À1according to the peritectic transformation mentioned in Introduction and described in Refs.[8,9].3.1.Dependence on supersaturationWe have studied the dependence of the steady-state velocity on the supersaturation j D j (Fig.8).The calcula-tions were performed with a large liquid phase fraction in order to consider the case were the channel width has almost no inﬂuence on the velocity.This eliminates the channel width as an additional length scale in theproblem.Fig.6.Simulations at D ¼À0:07ðT 0<T P Þfor diﬀerent average dimensionless concentrations ~c .The c phase (white)grows to the right along the d phase (black).The liquid is represented in grey.Left row from top to bottom:~c ¼À0:5;À0:7;À0:9;À1:2;right row from top to bottom:~c ¼À1:5;À1:7;À1:8;À1:9.G.Boussinot et al./Acta Materialia 58(2010)1750–17601755One observes that the velocity increases when j D j increases, which is expected.The curve approximately follows a j D j4 behavior.However,obtaining a scaling law from our phase-ﬁeld simulations is a challenging task.Indeed,the range of investigated supersaturation is relatively small, and smaller supersaturations are unreachable due to the large characteristic length scales.4.Transformation at T0>T PWhen T0>T P the thermodynamic equilibrium consistsof aðdþLÞmixture.The metastable equilibrium consists of aðcþLÞmixture or a single c phase if the operating point lies outside of theðcþLÞtwo-phase region.The ini-tial conﬁguration of the simulations is made of the metasta-ble equilibrium on the right side of the boxðþ1Þ,of the stable equilibrium on the left sideðÀ1Þ,and of a liquid ﬁlm in between.In Fig.9,we present the initial conﬁgura-tion for D¼0:14and~c¼À1:15which lies in theðcþLÞtwo-phase region.The concentration in the liquid is set to c L c on the right side of the box(dark grey),to c L d every-where else(light grey).The channel width is equal to k¼205d0,with W¼1:6d0and a liquid layer with a width k=8.In Fig.10,we present the steady-state pattern corre-sponding to the initial conﬁguration shown in Fig.9.The transformation consists inðcþLÞ!ðdþLÞand proceeds via an LFM process.In this case,the liquid layer between the two solid phases survives from the initial conﬁguration until the steady state is reached.The LFM process is a physical mechanism often reported in experimental works on sintering[18]or partial melting[19].A schematic illustration is given in Fig.2b. LFM is related to the problem of grain-boundary migra-tion,where a thin intergranular liquidﬁlm can propagate during grain evolution[20].In Ref.[21]a description of LFM is made within the assumption of vanishing surface tension.A velocity selection theory was then developed [22],introducing the eﬀect of surface tension(especially its anisotropy).Recently,a stability analysis of the liquid ﬁlm in the presence of stress has been carried out[23]. The LFM process occurs via the co-operative motion of two fronts,one solidiﬁcation front and one melting front, with a liquidﬁlm in between.The liquid concentration should be diﬀerent at the two fronts(diﬀerent equilibrium conditions),leading to a concentration gradient within the liquid necessary for a steady-state motion.Here,in the peritectic system(Fig.10),the d phase solidiﬁes and the c phase melts,and the concentration gradient in the liquid is the consequence of the diﬀerence between c L d and c L c.We also obtained an LFM process when the average concentration is on the solidus line of theðcþLÞtwo-phase region.In Fig.11the corresponding steady-state pattern for D¼0:14is presented.The channel width is k¼102:5d0.4.1.Dependence on supersaturationIn order to save CPU time,we investigated the depen-dence of steady-state velocity on supersaturation for a metastable equilibrium comprising a single c phase,which allows smaller channel widths to be investigated.The variation of the dimensionless velocity Vd0=D with respect to D is shown in Fig.12.Within the investigated range of D the curve exhibits a weak non-linear behavior. In addition,we investigated the variation of the thickness h of the liquidﬁlm.As a characteristic value of h we take the distance between the d and c phase at the symmetry axis of the simulation.In Fig.12the variation of the dimension-less quantity Vh=D is also shown,and it exhibits a linear behavior.Within the investigated range of D the behavior of Vh=D corresponds to a variation of2%of h,while the velocity exhibits a variation of approximatively50%.The quasilinear behavior of the steady-state velocity can be understood using the balance equation where the concen-tration gradient is approximately proportional to D(since the liquidﬁlm thickness depends weakly on D).We focus now on the solution of the problem in free space.An Ivantsov solution exists[21],consisting of two confocal parabolas,one for the melting front and one for the solidiﬁcation front.When making theappropriate Fig.9.Initial conﬁguration for T0>T P with D¼0:14and~c¼À1:15.Fig.10.TransformationðcþLÞ!ðdþLÞcorresponding to the initialconﬁguration shown in Fig.9.The velocity is indicated by the blackarrow.This transformation proceeds via a liquid-ﬁlm-migrationprocess.Fig.11.Liquid-ﬁlm-migration process for the transformation c!ðdþLÞwhere the metastable equilibrium consists of a single c phase,and theaverage concentration is exactly on the c solidus line.1756G.Boussinot et al./Acta Materialia58(2010)1750–1760。

材料学专业英语

加工方法Manufacturing Method 拉力强度Tensile Strength 机械性能Mechanical Properites低碳钢或铁基层金属Iron & Low Carbon as Base Metal 镀镍Nickel Plated 镀黄铜Brass Plated马氏铁体淬火Marquenching 退火Annealing 淬火Quenching高温回火High Temperature Tempering 应力退火温度Stress –relieving Annealing Temperature晶粒取向(Grain-Oriented)和非晶粒取向(Non-Oriented硬磁材料Hard Magnetic Material表面处理Surface Finish硬度Hardness 电镀方法Plating type 锌镀层质量Zinc Coating Mass表面处理Surface Treatment 拉伸应变Stretcher Strains焊接Welding 防止生锈Rust Protection硬度和拉力Hardness & Tensile strength test 连续铸造法Continuous casting process珠光体Pearlite 单相金属Single Phase Metal Ferrite渗碳体Cementitle奥氏体Austenite软磁Soft Magnetic硬磁Hard Magnetic疲劳测试Impact Test热膨胀系数Coefficient of thermal expansion 比重Specific gravity化学性能Chemical Properties物理性能Physical Properties 再结晶Recrystallization硬化Work Hardening包晶反应Peritectic Reaction包晶合金Peritectic Alloy 共晶Eutectic临界温度Critical temperature 自由度Degree of freedom相律Phase Rule金属间化物Intermetallic compound 固熔体Solid solution 置换型固熔体Substitutional type solid solution米勒指数Mill's Index晶体结构Crystal structure金属与合金Metal and Alloy金属特性Special metallic featuresStrength抗腐蚀和耐用Corrosion & resistance durability强度Strengthen 无机非金属inorganic nonmetallic materials 燃料电池fuel cell新能源new energy resources材料科学专业学术翻译必备词汇材料科学专业学术翻译必备词汇编号中文英文1 合金alloy2 材料material3 复合材料properties4 制备preparation5 强度strength6 力学mechanical7 力学性能mechanical8 复合composite9 薄膜films10 基体matrix11 增强reinforced12 非晶amorphous13 基复合材料composites14 纤维fiber15 纳米nanometer16 金属metal17 合成synthesis 18 界面interface19 颗粒particles20 法制备prepared21 尺寸size22 形状shape23 烧结sintering24 磁性magnetic25 断裂fracture26 聚合物polymer27 衍射diffraction 28 记忆memory29 陶瓷ceramic30 磨损wear31 表征characterization32 拉伸tensile33 形状记忆memory34 摩擦friction35 碳纤维carbon36 粉末powder37 溶胶sol-gel38 凝胶sol-gel39 应变strain40 性能研究properties41 晶粒grain42 粒径size43 硬度hardness44 粒子particles45 涂层coating46 氧化oxidation47 疲劳fatigue48 组织microstructure49 石墨graphite50 机械mechanical51 相变phase52 冲击impact53 形貌morphology54 有机organic55 损伤damage56 有限finite57 粉体powder58 无机inorganic59 电化学electrochemical60 梯度gradient61 多孔porous62 树脂resin63 扫描电镜sem64 晶化crystallizatio n65 记忆合金memory66 玻璃glass67 退火annealing68 非晶态amorphous 69 溶胶-凝胶sol-gel70 蒙脱土montmorillon ite71 样品samples 72 粒度size73 耐磨wear74 韧性toughness75 介电dielectric76 颗粒增强reinforced77 溅射sputtering78 环氧树脂epoxy79 纳米tio tio80 掺杂doped81 拉伸强度strength82 阻尼damping83 微观结构microstructure84 合金化alloying85 制备方法preparation86 沉积deposition87 透射电镜tem88 模量modulus89 水热hydrothermal90 磨损性wear91 凝固solidification92 贮氢hydrogen93 磨损性能wear94 球磨milling95 分数fraction96 剪切shear97 氧化物oxide98 直径diameter99 蠕变creep100弹性模量modulus留纞銅雀樓12:53:02101储氢hydrogen102压电piezoelectric103电阻resistivity104纤维增强composites 105纳米复合材料preparation 106制备出prepared107磁性能magnetic 108导电conductive 109晶粒尺寸size110弯曲bending111光催化tio 112非晶合金amorphous 113铝基复合材料composites 114金刚石diamond115沉淀precipitation116分散dispersion117电阻率resistivity118显微组织microstructure119sic复合材料sic120硬质合金cemented121摩擦系数friction122吸波absorbing123杂化hybrid124模板template125催化剂catalyst126塑性plastic127晶体crystal128sic颗粒sic129功能材料materials130铝合金alloy131表面积surface132填充filled133电导率conductivity134控溅射sputtering135金属基复合材料composites136磁控溅射sputtering137结晶crystallization138磁控magnetron139均匀uniform140弯曲强度strength141纳米碳carbon142偶联coupling143电化学性能electrochemical144和性能properties 145al复合材料composite 146高分子polymer147本构constitutive 148晶格lattice149编织braided150断裂韧性toughness 151尼龙nylon152摩擦磨损性friction 153耐磨性wear154摩擦学tribological 155共晶eutectic156聚丙烯polypropylene157半导体semiconductor158偶联剂coupling159泡沫foam160前驱precursor161高温合金superalloy162显微结构microstructure163氧化铝alumina164扫描电子显微镜sem165时效aging166熔体melt167凝胶法sol-gel168橡胶rubber169微结构microstructure170铸造casting171铝基aluminum172抗拉强度strength173导热thermal174透射电子显微镜tem175插层intercalation176冲击强度impact177超导superconducting178记忆效应memory179固化curing180晶须whisker181溶胶-凝胶法制sol-gel182催化catalytic183导电性conductivity184环氧epoxy185晶界grain186前驱体precursor 187机械性能mechanical 188抗弯strength189粘度viscosity190热力学thermodyna mic191溶胶-凝胶法制备sol-gel 192块体bulk 193抗弯强度strength194粘土clay 195微观组织microstructu re 196孔径pore197玻璃纤维glass198压缩compression199摩擦磨损wear200马氏体martensitic留纞銅雀樓12:53:57201制得prepared202复合材料性能composites203气氛atmosphere204制备工艺preparation205平均粒径size206衬底substrate207相组成phase208表面处理surface209杂化材料hybrid210材料中materials211断口fracture212增强复合材料composites213马氏体相变transformation214球形spherical215混杂hybrid216聚氨酯polyurethane217纳米材料nanometer218位错dislocation219纳米粒子particles220表面形貌surface221试样samples222电学properties223有序ordered224电压voltage225析出phase226拉伸性tensile227大块bulk 228立方cubic229聚苯胺polyaniline 230抗氧化性oxidation 231增韧toughening 232物相phase233表面改性modification 234拉伸性能tensile235相结构phase 236优异excellent237介电常数dielectric238铁电ferroelectric239复合材料力学性能composites240碳化硅sic241共混blends242炭纤维carbon243复合材料层composite244挤压extrusion245表面活性剂surfactant246阵列arrays247高分子材料polymer248应变率strain249短纤维fiber250摩擦学性能tribological251浸渗infiltration252阻尼性能damping253室温下room254复合材料层合板composite255剪切强度strength256流变rheological257磨损率wear258化学气相沉积deposition259热膨胀thermal260屏蔽shielding261发光luminescence262功能梯度functionally263层合板laminates264器件devices265铁氧体ferrite266刚度stiffness 267介电性能dielectric 268xrd分析xrd269锐钛矿anatase270炭黑carbon271热应力thermal272材料性能properties 273溶胶-凝胶法sol-gel 274单向unidirectiona l275衍射仪xrd 276吸氢hydrogen277水泥cement278退火温度annealing279粉末冶金powder280溶胶凝胶sol-gel281熔融melt282钛酸titanate283磁合金magnetic284脆性brittle285金属间化合物intermetallic286非晶态合金amorphous287超细ultrafine288羟基磷灰石hydroxyapatite289各向异性anisotropy290镀层coating291颗粒尺寸size292拉曼raman293新材料materials294tic颗粒tic295孔隙率porosity296制备技术preparation297屈服强度strength298金红石rutile299采用溶胶-凝胶sol-gel300电容量capacity301热电thermoelectric302抗菌antibacterial303聚酰亚胺polyimide304二氧化硅silica305放电容量capacity306层板laminates 307微球microspheres 308熔点melting309屈曲buckling 310包覆coated311致密化densification 312磁化强度magnetizatio n313疲劳寿命fatigue314本构关系constitutive 315组织结构microstructure316综合性能properties317热塑性thermoplastic318形核nucleation319复合粒子composite320材料制备preparation321晶化过程crystallization322层间interlaminar323陶瓷基ceramic324多晶polycrystalline325纳米结构nanostructures326纳米复合composite327热导率conductivity328空心hollow329致密度density330x射线衍射仪xrd331层状layered332矫顽力coercivity333纳米粉体powder334界面结合interface335超导体superconductor336衍射分析diffraction337纳米粉powders338磨损机理wear339泡沫铝aluminum340进行表征characterized341梯度功能gradient342耐磨性能wear343平均粒particle344聚苯乙烯polystyrene 345陶瓷基复合材料composites 346陶瓷材料ceramics 347石墨化graphitizatio n348摩擦材料friction 349熔化melting 350多层multilayer留纞銅雀樓12:55:33 351和其性能properties 352酚醛树脂resin353电沉积electrodeposition354分散剂dispersant355相图phase356复合材料界面interface357壳聚糖chitosan358抗氧化性能oxidation359钙钛矿perovskite360分层delamination361热循环thermal362氢量hydrogen363蒙脱石montmorillonite364接枝grafting365导率conductivity366放氢hydrogen367微粒particles368伸长率elongation369延伸率elongation370烧结工艺sintering371层合laminated372纳米级nanometer373莫来石mullite374磁导率permeability375填料filler376热电材料thermoelectric377射线衍射ray378铸造法casting379粒度分布size380原子力afm381共沉淀coprecipitation382水解hydrolysis 383抗热thermal 384高能球ball385干摩擦friction 386聚合物基polymer 387疲劳裂纹fatigue388分散性dispersion 389硅烷silane390弛豫relaxation 391物理性能properties 392晶相phase 393饱和磁化强度magnetization394凝固过程solidification395共聚物copolymer396光致发光photoluminescence397薄膜材料films398导热系数conductivity399居里curie400第二相phase401复合材料制备composites402多孔材料porous403水热法hydrothermal404原子力显微镜afm405压电复合材料piezoelectric406尼龙6nylon407高能球磨milling408显微硬度microhardness409基片substrate410纳米技术nanotechnology411直径为diameter412织构texture413氮化nitride414热性能properties415磁致伸缩magnetostriction416成核nucleation417老化aging418细化grain419压电材料piezoelectric420纳米晶amorphous421si合金si422复合镀层composite 423缠绕winding 424抗氧化oxidation 425表观apparent 426环氧复合材料epoxy 427甲基methyl428聚乙烯polyethylene 429复合膜composite 430表面修饰surface431大块非晶amorphous 432结构材料materials 433表面能surface434材料表面surface435疲劳性能fatigue436粘弹性viscoelastic437基体合金alloy438单相phase439梯度材料material440六方hexagonal441四方tetragonal442蜂窝honeycomb443阳极氧化anodic444塑料plastics445超塑性superplastic446sem观察sem447烧蚀ablation448复合薄膜films449树脂基resin450高聚物polymer451气相vapor452电子能谱xps453硅烷偶联coupling454团聚particles455基底substrate456断口形貌fracture457抗压强度strength458储能storage459松弛relaxation460拉曼光谱raman461孔率porosity462沸石zeolite463熔炼melting464磁体magnet465sem分析sem466润湿性wettability 467电磁屏蔽shielding 468升温heating469致密dense470沉淀法precipitation 471差热分析dta472成功制备prepared 473复合体系composites 474浸渍impregnation 475力学行为behavior 476复合粉体powders 477沥青pitch478磁电阻magnetoresistance479导电性能conductivity480光电子能谱xps481材料力学mechanical482夹层sandwich483玻璃化glass484衬底上substrates485原位复合材料composites486智能材料materials487碳化物carbide488复相composite489氧化锆zirconia490基体材料matrix491渗透infiltration492退火处理annealing493磨粒wear494氧化行为oxidation495细小fine496基合金alloy497粒径分布size498润滑lubrication499定向凝固solidification500晶格常数lattice留纞銅雀樓12:56:20501晶粒度size502颗粒表面surface503吸收峰absorption504磨损特性wear505水热合成hydrothermal506薄膜表面films507性质研究properties 508试件specimen 509结晶度crystallinity 510聚四氟乙烯ptfe511硅烷偶联剂silane 512碳化carbide 513试验机tester514结合强度bonding 515薄膜结构films516晶型crystal517介电损耗dielectric 518复合涂层coating519压电陶瓷piezoelectric520磨损量wear521组织与性能microstructure522合成法synthesis523烧结过程sintering524金属材料materials525引发剂initiator526有机蒙脱土montmorillonite527水热法制hydrothermal528再结晶recrystallization529沉积速率deposition530非晶相amorphous531尖端tip532淬火quenching533亚稳metastable534穆斯mossbauer535穆斯堡尔mossbauer536偏析segregation537种材料materials538先驱precursor539物性properties540石墨化度graphitization541中空hollow542弥散particles543淀粉starch544水热法制备hydrothermal545涂料coating546复合粉末powder547晶粒长大grain548sem等sem549复合材料组织microstructu re550界面结构interface 551煅烧calcined 552共混物blends553结晶行为crystallizatio n554混杂复合材料hybrid 555laves相laves556摩擦因数friction 557钛基titanium558磁性材料magnetic559制备纳米nanometer560界面上interface561晶粒大小size562阻尼材料damping563热分析thermal564复合材料层板laminates565二氧化钛titanium566沉积法deposition567光催化剂tio568余辉afterglow569断裂行为fracture570颗粒大小size571合金组织alloy572非晶形成amorphous573杨氏模量modulus574前驱物precursor575过冷alloy576尖晶石spinel577化学镀electroless578溶胶凝胶法制备sol-gel579本构方程constitutive580磁学magnetic581气氛下atmosphere582钛合金titanium583微粉powder584压电性piezoelectric585晶须sic586应力应变strain587石英quartz588热电性thermoelectric589相转变phase590合成方法synthesis 591热学thermal 592气孔率porosity 593永磁magnetic 594流变性能rheological 595压痕indentation 596热压烧结sintering 597正硅酸乙酯teos 598点阵lattice599梯度功能材料fgm 600带材tapes601磨粒磨损wear602碳含量carbon603仿生biomimetic604快速凝固solidification605预制preform606差示dsc607发泡foaming608疲劳损伤fatigue609尺度size610镍基高温合金superalloy611透过率transmittance612溅射法制sputtering613结构表征characterization614差示扫描dsc615通过semsem616水泥基cement617木材wood618分析tem619量热calorimetry620复合物composites621铁电薄膜ferroelectric622共混体系blends623先驱体precursor624晶态crystalline625冲击性能impact626离心centrifugal627断裂伸长率elongation628有机-无机organic-inorganic629块状bulk630相沉淀precipitation631织物fabric632因数coefficient 633合成与表征synthesis 634缺口notch635靶材target636弹性体elastomer 637金属氧化物oxide 638均匀化homogenizati on639吸收光谱absorption 640磨损行为wear641高岭土kaolin 642功能梯度材料fgm643滞后hysteresis644气凝胶aerogel645记忆性memory646磁流体magnetic647铁磁ferromagnetic648合金成分alloy649微米micron650蠕变性能creep留纞銅雀樓12:56:46651聚氯乙烯pvc652湮没annihilation653断裂力学fracture654滑移slip655差示扫描量热dsc656等温结晶crystallization657树脂基复合材料composite658阳极anodic659退火后annealing660发光性properties661木粉wood662交联crosslinking663过渡金属transition664无定形amorphous665拉伸试验tensile666溅射法sputtering667硅橡胶rubber668明胶gelatin669生物相容性biocompatibility670界面处interface671陶瓷复合材料composite 672共沉淀法制coprecipitatio n673本构模型constitutive 674合金材料alloy675磁矩magnetic 676隐身stealth677比强度strength 678改性研究modification 679采用粉末powder 680晶粒细化grain681抗磨wear682元合金alloy683剪切变形shear684高温超导superconducting685金红石型rutile686晶化行为crystallization687催化性能catalytic688热挤压extrusion689微观microstructure690tem观察tem691缺口冲击impact692生物材料biomaterials693涂覆coating694纳米氧化nanometer695x射线光电子能谱xps696硅灰石wollastonite697摩擦条件friction698衍射峰diffraction699块体材料bulk700溶质solute701冲击韧性impact702锐钛矿型anatase703凝固组织microstructure704磨损试验机tester705丙烯酸甲酯pmma706光谱raman707减振damping708聚酯polyester709体材料materials 710航空aerospace 711光吸收absorption 712韧化toughening 713疲劳裂纹扩展fatigue 714超塑superplastic 715凝胶法制备gel716半导体材料semiconduct or717剪应力shear718发光材料luminescence 719凝胶法制gel720甲基丙烯酸甲酯pmma721硬质hard722摩擦性能friction723电致变色electrochromic724超细粉powder725增强相reinforced726薄带ribbons727结构弛豫relaxation728光学材料materials729sic陶瓷sic730纤维含量fiber731高阻尼damping732镍基nickel733热导thermal734奥氏体austenite735单轴uniaxial736超导电性superconductivity737高温氧化oxidation738树脂基体matrix739含能energetic740粘着adhesion741穆斯堡尔谱mossbauer742脱层delamination743反射率reflectivity744单晶高温合金superalloy745粘结bonded746快淬quenching747熔融插层intercalation748外加applied749钙钛矿结构perovskite750减摩friction751复合氧化物oxide21 / 21。

颗粒稳定泡沫法制备莫来石基多孔陶瓷的结构与性能

n(Al):n(Si)
2.0:1.0 2.2:1.0 2.4:1.0 3.0:1.0 n—Mole.
表 1 实验参量设计 Table 1 Experimental parameters
Solid content Gelatin addition PG addition (accounting for (account-ing for (accounting for total volume)/% solid mass)/% alumina mass)/%
资助。第一作者：王涵(1994—)，女，硕士研究生。通信作者：李翠伟(1973—)，女，博士，教授。
Received date: 2019–07–17. Revised date: 2019–08–18. First author: WANG Han (1994–), female, Master candidate. E-mail: 17121355@ Correspondent author: LI Cuiwei (1973–), female, Ph.D., Professor. E-mail: cwli@
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《硅酸盐学报》 J Chin Ceram Soc, 2020, 48(6): 824–832
2020 年
图 1 颗粒稳定泡沫法制备莫来石基多孔陶瓷工艺流程图 Fig. 1 Flow chart for preparation of mullite-based porous ceramics by particle-stabilized foaming method
Keywords: mullite; porous ceramics; particle-stabilized; freeze drying; thermal conductivity

AlPO4-coated mullitealumina fiber reinforced

Available online at Journal of the European Ceramic Society28(2008)3041–3048AlPO4-coated mullite/aluminaﬁber reinforcedreaction-bonded mullite compositesYahua Bao∗,Patrick S.NicholsonCeramic Engineering Research Group,Department of Materials Science and Engineering,McMaster University,Hamilton,Ontario,L8S4L7CanadaReceived23February2008;received in revised form27May2008;accepted30May2008Available online22July2008AbstractA precursor for reaction-bonded mullite(RBM)is formulated by premixing Al2O3,Si,mullite seeds and mixed-rare-earth-oxides(MREO).An ethanol suspension thereof is stabilized with polyethyleneimine protonated by acetic acid.The solid in the suspension is inﬁltrated into unidirectional mullite/aluminaﬁber-preforms by electrophoretic inﬁltration deposition to produceﬁber-reinforced,RBM green bodies.Crack-free composites with≤25%porosity were achieved after pressureless sintering at1300◦C.Pre-coating theﬁbers with AlPO4as a weak intervening layer facilitates signiﬁcantﬁber pullout on composite fracture and confers superior damage tolerance.The bend strength is∼170MPa at25◦C≤T≤1100◦C.At 1200◦C,the composite fails in shear due to MREO-based,glassy phase formation.However,the AlPO4coating acts as a weak layer even after thermal aging at1300◦C for100h.©2008Elsevier Ltd.All rights reserved.Keywords:AlPO4;Weak layer;Nextel720ﬁber;Reaction-bonded mullite;Ceramic matrix composites1.IntroductionFiber-reinforced ceramic matrix composites(CMC’s)are candidates for structural application due to their damage toler-ance.Currently,most denseﬁber-reinforced CMC’s are based on non-oxide systems which oxidize in air at high temperatures. Oxidation-resistant,ﬁber-reinforced CMC’s are required for use in oxidizing atmospheres at high temperatures.To optimize the fracture work necessary to breakﬁber-reinforced CMC’s, the bonding betweenﬁber and matrix must allowﬁber pullout from the matrix via a weak layer therebetween.1LaPO4has been reported as one of the most successful dense weak layer candidates.2–4However,Bao and Nicholson5recently reported application of another phosphate,AlPO4,as the weak layer on theﬁber surface.They demonstrated signiﬁcantﬁber pullout on the fracture surface of hot-pressed,AlPO4-coated mullite/alumina(Nextel TM720)ﬁber-reinforced Al2O3.Highly covalently-bonded AlPO4displays poor sintering behavior even sintered at1550◦C.5Thus it should be a stable porous weak ∗Corresponding author.Current address:SII MegaDiamond,275West2230 North,Provo,UT84604,United States.E-mail address:baoyahua@(Y.Bao).layer onﬁber surface for crack deﬂection andﬁber e of a reaction-bonded matrix with near-zero sintering shrinkage should avoid the necessity to hot-press.Mullite is an ideal matrix at elevated temperatures because of high temperature strength,low thermal expansion coefﬁ-cient and good creep resistance.Due to the volume stability of mullite/aluminaﬁbers,the matrix sintering shrinkage must be low to avoid cracks on pressureless sintering.Reaction-bonded mullite(RBM)explored as near-zero shrinkage is achieved by mixing alumina,silicon and aluminum precursors.6–8In reaction-bonded mullite of composition3Al2O3–2Si,the pre-cursor Si oxidizes to SiO2at high temperatures which reacts with the Al2O3to form mullite.Two volume expansion reac-tions are involved,i.e.,Si→SiO2(+134%),and3Al2O3–2SiO2→mullite(+1.3%).The reaction shrinkage can be the-oretically calculated as;s=1−1.013+1.340×1.013×V SiV Al2O3+V Siρ0ρ1/3(1)whereρ0andρare the theoretical-green andﬁred-densities (V Al2O3and V Si are the volume fractions of Al2O3and Si powder in the mixture,respectively).Generally,after green process-ing,the ceramic density is∼55%.For3%sintering shrinkage,0955-2219/$–see front matter©2008Elsevier Ltd.All rights reserved. doi:10.1016/j.jeurceramsoc.2008.05.0323042Y.Bao,P.S.Nicholson/Journal of the European Ceramic Society28(2008)3041–3048theﬁred density can be80%theoretical.Thus reaction-bonded-mullite with less than20%porosity and less than3%shrinkage should not be problem if the SiO2comes from the Si precursor, i.e.,ﬁber-reinforced,reaction-bonded-mullite composites can be realistically fabricated by pressureless sintering free of macro cracks.The mullite-formation temperature is∼1500◦C,how-ever,the strength of the mullite/alumina(Nextel TM720)ﬁber degrades severely on heat-treatment>1300◦C.9Thus the mul-lite matrix sintering temperature must be modiﬁed to≤1300◦C forﬁber strength retention.Recently,rare earth oxides added to RBM reduced the mullite-formation temperature to1350◦C.10 So mixed-rare-earth-oxides(MREO)were added to the RBM mixture to form mullite<1300◦C.11In this paper,AlPO4-coated,Nextel720ﬁber/reaction-bonded mullite composites are fabricated by constant-current electrophoretic-inﬁltration-deposition(EPID)and subsequently pressureless-sintered.2.ExperimentalReaction-bonded mullite was prepared with submicron TM-DAR alumina powder(Taimei Chemicals,Tokyo,Japan)and micron Si metal powder(Atlantic Equipment Engineers,Ner-genﬁeld,NJ)which was pre-ball-milled in ethanol to a surface area>20m2/g(to promote oxidation during sintering).Table1 lists composition of mixed-rare-earth-oxides(MREO,Lan-thanide oxide,Molycorp,Fairﬁeld,NJ).These were added as sintering and mullite-formation aids.A mullite precursor (Siral28M,SASOL GmbH,Hamburg,German),pre-sintered at1300◦C for2h to form pure mullite,was ground and added as mullite-promotion seeds.The molar ratio of Al:Si was set to that of mullite.Anhydrous polyethyleneimine dispersant(PEI,M.W.10,000, Polysciences,Warrington,PA),protonated with glacial acetic acid was added to stabilize the RBM-precursor,ethanol suspension.The optimal addition was determined via the elec-trophoretic mobility value for the RBM precursors with a zeta potential analyzer(ZetaPALS,Brookhaven Instruments, Holtsville NY).The electrokinetic sonic amplitude(ESA)was measured on the mixed suspension(ESA-8000,Matec Applied Sciences,Hopkinton,MA).RBM pellets were also uniaxi-ally pressed then cold isostatically pressed at140MPa and heated to1175–1200◦C for10h in air to oxidize the Si to SiO2.Finally they were sintered at1250–1350◦C for2h to study mullite phase formation and shrinkage.Shrinkage was calculated from the change of the diameter of the pel-lets.Table1Composition of the mixed-rare-earth-oxides(MREO)provided by Molycorp Oxide Concentration(wt%)CeO249La2O333Nd2O313Pr6O114Other1AlPO4was coated onto the mullite/aluminaﬁbers(Nextel TM 720,3M,St.Paul,MN)by a layer-by-layer electrostatic method.5Desizedﬁbers were pre-treated with0.5wt%cationic polyelectrolyte solution(polydiallyldimethylammonium chlo-ride,Aldrich,M.W.400,000–500,000)to induce a positive surface charge.The latter attracts the negatively-charged AlPO4 nano-particles to produce the coating.The coatedﬁbers were heat-treated at1100◦C.AlPO4was coated for10cycles (∼10wt%gain)to give an acceptable thickness.The coated ﬁbers were unidirectionally mounted in a rectangular,plastic holder(25mm×5mm×3mm),the back of which was attached to a metal plate as cathode to draw particles through theﬁber preform and accomplish electrophoretic-inﬁltration-deposition (EPID).12,13The inter-electrode distance was2cm and EPID was conducted at a constant current of0.07mA/cm2.After depo-sition,the composite was cold isostatically pressed at140MPa then dried in an atmosphere-controlled closed container to avoid cracking.Uncoatedﬁbers were also inﬁltrated with RBM matrix for comparison.The green composites were heated to1175◦C for10h in air to convert the Si to SiO2then sintered at1300◦C for2h.Their thermal stability was determined by heat-treatment for100h at 1300◦C.Theﬁred density and open porosity were measured by Archimedes’method in water.Four-point bend testing was performed at0.10mm/min in a screw-driven,ultra-hard compression machine(Model10053 &10055,Wykeham Farrance Engineering Ltd.,UK)using an aluminaﬁxture with outer span,20mm,and inner span,10mm. The sample thickness was2.0–2.5mm.Fracture surfaces were observed by SEM and the degree ofﬁber-pullout determined. Ten samples were tested at room temperature andﬁve at elevated temperatures.3.Results and discussionFig.1shows the DTA curve for32wt%MREO,22wt% Al2O3and46wt%SiO2mixture.The endothermic peak around 1200◦C is due to eutectic liquid-phase formation.Mullite phase can be promoted by formation of the MREO–Al2O3–SiO2eutec-tic liquid phase.Fig.2tracks the phase evolution in the RBM matrix versus sintering temperature for a mixture containing 7.5wt%MREO.Mullite appears at1270◦C,and is the major phase at1300◦C.Traces of alumina and silica remain.Thesil-Fig.1.DTA curve for32wt%MREO,22wt%Al2O3and46wt%SiO2mixture.Y.Bao,P .S.Nicholson /Journal of the European Ceramic Society 28(2008)3041–30483043Fig.2.Phase evolution of the reaction-bonded mullite containing 7.5wt%MREO.ica is totally consumed >1300◦C but the alumina trace persists.The mullitization temperature was further decreased by adding mullite “seeds”(Fig.3).With 0.5wt%seeds,the mullitization process was complete at 1270◦C.However,the sintering shrink-age (6–8%)was still too large (should be <∼3%to avoid matrix cracks around the volume stable ﬁbers).Fig.4tracks the effect of mullite seed content on the shrinkage,density and open porosity,for RBM with 7.5wt%MREO.A small amount of mullite seeds can prompt the formation of mullite.However,when seed con-centration is high,it serves as refractory inclusion and retards the RBM sintering.Fig.5shows the dilatometric measurement curve for RBM with 7.5wt%MREO and 5.0wt%seeds.Length expansion occurs due to oxidation of Si metal powder and reaches maximum (<3%)at 1170◦C,close to the eutectic point of MREO–Al 2O 3–SiO 2.Formation of MREO–Al 2O 3–SiO 2liq-uid phase prompts the sintering shrinkage and a steep drop in length was detected at 1170–1200◦C.Si metal is completely oxi-dized into silica when soaked at 1200◦C for 10h.WithfurtherFig.3.Effect of mullite seed addition on the mullitizationtemperature.Fig.4.Inﬂuence of mullite seed addition on the sintering shrinkage,density and open porosity for RBM sintered at 1300◦C for 2h.increase of temperature,RBM sintering shrinkage takes place and ﬁnal length change is <3%.When [seeds]=5wt%,the RBM shrinkage is <3%and open porosity,<20%.This composition is optimum and is used in the EPID processing.Alumina particles in ethanol are positively charged whereas Si particles are negatively charged.These particles tend to hetero-coagulate.A dispersant was added to induce a com-mon surface charge sign and stabilize the mixed suspension.PEI,protonated with acetic acid,adsorbs on both surfacesren-Fig.5.Dilatometric measurement of RBM with 7.5wt%MREO and 5.0wt%seeds.3044Y.Bao,P.S.Nicholson/Journal of the European Ceramic Society28(2008)3041–3048Fig.6.Effect of PEI on mobility.dering them the same sign.Fig.6illustrates the inﬂuence ofPEI on the RBM-precursor-particle-mobility in ethanol.When[PEI]>0.15wt%,all particles are positively charged.No changeof particle mobility is observed for[PEI]>0.4wt%.Fig.7shows the electrokinetic sonic amplitude(ESA)values for a2.5vol%mixed suspension of alumina,Si,MREO and mulliteseeds(3Al2O3–2Si+5wt%mullite seeds+7.5wt%MREO)ver-sus PEI concentration.The ESA value increases with PEI andplateaus at[PEI]>0.3wt%.Thus0.5wt%PEI was used to ensurewell-dispersed suspensions for the EPID process.As PEI alsoabsorbs on the AlPO4-coatedﬁber surface,it renders theﬁbersrepulsive to the particles.Theﬁber adsorption of PEI causesparticles passing through them to be repelled as they pass,i.e.,particles are“streamed”.Fig.8shows the morphology of an AlPO4-coated Nextel720ﬁber-reinforced composite prepared by EPID(current density0.07mA/cm2)pressurelessly sintered at1300◦C for2h.Sin-teredﬁber-reinforced composites have∼25%open porosity and ∼25vol%ﬁber.Macro cracks do not occur and particles are well inﬁltrated into theﬁber preform.After polishing,the weakly-bonded AlPO4coating polished away and grooves appearedaround theﬁbers.An uncoated Nextel720/RBM compositewas employed for matrix phase analysis due to the very similarXRD patterns for AlPO4and SiO2.Fig.9is an XRD pattern foruncoated Nextel720Fiber/RBM composite sintered at1300◦CFig.7.Effect of PEI on ESA of2.5vol%suspension.Fig.8.Morphology ofﬁber/RBM composite prepared by EPID.for2h.It is a mixture of alumina,cristobalite and mullite,sug-gesting non-uniform distribution of the MREO in the matrix.MREO promotes matrix mullite formation via low temperatureeutectics formed with alumina and silica.11As the MREO pow-der has high density,it also may sediment duringﬁberpreformFig.9.X-ray diffraction pattern ofﬁber/RBM composite prepared by EPID andsintered at1300◦C.Y.Bao,P.S.Nicholson/Journal of the European Ceramic Society28(2008)3041–30483045Fig.10.Effect of AlPO4coating on fracture surface ofﬁber/RBM composites tested at R.T.(a)without AlPO4coating and(b)with AlPO4coating.inﬁltration.In fact,a yellow layer was noted on the base of a green body after inﬁltration.A non-uniform distribution,or, less-MREO-present-than-designed will locally retard formation of mullite at1300◦C.Fig.10compares the effect of AlPO4-coating onﬁber pullout, i.e.,without AlPO4on theﬁbers,strong bonds form between theﬁbers and the RBM matrix so that a planar fracture sur-face results.However,when AlPO4is coated on theﬁbers, signiﬁcantﬁber pullout is observed.The high covalent bond-ing level in AlPO4retards its sinterability so AlPO4ceramics remain very porous even when sintered at1550◦C.5Thus the inherently-porous AlPO4coating on Nextel720ﬁber surface, serves as a porous weak layer for crack deﬂection andﬁber pull-out.Fig.11show porous AlPO4coating attached to the matrix after theﬁber pullout,indicating that the AlPO4/ﬁber bonding is also weak and cracks deﬂect therefrom.Fig.12illustrates theﬁber bridging effect.Though the crack opening distance is ∼150␮m,theﬁbers still bridge across it.Fiber pullout reaches ∼150␮m(>10times theﬁber diameter).The pullout length is shorter on fracture at1100◦C(Fig.13).Fig.14showstheFig.11.Porous AlPO4coating after sintering at1300◦C for2h. Fig.12.Fiber bridging and pullout across a matrix crack in AlPO4-coated ﬁber/RBM composite tested at room temperature.3046Y.Bao,P .S.Nicholson /Journal of the European Ceramic Society 28(2008)3041–3048Fig.13.Fracture surface of AlPO 4-coated-ﬁber/RBM composite with AlPO 4weak layer tested at 1100◦C.Fig.14.4-point bending strength of AlPO 4-coated-ﬁber/RBM composites tested at R.T.and 1100◦C.load/cross-head-displacement curves for the composite tested at room temperature and 1100◦C.The composite exhibits dam-age tolerance at both temperatures.The ultimate bend strength of the composite is 175±20MPa and 170±25MPa at room temperature and 1100◦C,respectively.At 1200◦C,the MREO-induced MREO–Al 2O 3–SiO 2glassy phase occurs in the matrix,thus the composite fails in shear at 1200◦C (Fig.15).The ulti-mate bend strength of RBM and AlPO 4-coated ﬁber/RBM is listed in Table 2.Antti et al.14reported thermal degradation of commercial ﬁber reinforced porous aluminosilicate matrix composites.AfterTable 2The bend strengths of RBM and AlPO 4-coated Nextel 720ﬁber/RBM composites SampleBend strength (MPa)25◦C1100◦C 1200◦C RBM105±1570±1017±5AlPO 4-coated ﬁber/RBM175±20170±25ShearfailureFig.15.A morphology of an AlPO 4-coated-ﬁber/RBM composite tested at 1200◦C.exposure at 1100◦C in air,the composite embrittles due to localized matrix densiﬁcation and increased bonding of ﬁbers therewith.Although the RBM matrix is still with ∼20%poros-ity,the composite thermal stability can be signiﬁcantly increased with AlPO 4weak layer coating.Fig.16shows the load/cross-head-displacement curve for an AlPO 4-coated Nextel 720/RBM composite after heat-treatment at 1300◦C for 100h.The com-posite still exhibits damage tolerance with a bend strength of 160MPa.The AlPO 4coating is still porous (Fig.17).Almost no AlPO 4grain growth occurs during heat-treatment at 1300◦C for 100h.But the ﬁber pullout length is shorter (Fig.18).Short ﬁber pullout length should be mostly due to the severe ﬁber strength degradation on thermal aging at 1300◦C.The effect of thermal aging on interfacial bonding needs to be further evaluated.Kriven and Lee 15proposed phase transformation causes weakening in mullite/cordierite laminates with ␤-cristobalite (SiO 2)as the interface.The ␣↔␤,“cristobalite”,AlPO4Fig.16.Load/cross-head displacement curve for AlPO 4-coated-ﬁber/RBM composite tested at room temperature after heat-treatment at 1300◦C for 100h.Y.Bao,P .S.Nicholson /Journal of the European Ceramic Society 28(2008)3041–30483047Fig.17.Porous AlPO 4coating around ﬁber after heat-treatment at 1300◦C for 100h.transformation occurs at 220◦C with 4.6%volume change.Microcracks due to the phase transformation are minimized as the coating is porous.The AlPO 4coating should be pure “␤-cristobalite”phase at 1100◦C,however,it will still serve as a weak layer between the ﬁbers and matrix.Thus,phase-transformation weakening in porous AlPO 4coating can be neglected.The smooth pullout ﬁber surface and AlPO 4coat-ing attached to the matrix suggest weak bonding between the ﬁbers and the AlPO 4coating even after heating to high temper-atures.Therefore,an approaching crack will deﬂect along the ﬁber/AlPO 4surface.The interfacial sliding resistance depends on the AlPO 4-coating elastic modulus.A low elastic modulus signiﬁcantly decreases the ﬁber/coating sliding resistance,promoting ﬁber pullout.16,17Elastic modulus is a function of porosity 18;E p =E (1−1.9f p +0.9f 2p )where E and E p are the elastic modulus of the fully dense and porous materials,respectively and f p the porosity.TheelasticFig.18.Fiber pullout after heat-treatment at 1300◦C for 100h,fractured at R.T.modulus of AlPO 4is 57GPa,19so,assuming the porosity is approximately that of sintered AlPO 4,i.e.,∼30%,5the coating elastic modulus is ∼29GPa.This low value explains why sig-niﬁcant ﬁber pullout occurs from the matrix.It can be concluded that the high covalent bonding (poor sinterability)in AlPO 4is key to its performance on the ﬁber pullout.The latter results from both the low elastic modulus of porous AlPO 4coating and the weak bonding between the ﬁbers and the AlPO 4coating.4.SummaryReaction-bonded mullite was formed at <1300◦C via incor-poration of 7.5wt%mixed-rare-earth-oxides into an Al 2O 3–Si mixture.Inclusion of 5wt%mullite seeds decreased the sinter-ing shrinkage to <2%and give a density ∼2.6g/cm 3with open porosity <20%.A PEI dispersant produced a stable ethanol sus-pension of the reaction-bonded mullite precursor.Crack free,unidirectional ﬁber-reinforced RBM composites with 25vol%ﬁbers and 25%porosity were achieved by EPID followed by pressureless sintering at 1300◦C.AlPO 4was coated onto mul-lite/alumina ﬁbers and the ﬁber/RBM composites exhibited superior damage tolerance with signiﬁcant ﬁber pullout between room temperature and 1100◦C.The ultimate bend strength of the composites was ∼170MPa.At 1200◦C,the composite failed in shear due to glassy phase formation in the matrix.The composite still displayed damage tolerance with ﬁber pullout after thermal aging at 1300◦C for 100h,and testing at room temperature.It is concluded AlPO 4is an effective oxidation-resistant,weak layer between the ﬁbers and matrix of oxide-ﬁber/oxide-matrix CMC’s.AcknowledgementYahua Bao would like to thank Prof.D.S.Wilkinson and Prof.J.Barbier for fruitful discussions.Referencesmittee on Advanced Fibers for High-Temperature Ceramic Compos-ites,National Research Council,Ceramic Fibers and Coatings:Advanced Materials for the Twenty-First Century .NMAB-494,National Academies Press,Washington,DC,1998.2.Morgan,P.E.D.and Marshall,D.B.,Ceramic composites of mon-azite and alumina.Journal of the American Ceramic Society ,1995,78(6),1553–1563.3.Marshall,D.B.,Davis,J.B.,Morgan,P.E.D.,Waldrop,J.R.and Porter,J.R.,Properties of La-monazite as an interphase in oxide composites.Zeitschrift Fur Metallkunde ,1999,90(12),1048–1052.4.Kerans,R.J.,Hay,R.S.,Parthasarathy,T.A.and Cinibulk,M.K.,Interface design for oxidation-resistant ceramic composites.Journal of the American Ceramic Society ,2002,85(11),2599–2632.5.Bao,Y .and Nicholson,P.S.,AlPO 4coating on alumina/mullite ﬁbers as a weak interface in ﬁber-reinforced oxide composites.Journal of the American Ceramic Society ,2006,89(2),465–470.6.Wu,S.X.and Claussen,N.,Fabrication and properties of low-shrinkage reaction-bonded mullite.Journal of the American Ceramic Society ,1991,74(10),2460–2463.7.Wu,S.and Claussen,N.,Reaction bonding and mechanical properties of mullite/silicon carbide composites.Journal of the American Ceramic Soci-ety ,1994,77(11),2898–2904.3048Y.Bao,P.S.Nicholson/Journal of the European Ceramic Society28(2008)3041–30488.Holz,D.,Pagel,S.,Bowen,C.,Wu,S.and Claussen,N.,Fabrication oflow-to-zero shrinkage reaction-bonded mullite composites.Journal of the European Ceramic Society,1996,16(2),255–260.9.Petry,M.D.and Mah,T.I.,Effect of thermal exposures on the strengths ofNextel TM550and720ﬁlaments.Journal of the American Ceramic Society, 1999,82(10),2801–2807.10.Mechnich,P.,Schneider,H.,Schmucker,M.and Saruhan,B.,Acceleratedreaction bonding of mullite.Journal of the American Ceramic Society,1998, 81(7),1931–1937.11.Kim,H.S.and Nicholson,P.S.,Use of mixed-rare-earth oxide in the prepa-ration of reaction-bonded mullite at≤1300◦C.Journal of the American Ceramic Society,2002,85(7),1730–1734.12.Bao,Y.and Nicholson,P.S.,Electrophoretic Inﬁltration deposition forﬁber-reinforced ceramic composites under constant current.Journal of the American Ceramic Society,2007,90(4),1063–1070.13.Y.Bao,Strong,Damage-Tolerant Oxide-Fiber/Oxide-Matrix Composites.Ph.D.Thesis,McMaster University,Hamilton,Ontario,Canada2006.14.Antti,M.-L.,Lara-Curzio,E.and Warren,R.,Thermal degradation of anoxideﬁbre(Nextel720)/aluminosilicate composite.Journal of the European Ceramic Society,2004,24(3),565–578.15.Kriven,W.M.and Lee,S.J.,Toughening of mullite/cordierite laminatedcomposites by transformation weakening of beta-cristobalite interphases.Journal of the American Ceramic Society,2005,88(6),1521–1528.16.Hsueh,C.-H.,Becher,P.F.and Angelini,P.,Effects of interfacialﬁlms onthermal stresses in whisker-reinforced ceramics.Journal of the American Ceramic Society,1988,71(11),929–933.17.Kerans,R.J.,Viability of oxideﬁber coatings in ceramic composites foraccommodation of misﬁt stresses.Journal of the American Ceramic Society, 1996,79(6),1664–1668.18.Kingery,W.D.,Bowen,H.K.and Uhlmann,D.R.,Introduction to Ceramics.John Wiley&Sons,New York,1976.19.Hanada,T.,Bessyo,Y.and Soga,N.,Elastic-constants of amorphousthin-ﬁlms in the systems SiO2-Al2O3and AlPO4-Al2O3.Journal of Non-Crystalline Solids,1989,113(2–3),213–220.。

无机非金属材料专业英语

被铝取代的氧O2-ion replaced by Al3+比热specific heat波函数wave function玻璃态的vitreous玻璃组成glass composition 不完整的配位incomplete coordination长石feldspar成对电子paired electrons 初晶相the primary phase 磁光效应magneto-optic effect缔合缺陷associated defects 电导conductivity电光效应electro-optic effect电子空穴electron holes电子排布electronconfiguration断裂韧性fracture-toughness二价阳离子divalentcation钙铝硅酸盐玻璃calcium-aluminateglass刚性体rigid body锆英石zircon共沉淀和过饱和coprecipitationandsupersaturation共价键covalent bonds固体电解质Solid electrolyte硅铝酸盐alumina-silica红外投射infraredtransmission互溶体mutual solution化学方程式chemical formulate碱金属alkali metal碱金属硅酸盐玻璃alkali silicateglass碱金属卤化物hailde of alkalimetals角连接的硅氧四面体[SiO4]tetrahedrawith shared corners介电常数、强度、损耗dielectricconstant、strength、losses紧密堆积结构closed-packedstructure近似立方紧密堆积nearly cubicclose-packedstructure净化工艺purificatinprocedures颗粒尺寸分布particle sizedistribution颗粒的重排和团聚particlerearrangement andagglomerate快离子导体Fast ion conductor冷却速率cooling rate离子键ionic bonds链状排列chain arrangement莫来石mullite母体玻璃parent glass钠钙硅玻璃soda-lime-silicaglass配位数coordinationnumber喷雾干燥和煅烧spray-drying andcalcination缺乏absence of缺陷化学defect chemistry热历史the thermal history热能thermal energy热膨胀系数thermal expansioncoefficient熔点melting point软化范围softening range三元系统the ternary system受控结晶controlledcrystallization水软化water softener四面体tetrahedron体积核化volume nucleation退火玻璃annealing glass退火和烧结温度annealing andsinteringtemperature网络结构network structure网络条整体network modifier相图phase diagram学说theory学说解释account for压敏电阻和热敏电阻varistor andthermistor亚原子粒子subatomicparticles衍生结构derivationstructure阳离子cation氧化锆陶瓷zirconia-basedceramics氧离子oxygen ions液相温度liquidustemperature一价阴离子univalent anion异质核化heter ogeneousnucleation阴离子anion阴离子空位vacant anion sites有效电荷effective charges折射率和色散index of refractionand dispersion中间体intermediate转变温度transmissiontemperatureact as作为，冲当aggregation of finepowder细粉团聚alumina-silica铝硅酸盐as compared to与…比较ball-milled powers球磨粉末be based on以…为基础be regarded as被认为是chanrgedinterstitial site带电间隙位chemical formulate化学方程式cohesive fore内聚力commence with从……开始effectivelyneutral charge有效中性点荷fireclay products黏土烧制产品framework框架结构glassy andcrystalline grainboundary phases玻璃相和晶界相hexagonalclosed-packedstructure六方紧密堆积结构host lattice主晶格hot uniaxialpressing单轴热压hybridization ofthe atomic orbitals原子杂化轨道in particular of特别尤其in spite of尽管isotronic均质的isotrophicsubstitution均匀取代layed structure层状结构Low temperaturemodifications低温变体non-metal非金属octahedral hole八面体空隙olivine minerals橄榄石矿物on the basis of 以…为基础point defects点缺陷quantum mechanics 量子力学shrinkage and densification收缩和致密化solid solution固溶体tetrahedral coordinations四面体配位tetrahedral site 四面体位置Three dimensiona models三维结构模型transmission of light beams透过光束transparency、translucency、opacity透明、半透明、不透明universal acceptance普遍认可vacancy pair空位对Van der Waals forces范德华力vice versa反之亦然。

陶瓷专业英语词汇

[原创] 陶瓷专业英语materials（body composition)stoneware:粗瓷dolomite：白云土terracotta:红土construction/building material :建白材料type of glaze:matt:哑光transparent:透明釉opaque: 不透明釉pigmented:色釉crackled:裂纹釉pearlized :珍珠釉under glazed:釉下彩on-glazed:釉上彩浮雕：relief, embossbisque firing:素烧glost firing:釉烧gild ：镀金decal：贴花trinket box：首饰盒silk screen printing :丝网印素彩瓷plain porcelain高温陶瓷refractory china窑：kilntunnel:隧道窑印花：stamping又收集了一些陶瓷术语陶瓷原料Ceramic Material长石feldspar瓷泥petunse, petuntse, petuntze瓷漆enamel paint, enamel封泥lute高岭土kaolin, china clay硅石,二氧化硅silica, SiO2堇青石cordierite莫来石,红柱石andalusite泥果,坯体clay body 泥釉slip石灰,生石灰,氧化钙lime, calcium oxide, CaO氧化锡tin oxide釉glaze原材料raw material云母mica皂石,块滑石steatite陶瓷类型Types of Pottery碧玉细炻器jasper薄胎瓷thin china彩陶器,釉陶faience陈设瓷,摆设瓷display china瓷porcelain, china (China ‘中国’来自’Chin’’秦’,在英文中’中国’和’瓷’同一单词) 赤陶terracotta, terracotta, red earthenware代尔夫精陶delft德化陶瓷Te-hua porcelain, Dehua pottery高温陶瓷refractory china工业陶瓷industrial ceramics工艺瓷,美术瓷,艺术瓷art porcelain, art and craft china, art pottery, artistic china 骨瓷bone china官瓷mandarin porcelain光瓷lusterware黑色陶器basalt裂变瓷crackled porcelain裂纹瓷crazed china米色陶器creamware青瓷celadon青花瓷bule and white porcelain轻质瓷、轻瓷light china日用瓷household china, table ware软瓷soft porcelain杀菌陶瓷antiseptic pottery绳纹陶器Jomon pottery施釉陶器slipware炻瓷stoneware素彩瓷plain porcelain陶earthenware陶瓷pottery无釉陶、陶瓷素烧坯biscuit, unglazed ware锡釉陶majolica细瓷fine china硬瓷hard porcelain赭色粘土陶器terra sigillata紫砂purple granulated, purple sand, terra-cotta工艺技术Technology凹雕intaglio标记marking玻璃化vitrify车削turning成型forming冲压,冲压花repousse瓷土加工clay processing雕刻carving浮雕relief隔焰窑muffle工艺技术technology硅氧键silicon-oxygen bond技艺technique, craft间断窑intermittent kiln浇铸casting拉毛sgraffito连续窑continuous kiln镂雕、镂空piercing辘轳车jigger泥釉彩饰法trailing碾磨grinding抛光burnishing, polishing破裂chip嵌入inlay切刻incising筛子sieve烧制firing陶瓷科技ceramics陶轮potter’s wheel贴花、嵌花appliqué, decal凸雕，底切，拉底，底部掏槽undercut细裂纹craze性能property压印impressing窑kiln印花stamping釉上彩overglazed color figure釉下彩underglazed color figure预加工pre-processing 粘性，粘滞性viscosity, stiffness 转模片jiggered piece转印transfer print装饰décor, decoration其它Others斑点speck半透明translucence, translucency, translucent不渗透的nonporous不透明的opaque茶叶罐caddy单色的monochrome多色的polychrome高白high white, Gaobai工艺品artware鬼工，鬼爷神工demon’s work, kuei kung景德镇Jingdexhen, Ching-te-chen景泰蓝cloisonné绝缘子insulator考古学archaeology可塑的plastic流变学rheology琉璃瓦glazed tile 模型、模特model模子mould耐热heat-proof配方formula盆栽bonsai漆器lacquer work器皿ware秦始皇陵兵马俑life-size terra-cotta soldiers and horses in Chin tomb青铜器bronze work人类学antropology渗透的porous手印，指印finger mark丝网印刷silk screen printing四面体tetrahedral搪瓷，珐琅enamel陶瓷的ceramic陶瓷专家，陶瓷艺术家ceramist陶工potter瓦tile 碗bowl卫生洁具sanitary ware温度temperature硬度hardness釉工glazier圆块，雕球，瘤knob砖brick爱比克泰德Epicteus（活动于520-500 BC）古希腊陶工兼画家何朝宗He Chaozhong（1522-1620）Chinese ceramist in Ming Dynasty,中国明代陶瓷艺术家韦奇伍德Wedgwood（1730-1795）英国著名陶瓷工匠和制造商希腊古瓮颂’Ode on a Crecian Urn’英国诗人济慈Keats（1795-1821）的名诗，惊叹古希腊陶器彩绘之精美陶瓷专业核心词汇AAbsolute alcohol 无水酒精Abrasive paper 砂纸Agate 玛瑙Ageing 陈腐Agitator 搅拌器Albite 钠长石Alkali 碱Alumina 氧化铝（?/FONT> Amber 琥珀Amethyst 紫水晶Allowance 公差Ammonia 氨Amorphous 无定形的Amphora 双耳瓶Ampoule 安瓿（bu，四声）Analytical balance 分析天平Angle brick 角砖Anhydrous gypsum 无水石膏Anisotropic 各向异性的Anneal 退火Anthemion 无烟煤Apparent porosity 显气孔率Aqua regia 王水Application of glaze 施釉Aerometer 比重计Artesian well自流井Asbestos 石棉Asphalt 沥青Atomizing dryer 喷雾干燥Autogenous ignition 自燃Avogadro’s number 阿伏加德罗常数Azote 氮Azure 天蓝色BBakclite 电木Balancing 配平Ball-float valve 浮球阀Ball mill 球磨机Barite 重晶石（硫酸钡）Basalt 玄武岩Batch bucker 料斗Beaker flask 锥形瓶Bearing 轴承Bending stress 弯曲应力Bentonite 膨润土Biotite 黑云母Biscuit fire 素烧Bitumen 沥青Black lead 石墨Blanket feeder 毯式投料机Blank test 空白试验Blende 闪锌矿Blender 混料器Blinding 失透Blue copperas 胆矾Bogie 窑车Bond angle 键角Bond distance 键长Borax 硼砂Boyle’s law 波义耳定律Brinell hardness 布氏硬度Brittle fracture 脆性断裂Brown coal 褐煤Burette 滴定管Burned lime 生石灰补充陶瓷原料Ceramic Material长石feldspar瓷泥petunse, petuntse, petuntze 瓷漆enamel paint, enamel封泥lute高岭土kaolin, china clay硅石,二氧化硅silica, SiO2堇青石cordierite莫来石,红柱石andalusite泥果,坯体clay body 泥釉slip石灰,生石灰,氧化钙lime, calcium oxide, CaO氧化锡tin oxide釉glaze原材料raw material云母mica皂石,块滑石steatitedolomite：白云土terracotta:红土construction/building material :建白材料陶瓷质地和类型Qualities & Types of PotteryWhite Body 白胎碧玉细炻器jasper薄胎瓷thin china / egg-shell porcelain彩陶器,釉陶faience陈设瓷,摆设瓷display china / ornamental porcelain瓷porcelain, china (China ‘中国’来自’Chin’’秦’,在英文中’中国’和’瓷’同一单词) 赤陶terracotta, terracotta, red earthenware代尔夫精陶delft德化陶瓷Te-hua porcelain, Dehua pottery高温陶瓷refractory china低温陶瓷low-fired porcelain工业陶瓷industrial ceramics工艺瓷,美术瓷,艺术瓷art porcelain, art and craft china, art pottery, artistic china骨瓷bone china古瓷ancient porcelain官瓷mandarin porcelain光瓷lusterware黑色陶器basalt裂变瓷crackled porcelain裂纹瓷crazed china米色陶器creamware青瓷celadon青花瓷blue and white porcelain轻质瓷、轻瓷light china日用瓷household china, table ware软瓷soft porcelain杀菌陶瓷antiseptic pottery绳纹陶器Rope figure pottery施釉陶器slipware粗瓷stoneware素彩瓷plain porcelain陶earthenware陶瓷pottery无釉陶、陶瓷素烧坯biscuit, unglazed ware 锡釉陶majolica细瓷fine china硬瓷hard porcelain赭色粘土陶器terra sigillata紫砂purple granulated, purple sand, terra-cotta 长石瓷feldspar porcelain瓷牙；牙科用瓷dental porcelain瓷质纱网装饰瓷器lace work porcelain电瓷electrical porcelain雕塑瓷；象牙色帕利安瓷statuary porcelain 高铝瓷high-alumina porcelain锆质瓷zircon porcelain滑石瓷steatite porcelain化学瓷chemical porcelain尖晶石瓷spinel porcelain堇青石瓷cordierite porcelain镁橄榄石瓷forsterite porcelain镁质瓷；氧化镁瓷magnesia porcelain莫来石瓷mullite porcelain耐热瓷器refractory porcelain缥瓷faint coloured porcelain青白瓷greenish white porcelain熔块瓷fritted porcelain三成分瓷器triaxial porcelain钛质瓷titania porcelain天然原料制成的瓷器natural porcelain 透辉石瓷diapside porcelain卫生瓷器sanitary porcelain鸭蛋青瓷duck-egg porcelain氧化铝瓷alumina porcelain原始瓷proto-porcelain皂石瓷soapstone porcelain中火度瓷intermediate porcelainIC管瓷IC packages白云石瓷dolomitic保温耐火砖insulating firebrick--釉面transparent:透明釉的opaque: 不透明釉的pigmented:色釉的crackled:裂纹釉的pearlized :珍珠釉的Color Glaze 花釉White Glaze 白釉Red Glaze 红釉Pearl Glaze 珍珠釉翡翠釉;翠绿釉Kingfisher blue glaze 分相釉phase separation glaze粉青釉lavender grey glaze复合釉;混合铀composite glaze钙釉.石灰釉calcareous glaze盖底釉cladding glaze高硅质釉;硅酸质釉siliceous glaze高温釉hard glaze瓜皮绿釉cucumber green glaze光泽釉bright glaze光泽釉glossy glaze孩儿脸(铜红釉)crushed strawberry red 孩儿脸釉strawberry red glaze海参釉trepang glaze海棠红釉begonia red glaze虹彩釉luster glaze弧坑釉crater glaze虎斑釉tiger-skin glaze琥珀釉amber glaze花釉；复色釉fancy glaze滑石釉talc glaze灰釉ash glaze挥发釉vapour glaze火焰红釉flamboyant red glaze鸡皮釉fowl-skin glaze基础釉parent glaze祭红altar red祭红釉sacrificial red glaze祭蓝altar blue祭蓝釉sacrificial blue glaze荠红釉shiny red glaze荞蓝釉deep blue glaze贾丁尼尔釉Jardiniere glaze碱石灰釉alkaline-calcareous碱釉alkaline glaze豇豆红釉cowpea red glaze豇豆红釉haricot red glaze酱釉；棕釉brown glaze结晶釉crystalline glaze金丝黄釉gold filament yellow glaze金星绿釉aventurine green glaze金星釉；砂金釉aventurine glaze金属釉metallic glaze桔皮釉orange-peel glaze钧釉chun glaze可鲁宾釉（红蓝混合釉）Columbine glaze可帕尔塔釉（盖面透明釉）coperta glaze可溶性釉soluble glaze孔雀蓝釉peacock blue glaze孔雀绿釉peacock green glaze辣椒红釉chilli red glaze郎窑红釉Lang yao red glaze郎窑绿釉Lang yao green glaze梨皮釉pear peel glaze锂辉石釉spodumene glaze流纹釉flowing glaze龙皮釉dragon-skin glaze卵青釉egg and spinach glaze罗宾蛋壳釉（乳白青绿色）robin\'s-egg glaze罗金厄姆釉（紫褐色铅釉）Rockingham glaze麻点釉sesame pot glaze梅子青釉plum green glaze美人醉釉beauty\'s flush glaze米黄色釉；奶油色釉cream glaze面釉cover glaze墨地三彩tricolour with china-ink南京黄釉（金黄-棕色）Nankin yellow glaze泥浆釉；易熔粘土釉slip glaze凝固釉（釉浆与明胶的混合物）solidified glaze牛血红釉ox-blood glaze牛血红釉sang-de-boeuf (法语）硼釉boracic glaze葡萄紫釉grap purple glaze铅硼釉lead borate glaze铅釉lead glaze茄皮紫釉aubergine glaze青瓷釉celadon glaze日本风格的铁系花釉；天目釉的日本译名tessha glaze 熔块釉fritted glaze乳白釉opaline glaze乳光釉opalescence glaze乳鼠皮釉mousie skin glaze砂金釉；金星釉gold stone glaze鲨皮釉（有皱纹特征）[日本]shark-shin glaze 鳝皮绿釉green eel-skin glaze鳝皮釉eel-skin glaze鳝鱼黄釉eel yellow鳝鱼青釉eel bluish green glaze蛇皮绿釉green snake-skin glaze蛇皮釉snake-skin glaze生料釉raw glaze生铅釉raw lead glaze失透釉devitrification glaze石灰釉lime glazeCrack Glaze 纹片釉Color Glaze with gold 色釉金彩工艺技术Technology凹雕intaglio标记marking玻璃化vitrify车削turning成型forming冲压,冲压花repousse瓷土加工clay processing雕刻carving浮雕relief隔焰窑muffle工艺技术technology硅氧键silicon-oxygen bond技艺technique, craft间断窑intermittent kiln浇铸casting拉毛sgraffito连续窑continuous kiln镂雕、镂空piercing辘轳车jigger泥釉彩饰法trailing碾磨grinding抛光burnishing, polishing破裂chip嵌入inlay切刻incising筛子sieve烧制firing陶瓷科技ceramics陶轮potter’s wheel贴花、嵌花appliqué, decal凸雕，底切，拉底，底部掏槽undercut细裂纹craze性能property压印impressing窑kiln印花stamping釉上彩overglazed color figure釉下彩underglazed color figure预加工pre-processing 粘性，粘滞性viscosity, stiffness 转模片jiggered piece转印transfer print装饰décor, decorationCut edge切割边缘Scalloped扇形边的陶瓷成品杯mugMeat Plate, Round-edge 荷口汤盘Soup Plate 汤盘Tea cup 茶杯Tea Saucer 茶杯碟Creamer 奶壶Bowl ( Rice pot , Fan-zong) 饭碗Coffee cup / Saucer 咖啡杯碟Tea cup/Saucer茶杯/碟Milk pot 奶壶Salad Bowl 沙拉碗Shaker 筛Tureen 汤窝Duck bowl 鸭碗Rice Bowl 饭碗Rice bowl Flaring 反口碗Rice bowl rope 反口碗Saucer, thick body 厚碟Bowl with cover 盖碗Cylindrical Decor 直身杯碟Tea Pot, Persimmon Shape No. 2 2号筛壶Plate with handle/ heat plate, thick带把手的厚碟4.5”Rice Bowl Flaring 4.5`反口碗Platter 大浅盘Vase 花瓶Fish Bowl 鱼缸Fruit Plate 水果盘Deep Plate深盘Pepper Jar 胡椒瓶pepper cellarSalt Jar 盐瓶Salt sellarSeasoning Dish 格碟Food Mixing Bowl 斗碗Table wares Set 餐具Teacup with saucer 茶杯碟Christmas Toys 小丑Goddess of Mercy 观音Eight Immortals of ancient Figures 八仙God of longevity with lad 寿星God of wealth/Treasure/Longevity 财神Porcelain figuring 人物Mask 面具Oval plate 鱼盘Meat Plate 拼盘Flat Plate 平盘Soup Plate 汤盘Porcelain for daily use 日用陶瓷Porcelain for artistic 工艺陶瓷Porcelain for display 陈列陶瓷the imitation of antique porcelain 仿古陶瓷labels for porcelain 花纸Gold-embossed porcelain 刷金vats 盆Corrugated cardboard boxes 皱硬纸板箱Fleur bouquet 花篮13pec Chinese Tea Pot6pec Decorated Mug setCow mobile 6/s 动态牛Single unicorn独角兽（麒麟）Spoon RestCandle holder 烛座Vase open-work 通花瓶flower pot 花瓶Pair of horse, brown 棕色对马15pes “chaozhou”tea service 15头潮州茶具Sugar pot 糖罐inner box bubble 气泡袋Polly foam polybag 胶袋Lamp stand 灯台Foamed Plastic 泡莫塑料Lampshade 灯罩Candlestick 烛台其它Others斑点speck半透明translucence, translucency, translucent不渗透的nonporous不透明的opaque茶叶罐caddy单色的monochrome多色的polychrome高白high white, Gaobai工艺品artware鬼工，鬼爷神工demon’s work, kuei kung景德镇Jingdexhen, Ching-te-chen景泰蓝cloisonné绝缘子insulator考古学archaeology可塑的plastic流变学rheology琉璃瓦glazed tile 模型、模特model模子mould耐热heat-proof配方formula盆栽bonsai漆器lacquer work器皿ware秦始皇陵兵马俑life-size terra-cotta soldiers and horses in Chin tomb青铜器bronze work人类学antropology渗透的porous手印，指印finger mark丝网印刷silk screen printing四面体tetrahedral搪瓷，珐琅enamel陶瓷的ceramic陶瓷专家，陶瓷艺术家ceramist陶工potter瓦tile 碗bowl卫生洁具sanitary ware温度temperature硬度hardness釉工glazier圆块，雕球，瘤knob砖brick爱比克泰德Epicteus（活动于520-500 BC）古希腊陶工兼画家何朝宗He Chaozhong（1522-1620）Chinese ceramist in Ming Dynasty,中国明代陶瓷艺术家韦奇伍德Wedgwood（1730-1795）英国著名陶瓷工匠和制造商希腊古瓮颂’Ode on a Crecian Urn’英国诗人济慈Keats（1795-1821）的名诗，惊叹古希腊陶器彩绘之精美在线英汉－汉英科技大词典中含有“瓷”的词汇瓷porcelain; china瓷把手porcelain handle瓷板vitrolite瓷保险丝盒porcelain fuse box瓷杯porcelain cup瓷插入式熔丝porcelain plug fuse瓷插头porcelain plug瓷插座porcelain receptacle瓷茶壶china tea-pot瓷衬ceramic liner; porcelain lining瓷衬底印刷电路ceracircuit瓷衬球磨机porcelain mill瓷充填porcelain filling瓷灯头porcelain lamp holder瓷灯罩porcelain lamp shade瓷电容器porcelain condenser瓷雕porcelain carving; sculpture porcelain瓷断流器porcelain cut-out瓷断面porcelain fracture瓷防水拉线开关porcelain water proof cord switch瓷粉充填器porcelain plugger瓷粉调板porcelain slab瓷粉调刀porcelain cement spatula瓷管porcelain tube; porcelain pipe瓷管道porcelain duct瓷管型油断路器porcelain type oil circuit breaker瓷罐porcelain jar; porcelain pot瓷辊porcelain roller瓷过滤棒porcelain filtering stick瓷过滤漏斗porcelain filtering funnel瓷合金刀具porcelain alloy cutter瓷横担绝缘子porcelain-arm insulator瓷花金属板plaque瓷环过滤器porcelain ring filtre瓷火花塞porcelain spark plug; stone plug瓷基体porcelain basal body瓷夹cleat insulator; porcelain clip瓷夹布线cleat wiring瓷件粘接剂binder for porcelain paste瓷胶porcelain cement瓷胶结砂轮vitrified grinding stone瓷接线盒porcelain covered connector瓷进线套管porcelain lead-in bushing瓷绝缘porcelain insulating; ceramic type insulation瓷绝缘体porcelain insulator瓷绝缘子式断流器porcelain clad type circuit breaker 瓷绝缘子线轴porcelain spool insulator瓷壳保险丝porcelain fuse; porcelain cartridge fuse瓷壳接线盒porcelain connector瓷蓝ceramic cobalt blue; cobalt blue瓷料porcelain瓷漏斗Buchner filter; Buchner funnel; porcelain funnel瓷螺丝灯头porcelain screw holder瓷螺丝灯座porcelain screw socket瓷面porcelain facing瓷面具porcelain mask瓷面纸porcelain paper瓷皿porcelain dish瓷模ceramic former瓷磨porcelain mill瓷磨球porcelain media瓷钮porcelain knob瓷盘电阻器porcelain disc resistor瓷盘绝缘子porcelain cup瓷配件porcelain accessory; porcelain fitting瓷硼钙石bakerite瓷片ceramic chip瓷平底漏斗Buchner filter瓷瓶porcelain insulator瓷瓶穿钉insulator bolt瓷瓶球磨机pot mill瓷瓶式电流互感器porcelain clad current transformer瓷瓶式断路器porcelain clad circuit breaker瓷漆enamel lacquer; enamel paint; enamel varnish瓷漆刷enamel brush瓷器china; chinaware; porcelain; porcelain ware; hard paste porcelain; true porcelain在线英汉－汉英科技大词典中含有“瓷”的词汇(2)瓷器餐具porcelain dinner-ware瓷器残片fragment of pottery瓷器店china shop瓷器结合剂porcelain cement瓷器制造日期印记date-letter; date mark瓷器装饰decoration of porcelain瓷嵌体porcelain inlay瓷球磨机porcelain ball mill瓷裙porcelain patticoat瓷燃烧管porcelain combustion tube瓷燃烧舟porcelain combustion boat瓷绕线管绝缘子porcelain bobbin insulator瓷三角架porcelein triangle瓷色料porcelain color瓷砂石porcelain sandstone瓷勺皿porcelain casserole瓷石china stone; porcelain stone瓷室porcelain chamber瓷首饰porcelain jewellery瓷栓porcelain knob瓷胎porcelain body瓷胎画珐琅enamel painting on porcelain瓷胎竹编bamboo over porcelain瓷套冠porcelain jacket crown瓷套管porcelain bushing shell; porcelain through insulator; porcelain sleeve; porcelain bushing 瓷套管绝缘子porcelain through insulator瓷天线绝缘子porcelain antenna insulator瓷贴花印刷油墨ceramic ink瓷头螺丝porcelain head screw瓷头铜螺丝pottery topped brass screw瓷土porcelain clay; china-clay; kaoline瓷土粉china clay in powder瓷土熟料china-clay chamotte瓷托porcelain plate瓷托托牙porcelain base denture瓷外罩porcelain housing瓷弯管porcelain elbow tube瓷相学ceramography瓷芯porcelain core瓷旋流器porcelain cyclone瓷牙porcelain teeth; mineral teeth瓷牙的odontoceramics瓷牙面facing porcelain瓷牙学ceramodontia瓷牙制作术odontoceramotechny瓷研钵porcelain mortar瓷眼porcelain eye瓷样的porcellaneous; porcellanous瓷仪器porcelain utensil瓷用色料porcelain colour瓷釉ceramic glaze; porcelain glaze瓷釉剥落shivering瓷釉电容器glaze capacitor瓷圆柱体porcelain cylinder瓷渣pitchers瓷罩式变流器porcelain clad type current transformer瓷罩式电流互感器porcelain clad current transformer瓷罩式断流器porcelain clad circuit breaker瓷枕porcelain pillow瓷支持绝缘子porcelain support insulator瓷制多孔滤筒bougie瓷制门把手door-handle of porcelain瓷制球磨罐jar mill瓷制珍珠porcelain pearl瓷质porcelain瓷质胶粉porcelain cement瓷质绝缘子porcelain insulator瓷质卫生器vitreous-china sanitary ware瓷质悬挂隔电子suspension porcelain瓷珠porcelain bead; knob insulator瓷柱porcelain knob; knob insulator瓷柱布线knob wiring柱瓷管布线knob-and-tube wiring瓷砖ceramic tile; enamel tile; porcelain tile; tile 瓷砖锤tile hammer瓷砖刀tile cutter瓷砖机wall tile press瓷砖胶glue for tile瓷砖嵌镶tile fillet瓷砖清洗剂tile cleaner瓷座放大镜magnifying mirror with ceramic foot。

第五次课-《材料科学导论》第02章-材料的结构基础02-晶体学与晶体化学(原子规则排列)-2015-骆军

crystobalite +L
mullite
+L
mullite
alumina
+
1600 1400 0
+ crystobalite
mullite
20
40
60
80
100
alumina
crystobalite
Composition (wt% alumina)
图2.10.8 石英-氧化铝中间相相图
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3. 固溶体的结构（Structure of solid solution）
虽然固溶体仍保持溶剂的晶格类型，但与纯溶剂组元的晶体结构相比，其结构还是发生了变化。（1）晶格畸变
由于溶质与溶剂的原子半径不同，因而在溶质原子附近的局部范围内形成一弹性应力场，造成晶格畸变(图2.10.3)。晶格畸变程度可通过溶剂晶格常数的变化反映出来(图2.10.4)。
27
图2.10.11 拓扑密排相中的配位多面体
28
表2.10-1 固溶体与中间相的比较
类型固溶体中间相
单相固溶体位于相图相图中的位置位于相图中部两侧，紧挨纯组元晶体结构特点结构与溶剂相同成分特点原子分布特点性能特点标记符号
2015/4/24
与形成中间相的元素均不同符合特定比例(或在成分不符合特定比例，可连续变化这比例附近连续变化)
20
图2.10.9
AB型（NaCl）正常价化合物的晶体结构
21
2. 电子化合物
（elctrides)
电子化合物是指按照一定价电子浓度的比值组成一定晶格类型的化合物。电子浓度是指化合物中每个原子平均所占有的价电子数（e/a）表示。当价电子浓度为3/2时，电子化合物具有体心立方晶格；价电子浓度达到7/4时，电子化合物具有密排六方晶格。常见的电子化合物有：Fe-Al、Ni-Al等。电子化合物的熔点和硬度都很高，但塑性较差，一般是有色金属的强化相。

耐火材料中英文对照

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英文文献

The role of fine alumina and mullite particles on the thermomechanical behaviour of alumina–mullite refractory materialsCemail Aksel,Department of Ceramic Engineering, Anadolu Univers ity, Iki Eylül Campus, 26470, Eskisehir, TurkeyReceived 29 January 2002;revised 7 May 2002;accepted 23 May 2002. ;Available online 16 June 2002.AbstractFine grain alumina and mullite particles ( 5 μm) were incorporated into slip-cast alumina–mullite refractories in order to investigate their effects on the microstructure, mechanical properties and thermal shock behaviour of the refractories. The incorporation of alumina particles improved both densification and mechanical properties markedly, compared to that of mullite particles, which resulted in high porosity and low mechanical properties. The crack initiation resistance increased with the addition of alumina, which is supported by R parameter. Therefore, more fracture surface energy was required to connect the cracks for propagation, associated with a high value of fracture toughness. The strength and Young's modulus were also improved with increasing quench temperature, leading to a high thermal shock resistance.Author Keywords: Alumina; Mullite; Slip casting; Refractory; Mechanical properties; Thermal shockArticle Outline1. Introduction2. Experimental3. Results and discussion4. ConclusionsReferences1. IntroductionMullite has been studied for structural ceramic applications, because of its high thermal shock resistance, high resistance to creep and chemical corrosion [1, 2, 3 and 4]. However, the potential capability of mullite ceramics cannot be easily exploited for high-temperature structural ceramics because of difficulties in attaining complete densification without additives and the formation of a glassy phase at grain boundaries [5 and 6]. Alumina could be a good candidate to improve the mechanical properties of mullite-based materials, because the formation of a glassy phase may be controlled by alloying alumina into mullite ceramics [3, 5 and 6]. Relative density, hardness and fracture toughness of mullite solid solution sintered bodies are fairly low but increase with increasing alumina concentration [6]. It has been reported that a high loading of fine alumina particles (>10 vol.%) results in a dramatic reduction in the compaction shrinkage, leading to an increase in the compacted density with increasing volume fraction of particles [7]. Furthermore, it appears that the presence of moderate quantities of both mullite and alumina in the refractory materials provides highest thermal shock resistance [8]. This paper explains how mechanical properties and thermal shock behaviour of slip-cast refractory materials are affected by the addition of fine alumina and mullite particles.2. ExperimentalIn the present work, two different batches were prepared by slip casting. The first batch consists of coarse (−2+1 mm) and medium (−1000+20 μm) alumina–mullite particles, including 5 μm fine alumina particles, designated as sample A. The second one is composed of the same range of coarse and medium alumina–mullite as well as fine mullite particles of5 μm, nominated as sample B. The phase compositions of the refractories by vol.% are given as follows: (a) 75% alumina, 20% mullite, 5% clay for sample A, and (b) 55% alumina, 40% mullite, 5% clay for sample B, wherethe chemical composition of the clay used is 48.3% SiO2, 37.6% Al2O3, 1.1%K 2O, 0.5% Fe2O3, 0.4% Na2O, 0.1% CaO and 12% combined water. Slips were takenand cast into bars of 100×10×15 mm3, dried at 110 °C for overnight and fired at 1600 °C for 2 h. Heating and cooling rates were 300 and 200 °C/h, respectively. Bulk density and apparent porosity were measured using the standard water immersion method [9]. Mechanical and thermal shock measurements were then carried out by water quenching the representative bars from room temperature to 1200 °C, at 300 °C intervals. Those bars were tested using a Mayes, SMT50 tensile testing machine in three-point bend configuration with a support roller span of 90 mm, at a crosshead speed of 1 mm min−1. Five specimens were tested to obtain a mean value for each quench temperature. Strength (σ) [10] and Young's Modulus (E) [11] values were determined using the standard equations:(1)E=L3m/(4WD3)where L is the length of support span, W is the specimen width, D is the specimen thickness, and m is the slope of the tangent of the initial straight-line portion of the load–deflection curve. Fracture toughness was measured using the single edge notched beam (SENB) technique [12, 13and 14], with a 1-mm-thick diamond blade producing a notch-to-depth ratio of 0.25. Fracture toughness (K1c) was calculated from the maximum load using Eq. (2):(2)where P is the load at fracture, c is the notch depth, γi is a measure of resistance to the initiation of crack propagation, C is the critical crack length and Y=[A0+A1(c/D)+A2(c/D)2+A3(c/D)3+A4(c/D)4]. Y is a dimensionless number, which is dependent on the geometry of the loading and the crack configuration with L/D≈8, A0=+1.96, A1=−2.75, A2=+13.66, A3=−23.98, A4=+25.22 [12, 13and 14]. The R parameter can be expressed as the difficulty of crack initiation and predicts the resistance of a material to fracture initiation by the thermal stresses. It has been calculated by the formula, {R=[σ(1−ν)]/(Eα)} [15], where ν is Poisson's ratio and α is the mean thermal expansion coefficient of the refractory. αvalues, depending on the Young's modulus, volume fraction and coefficient of thermal expansion of each component, were calculated from the equation derived by Turner [16]. Secondary and back scatter electron images were then made to examine microstructure and fracture surfaces using a CamScan 4 scanning electron microscope.3. Results and discussionThe bulk density values of the representative samples A and B are 2.9 and 2.6 g cm−3, respectively. Chemical compositions of these samples are alsogiven as follows: 93.4% Al2O3, 6.6% SiO2(A); and 87.5% Al2O3, 12.5% SiO2(B). Fig. 1a shows a general microstructure of alumina–mullite refractory materials with a map distribution of each element. Pores in the refractory material are largely concentrated within the coarse mullite grains, but also some are observable within the alumina grains(Fig. 1a). A possible cause for the formation of such large voids is the differential sintering of the mullite grains with respect to the alumina grains due to the density differences. The incorporation of fine alumina powder in the sintered sample A provides more improvement in both bulk density and strength values than that of sample B through a reduction in porosity by filling of interparticle voids and by faster sintering. Fig. 1b illustrates a high packing of fine alumina particles amongst the coarser grains in sample A. On the contrary, the pores in the sample B are mainly located between the interface of alumina and mullite grains, and thus a relatively weak interface was observed between the shorter length of needlelike mullite whiskers (Fig. 1c).Full-size image (84K)Fig. 1. Map distribution of alumina–mullite refractory materials (A: Alumina, M: Mullite, P: Porosity): (a) sample A, (b) sample A, (c) sample B.It is well known that the significant difference in the coefficients of thermal expansion, which are 5.3×10−6K−1for mullite and 7.8×10−6K−1 for alumina [17], leads to a marked crack development between alumina and mullite grains, due to the large tensile hoop stresses during cooling from fabrication temperature of 1600 °C. As can be seen from Fig. 1a, crack formation has been observed around both alumina and mullite grains, where when a crack reached to the coarse mullite grains, crack propagation decreased steadily. After the interlinked crack between the coarsergrains came across to the pores, crack blunting occurred and thus the resistance to further crack extension was increased, as shown in Fig. 1a. The apparent porosity values and mechanical properties of samples A and B are given in Table 1. The incorporation of fine mullite particles, resulting in a low packing density amongst the coarser grains and thus a high porosity value, led to relatively low mechanical properties. However, the addition of fine alumina particles increased both strength (≥70%) and Young's modulus (50%) significantly. This improvement indicates that critical crack length in sample A was limited, and slow crack propagation occurred, although the calculated critical crack lengths of these samples have the same values. This is because of both the sintering morphology and the distribution of the pores in the microstructure, where a strong interface was observed, and cracks were prevented from propagating around the pores. There was also a significant improvement on K1c value of sample A, risen by a factor of 1.6. The γi values confirmed that the amount of energy required to initiate crack propagation for sample A is much higher than B, and thus more energy was required to connect the cracks for propagation. As predicted and supported by the thermal stress resistance parameter R, material containing fine alumina particles appeared to have higher resistance to crack initiation caused by the thermal stresses, regarding to that of sample B (Table 1).Table 1. Apparent porosity values and mechanical properties of samples A and BFull-size table (5K)Fig. 2a and b showed that there was mostly intergranular, with some transgranular fracture in sample B, where shorter length of needlelikemullite whiskers have been observed. However, Fig. 3a and b showed that both intergranular, and a significant amount of larger transgranular cracks occurred through the alumina grains, where a strong bonding was observed between the interface of finer and coarser grains. The major change in the fracture path from intergranular to a marked transgranular fracture by the addition of fine alumina particles is associated with the improvement in the fracture surface energy and fracture toughness values of sample A.Full-size image (67K)Fig. 2. Fracture surface of sample B, quenched from 1200 °C (t: transgranular, i: intergranular): (a) scale bar: 20 μm (magnification: ×1250), (b) scale bar: 20 μm (magnification: ×1750).Full-size image (78K)Fig. 3. Fracture surface of sample A, quenched from 1200 °C (A: Alumina, t: transgranular, i: intergranular): (a) scale bar: 50 μm (magnification: ×770), (b) scale bar: 20 μm (magnification: ×1370).Fig. 4a and b showed a decreasing trend in strength and Young's modulus values with increasing quench temperature, where sample A had higher values of both strength and modulus from room temperature to 1200 °C. It has been seen that both strength (>40%) and modulus (>30%) values were increased markedly by the addition of fine alumina particles. This improvement indicated that retained strength of sample A is higher than B, with a limited crack size and slow crack propagation, which is associated with higher mechanical properties.Full-size image (8K)Fig. 4. Mechanical properties of alumina–mullite refractory materials as a function of quench temperature: (a) flexural strength, (b) Young's modulus.4. ConclusionsAlthough the addition of fine mullite resulted in a high porosity and low mechanical properties, both densification and mechanical properties increased with the incorporation of fine alumina particles significantly through improvement of interparticle bond strength amongst the coarser grains. Furthermore, the crack initiation resistance was also increased with higher alumina content, as supported by the R parameter. After thermal shock testing, there was a significant change in fracture path from mainly intergranular to a significant amount of transgranular fracture, due to the presence of fine alumina. This explains that moreenergy is required to initiate crack propagation, associated with the high values of γand K1c. Thermal shock tests confirmed that the addition ofifine alumina rather than fine mullite particles into an alumina–mullite slip-cast refractory material improved both strength and modulus values with increasing quench temperature, indicating better thermal shock resistance.References1. H. Schneider, K. Okada and J. Pask , Mullite and Mullite Ceramics. , Wiley, UK (1994).2. J.H. Chesters , Refractories: Production and Properties. , The Iron and Steel Institute, London (1973).3. K.S. Mazdiyasni and L.M. Brown , Synthesis and mechanical properties of stoichiometric aluminium silicate (mullite). J. Am. Ceram. Soc.55 (1972), p. 548. Full Text via CrossRef4. N. Tamari, I. Kondoh, T. Tanaka and H. Katsuki , Mechanical properties of alumina–mullite composites. J. Ceram. Soc. Jpn.101 (1993), p. 721. View Record in Scopus | Cited By in Scopus (3)5. M.D. Sacks and J.A. Pask , Sintering of mullite-containing materials: 1. Effect of composition. J. Am. Ceram. Soc.65 (1982), p. 65. Full Text via CrossRef | View Record in Scopus | Cited By in Scopus (30)6. T. Sato, M. Shizuka and M. Shimada , Sintering and characterisation of mullite alumina composites. Ceram. Int.12 (1986), p. 61. Abstract | PDF (405 K) | MathSciNet | View Record in Scopus | Cited By in Scopus (8)7. J. Wang, M.R. Piramoon, C.B. Ponton and P.M. Marquis , A study in short alumina-fiber-reinforced mullite composites. Trans. J. Br. Ceram. Soc.90 (1991), p. 105. View Record in Scopus | Cited By in Scopus (6)8. N.C. Biswas and S.P. Chaudhuri , Comparative study of zirconia–mullite and alumina–zirconia composites. Bull. Mater. Sci.22 1 (1999), p. 37. Full Text via CrossRef | View Record in Scopus | Cited By in Scopus (7) 9. British Standard Testing of Engineering Ceramics, BS 7134 Section 1.2, 1989.10. ASTM C1161-90, Standard Test Method for flexural strength of advanced ceramics at ambient temperature. In: Annual Book of ASTM Standards, 15.01 (1991), p. 327.11. ASTM D790M-86, Standard Test Methods for flexural properties of unreinforced and reinforced plastics and electrical insulating materials. In: Annual Book of ASTM Standards, 08.01 (1988), p. 290.12. D.R. Larson, J.A. Coppola and D.P.H. Hasselman , Fracture toughnessand spalling behaviour of high-Al2O3refractories. J. Am. Ceram. Soc.57(1974), p. 417. Full Text via CrossRef | View Record in Scopus | Cited By in Scopus (32)13. W.F. Brown and J.E. Srawley , Plane strain crack toughness testing of high strength metallic materials. ASTM Spec. Tech. Publ.410 (1967), p. 1.14. ASTM D5045-91, Standard Test Methods for plane-strain fracture toughness and strain energy release rate of plastic materials. In: Annual Book of ASTM Standards, 08.03 (1991), p. 728.15. D.P.H. Hasselman , Thermal stress resistance parameters for brittle refractory ceramics: a compendium. Am. Ceram. Soc. Bull.49 12 (1970), p. 1033.16. P.S. Turner , Thermal-expansion stresses in reinforced plastics. J. Res. Natl. Bur. Stand.37 4 (1946), p. 239.17. S.J. Burnett , Properties of Refractory Materials. , UKAEA Research Group Report, Harwell (1969).。

球磨法制备超细氧化铝外文原文及译文

北京联合大学毕业设计（论文）外文原文及译文题目：球磨法制备超细氧化铝的工艺优化分析专业：材料科学与工程（检测与质量管理工程）指导教师：王训伟学院：机器人学院学号： 2014090371026班级：材料1401B 姓名：马子悦一、外文原文In-situ synthesis and characterization of nacho-structured NiAl-Al2O3 composite during high energy ball millingIn this work, synthesis of NiAl–Al2O3nanocomposite powders via the mechanosynthesis route and by using Ni, NiO and Al is investigated. Ignition of the reaction inside the ball-mill vial happens after 110 min; NiO is totally finished and NiAl and Al2O3as product phases are formed. After 10 h of ball milling, raw materials are totally used in the reaction and only product phases exist in the vial. By continuing the ball milling process to 60 h, better mixing of the synthesized phases and decrement in their crystallite sizes plus particle size are observed. Crystal-lite sizes of the product phases are in the nanometer range in all ball milling times. Crystallite sizes of NiAl and Al2O3 after 10 h are around 11 nm and 19 nm respectively, and these are reduced to around 8 nm for both phases after 60 h of ball milling.1. IntroductionMechanical energy has been proved capable of providing the energy needed to start a reaction as well as thermal energy. It is reported that high energy ball milling provides energy for activation of a chemical reaction and also for full occurrence of a chemical reaction during milling. The first is defined as mechanical activation and the second is called as mechanochemical synthesis or reactive milling. In such a process, the intensely high mechanical impacts result in repeated welding and fracturing of powder particles. Therefore, the contact area between the reactant powder particles increases and fresh surfaces come into contact in each collision in the ball mill. This matter leads the reaction to proceed without the necessity for diffusion through the product layer. Intense plastic deformation, increment in strain energy and thus high defect density leads to activation/occurrence of a chemical reaction. High defect density accelerates the diffusion process during mechanical milling Hence, it is expected that kinetic and thermodynamic behavior of the aforementioned reactions may be very different from those reactions initiated by thermal energy.These reactions might undergo two different kinetics based on milling conditions:1- Gradual reactions or progressive type reactions in which materials transform in a gradual manner and in very small volumes during each collision.2- A self-propagating combustion reaction when the reaction enthalpy is high enough In the latter type, there is a critical milling time in which the combus-tion reaction initiates. In the beginning, the temperature of the milling vial shows a slow increase over time. Then, it increases abruptly which implies that the reaction ignition has occurred (ignition time = tig). After this time, temperature decreases slowly. Measurement of tig enables definition of the structural and chemical evolution during combustion reaction In the reactive milling method, there are varied and numerous benefits and these make this process economical. These benefits include the ability to produce solid solutions, different volume fractions of different types of submicron reinforcements as well as nascent clean interfaces . Researchers have focused on intermetallics such as nickel aluminides. They utilize advantages of this process to improve ductility of such materials by microstructure refinements and to enhance their creep resistance by the reinforced particulates . These goals have led researchers to concentrate on nanostructured composites.Alumina has been considered as a promising reinforcement for NiAl by researchers . That is because of its thermodynamic compati-bility with NiAl and also because its thermal expansion coefficient is relatively close to that of NiAl. And furthermore, alumina has desirable properties such as low density, high specific strength, high modulus and good oxidation resistance . Research has shown that NiAl–Al2O3composite can be utilized for high temperature applica-tions because of improved oxidation and creep resistance and thus is a good candidate for coating purposes .Researchers have reported the formation of NiAl–Al2O3composites by mechanical milling powder mixtures of Ni, Al and Al2O3, NiO and Al and Ni2(OH)2CO3·3H2O and Al . Anvari et al. devel-oped NiAl-Al2O3composite by high-energy ball milling of Ni, NiO and Al in a planetary ball mill without any process controlling agent (PCA). Udhayabanu et al. carried out reactive milling of Ni, NiO and Al powder mixture in toluene medium. Then they heated the as-milled powder. Even after the 20 h of ball milling, reaction between raw materials was still incomplete; significant amounts of Ni and NiO remained in the mixture and Al was totally consumed. NiAl-Al2O3com-posite was gained after the mixture was heated to 1200 °C. consolidated the as-milled powder containing NiAl, amorphous Al2O3 and unreacted NiO by spark plasma sintering and finally gained NiAl-Al2O3 composite. Because the above-mentionedre-searchers have conducted their milling in conditions different from the present work, the phase formation and reaction mechanism are ex-pected to be different in this study.The present research aims to develop NiAl-Al2O3 nano-structured composite by reactive milling of Ni, NiO and Al as raw materials and by completing the reaction totally in the ball mill vial. Ignition time of the reaction is measured. Influence of ball milling time is assessed on phase evolution, crystallite size and lattice strain of phases, microstruc-ture and chemical composition of different phases in the as-milled powder mixture and powder particle size (for milling times higher than 10 h).2. Experimental2.1. CharacterizationBall milling was carried out in a Retsch planetary ball mill. The constant ball milling conditions are mentioned in Table 1 and only ball milling time differs for different samples.Stearic acid (1 wt% of the powder mixture) was used as process controlling agent (PCA) to avoid oxidation and excessive cold welding of powders to vial and balls.Temperature variations of outer wall of the ball mill vial were studied as an indirect indicator of the temperature changes inside the vial. For this purpose, two RoHS Thermochromic Liquid Crystal Reversible Temperature Indicating Strips were attached on the outer wall of the vial.able 1The ball milling conditions.Vial Stainless steel vial with volume of 250 mlBalls WC balls with approximate weight of 8.4 g/ball Ball milling speed (rpm) 280ball-to-powder ratio 10/1Process controlling agent (PCA) Stearic acid (1 wt% of the powder mixture)Raw materials powder mixture Total weight: 5.85 (Ni:2.63, NiO:1.25, Al:1.97) weight (g)Ball milling time (h) 1, 2, 3, 4.5, 6, 10, 15, 20, 40, 60Ball milling atmosphere AirIn these strips, temperature ranges are seen in dark brown squares and color change of one square from brown to green indicates temperature change to that range. The temperature range under this survey was 60 °C. This assessment was intended to determine the ignition time of the exothermic reaction inside the ball mill vial.The morphologies, microstructures and chemical compositions were studied using a LEO 1530 scanning electron microscope (SEM) with a Gemini column and EDS detector from Oxford Instruments of type 50 mm2 XMax Silicon Drift Detector (SDD). To prepare samples for powder morphology investigations, powder particles were dispersed in ethanol. To study the microstructure of powder particles and phase investigations, a powder particle cross-section sample was prepared by conventional crystallography techniques. Phase evolution of samples was analyzed by PANalytical X'pert PRO XRD, using Cu Kα radiation with λ = 0.15406 nm. Using diffraction data of ball milled powder samples, the amount, crystallite size and lattice strain of existing phases were obtained from Rietveld refinement analysis by ap-plying a pseudo-Voigt function and by X'Pert HighScore Plus software. The particle size measurement of raw materials and powder samples was performed using a SHIMADZU SALD-2101 apparatus.2.2. MaterialsRaw-materials used in this work were Al (Merck-art no.1056), Ni (Merck-art no.12277) and NiO (Merck-art no.6723) powders. Particle size measurement showed that d50 of Al, Ni and NiO was 68 ± 0.5, 5 ± 0.1 and 9 ± 0.3 μm, respectively. Fig. 1 shows scanning electro n microscopy micrographs of the starting powder particles. The aluminum and nickel powder particles had spherical morphologies while nickel oxide powder particles showed an irregular shape.3. Results and discussion3.1. Phase analysesFig. 2 presents XRD patterns of the mixed powder mixture (0 h of high energy ball milling) and ball milled ones after different milling times. In the mixed and 1 h ball milled powder, peaks of Ni, Al and NiO are observed, so these were just reactant phases and none of reac-tion products were formed in powders. While increasing milling time to 2 h led to the disappearance of NiO peaks, NiAl and Al2O3peaks began to appear, which are formed “in-situ” during themilling process. It should be mentioned that in the binary phase diagram of NiO-Al2O3, there is no complex oxide composition other than NiAl2O4 and peaks of this phase were not observed in any of patterns in Fig. 2. Thus, know-ing that ball milling is not an equilibrium process, it seems that NiO is not consumed in any complex oxide in Ni-Al-O system and that Al2O3is the product of reduction reaction of NiO and Al. Therefore, within 1 and 2 h of ball milling, the exothermic reaction between starting pow-ders has taken place. With further milling, the intensity of Al and Ni peaks decreases. After 4.5 h of ball milling, Al peaks are not observable in XRD pattern while Ni peaks can still be seen. This has been reported in other references too. Peak broadening is a common phenomenon which is observed as a result of grain refinement and increase in lattice strain . Because of the high ductility of aluminum in comparison with Ni, it experiences rapid and severe deformation. Therefore, the broadening and disappearing of peaks happens faster for aluminum, while nickel diffraction lines are observable over longer times because of its greater hardness and so undergoes less deformation than aluminum. Abbasi et al. [28] observed the same phenomenon in milling of nickel and aluminum powder mixture. If solid solution of Ni(Al) forms during ball milling, because of radius difference between nickel and aluminum, it is expected that expansion in the nickel lattice happens and thus Ni diffraction peaks should be detected at lower 2θ. But since such a peak shift is not seen in phase analysis results, it can be concluded that Ni(Al) solid solution has not been formed during milling. Studies on heating of the nickel aluminum binary system dem-onstrated that the phase formation sequence is NiAl3/Ni2Al3/NiAl/Ni3Al. However, in XRD patterns of ball milled powder, the only formed nickel aluminide phase is NiAl. In reactive milling, mechanical energy is employed to initiate chemical reactions. Intense mechanical pressure in ball milling causes formation of new surfaces, plastic defor-mation and increase in lattice defect density which makes the kinetics and thermodynamics of reactions in ball milling very different from that of thermally activated reactions. Ball milling leads to particle refine-ment and thus reduction in diffusion distances (by reduction in interla-mellar spacing) which lowers the activation energy for diffusion significantly . During milling, the system temperature is more or less constant and around room temperature. However, intense pressure is imparted drastically on powder particles by balls and the achieve uniform distribution of phases in powder particles. In addition to refinement of crystallites and increment in lattice strain, the lower in-tensity of alumina peaks than that with NiAl ones is because of the lower atomic scattering factor of Al2O3 than NiAl. The atomic scattering factor is proportional to atomic number in a phase. The higher the atomic scattering factor of a phase is, the sharper and more distinguishable the peaks are. Thematerial with lower atomic scattering factor shows peaks with lower intensity and, in longer ball milling times, these peaks diminish. It seems that the atomic scattering factor is very impor-tant in affecting peak intensity . In other works, researchers faced the same problem in defining alumina peaks properly in XRD patterns. In a work by Marashi et al. in WC-Al2O3 system, despite the existence of 34.24 wt% of alumina, alumina peaks could not be distinguished even after 2 h of milling. Hosseinpour et al. in Mo-ZnS system in presence of alumina, could not observe alumina peaks in XRD patterns, although they had 26 wt% of alumina in their system.4. ConclusionIn the present work,the development of NiAl-Al2O3 nanostructured composite via a mechanochemical reaction in a powder mixtures of 13Al+8Ni+3NiO was investigated.The following conclusions can be derived from the experiments:•In early stages of ball milling,that is up to 110 min,powder particles experienced mixing besides intense plastic deformation.This process caused layer-by-layer cold welding of phases on new and fresh sur-faces of each other.•For between 100 and 110 min of ball milling,a mechanically induced self-propagating reaction(MSR)took place.Ignition of the exothermicreaction caused a sudden temperature increase,NiO was consumed totally,the Ni and Al amount decreased and NiAl and Al2O3as the re-action products phases were formed.V olume changes resulting from formation reaction of nickel aluminides were calculated.They showed that NiAl is the only phase which is expected to form in the system be-cause of its highest volume reduction.•By continuation of ball milling to 10 h, raw materials were no longer detected and just NiAl and Al2O3existed in the ball mill vial. Increasing ball milling time to 60 h resulted in better intermixing of synthesizedphases, decrement in their crystallite sizes and increment in lattice strains, as well as decrement in powder particles size. Crystallite sizes of NiAl and Al2O3 were in nano-metric scale in all ball milling times, and after 60 h they both reached around 8 nm.二、译文纳米结构NiAl-Al2O3复合材料在高能球磨过程中的合成与表征在这项工作中，研究了使用Ni，NiO和Al合成NiAl-Al2O3纳米复合粉末的机械合成路线。

相场法对含第二相颗粒的纳米结构AZ31镁合金晶粒生长的模拟研究(英文)

ABSTRACT Phase field models were established to simulate the grain growth of a nanostructured AZ31 magnesium alloy, which contain spherical particles of differing sizes and volume fractions, under realistic spatial and temporal scales. The effect of the second phase particles on the nanostructure evolution was studied. The simulated results were compared with those of the conventional microstructured alloy. The expression of the local free energy density was improved by adding a second phase particle term. The right input parameters were selected for proper physical meaning. It was shown that the rules that govern the pinning effect of the second phase particles during the grain growth were different for the nanostructure and microstructure. There was a critical particle size value that affected the grain growth within the nanostructure. If the particle size was lower than the critical value, the pinning effect on grain growth increased with decreasing particle size. When the particle size was greater than the critical value, the particles had almost no pinning effect. However, in the conventional microstructured material, the larger particle size resulted in an enhanced pinning effect during grain growth for particle sizes smaller than 1 µm. The effect was reversed when the particle size was larger than the critical value. For the nanostructure, the critical value was 200 nm when the particle content was 10 v.%, and the critical value decreased when the content increased. When the particle size was 30 nm, the particle pinning effect on the grain growth increased for increasing particle content. Keywords: phase field models, second phase particles, grain growth, magnesium alloy

气相法氧化铝英文

气相法氧化铝英文Gas-phase synthesis of alumina, also known as aluminum oxide, is a process that involves the oxidation of aluminum-containing precursors in a gaseous environment. This method offers several advantages over traditional solid-state synthesis techniques, including better control over particle size and morphology, as well as the ability to produce high-purity materials.One of the key advantages of gas-phase synthesis is the ability to precisely control the stoichiometry of the reaction. By adjusting the composition of the gas mixture, researchers can tailor the properties of the resulting alumina particles, such as their crystallinity and phase composition. This level of control is crucial for applications that require materials with specific characteristics, such as catalysis or electronic devices.Another benefit of gas-phase synthesis is the ability to produce nanoparticles with a narrow size distribution. In traditional solid-state methods, it can be challenging to achieve uniform particle sizes, which can lead to variations in material properties. By contrast, gas-phase synthesis allows for the precise control of nucleation and growth processes, resulting in particles with consistent sizes and shapes.Gas-phase synthesis also offers the advantage of producing high-purity alumina materials. Since the reaction takes place in a controlled gaseous environment, there is minimal contamination from impurities that are often present in solid-state precursors. This is particularly important for applications where purity is critical, such as in the production of advanced ceramics or catalysts.In addition to these advantages, gas-phase synthesis of alumina is a scalable and cost-effective process. The equipment required for gas-phase synthesis is relatively simple and can be easily scaled up for industrial production. This makes it an attractive option for large-scale manufacturing of alumina materials for a wide range of applications.Overall, gas-phase synthesis of alumina offers a number of advantages over traditional solid-state methods, including better control over particle size and morphology, high purity, and scalability. As researchers continue to explore new applications for alumina materials, gas-phase synthesis is likely to play an increasingly important role in the production of high-quality, tailored materials for a variety of industries.。

三种典型直流特高压用氧化铝电瓷组成、结构与力学性能的对比研究

第42卷第6期2023年6月硅㊀酸㊀盐㊀通㊀报BULLETIN OF THE CHINESE CERAMIC SOCIETY Vol.42㊀No.6June,2023三种典型直流特高压用氧化铝电瓷组成㊁结构与力学性能的对比研究袁志勇,阎法强,许承铭,吴佳莉,廖仓冬,郑㊀猛,吴英豪(芦溪高压电瓷电气研究院有限公司,萍乡㊀337200)摘要:电瓷材料的组成和结构是决定其性能的两个本质因素㊂本文系统研究了三种典型直流特高压(UHVDC)用工业氧化铝电瓷在组成㊁结构和力学性能方面存在的差异,并分析了原因㊂结果表明,与日本碍子株式会社(NGK)的瓷件相比,国内两个厂家在瓷件配方中加入了更多的煅烧工业氧化铝㊁更少的钾长石和石英,并且没有根据配方适当提高烧结温度,因此国内两厂家瓷件表现出如下特点:1)Al 2O 3质量含量分别为47.22%和45.24%,比NGK 高7.31%和5.33%,SiO 2质量含量分别为46.32%和48.44%,比NGK 低7.09%和4.97%㊁K 2O +Na 2O 总质量分别为3.21%和3.14%,比NGK 低0.41%和0.48%;2)真气孔率分别为5.24%和4.18%,比NGK 低3.46%和4.52%,内部气孔以5μm 以下的小气孔为主,孔径普遍更小,存在较多扁平㊁细长㊁裂纹状的异形气孔,以及环绕颗粒周边的环状气孔和裂纹;3)内部总晶相含量分别为48.10%和49.04%,比NGK 高近10%,其中刚玉相高10%,石英相高4%,莫来石相少4%,莫来石相更细更短㊁交联程度更低;4)瓷件头部试条三点弯曲强度平均值分别为155和158MPa,比NGK 高,但是国内两厂家瓷件的强度分散性更大㊂关键词:氧化铝;电瓷;直流特高压;组成;微观结构;力学性能中图分类号:TQ174㊀㊀文献标志码:A ㊀㊀文章编号:1001-1625(2023)06-2206-09收稿日期:2023-02-08;修订日期:2023-03-28作者简介:袁志勇(1986 ),男,博士㊂主要从事绝缘子和陶瓷材料的研究㊂E-mail:yuanzy_1986@Comparative Study on Composition ,Structure and Mechanical Properties of Three Typical Alumina Insulator for UHVDCYUAN Zhiyong ,YAN Faqiang ,XU Chengming ,WU Jiali ,LIAO Cangdong ,ZHENG Meng ,WU Yinghao (Luxi High Voltage Insulator and Electricity Research Institute Co.,Ltd.,Pingxiang 337200,China)Abstract :Composition and structure are the two most essential factors for insulator materials performance.In this paper,the differences in composition,structure and mechanical properties of three typical alumina insulator for ultra high voltage direct current (UHVDC)were systematically studied,and the reasons were analyzed.The results show that compared with NGK,two domestic manufacturers have added more calcined industrial alumina,less potash feldspar and quartz into the formula,and have not properly raised the sintering temperature according to the formula,the porcelains of two domestic manufacturers show the following characteristic:1)the content of Al 2O 3are 47.22%and 45.24%,respectively,7.31%and 5.33%higher than that of NGK,the content of SiO 2are 46.32%and 48.44%,respectively,7.09%and 4.97%lower than that of NGK,and the total content of K 2O +Na 2O are 3.21%and 3.14%,respectively,0.41%and 0.48%lower than that of NGK;2)the true porosity are 5.24%and 4.18%,respectively,3.46%and 4.52%lower than that of NGK.The internal pores are basically less than 5μm,and the pore size is generally smaller.There are many flat,slender,crack like pores,as well as annular pores and cracks surrounding the particles;3)the content of internal crystalline phase is 48.10%and 49.04%,respectively,which is nearly 10%higher than NGK,with corundum phase being 10%higher,quartz phase being 4%higher,mullite phase being 4%less.The mullite phase is thinner and shorter,and the degree of crosslinking is lower;4)the average value of the three-point bending strength of the porcelain head test strip are 155and 158MPa,respectively,which are higher than NGK,but the dispersion strength of the two domestic manufacturers is higher than NGK.㊀第6期袁志勇等:三种典型直流特高压用氧化铝电瓷组成㊁结构与力学性能的对比研究2207Key words:alumina;insulator;UHVDC;composition;microstructure;mechanical property0㊀引㊀言外绝缘技术是特高压的核心技术之一,大吨位瓷绝缘子是特高压外绝缘技术的关键材料和重要部件,它的性能直接影响特高压输配电技术的可行性㊁运营的可靠性和安全性[1]㊂对于特高压大吨位瓷绝缘子,目前日本碍子株式会社(NGK)拥有国际最先进的制造水平,其产品强度分散性小㊁质量稳定好㊁零值率低㊂虽然近年来国内特高压大吨位瓷绝缘子供货厂家在结构设计㊁原材料选择㊁瓷件性能㊁制造自动化水平等方面得到显著提升,但是国内厂家生产的特高压大吨位瓷绝缘子质量与国际先进水平仍存在差距[2]㊂近年来,国内外有关特高压大吨位瓷绝缘子的研究主要集中于以铝质高强度电瓷为基础㊂作为铝质高强度电瓷的重要原料,高铝原料主要采用含刚玉晶体的原料,包括煅烧铝矾土和高温氧化铝[3-8]㊂其中,煅烧铝矾土由于Fe2O3㊁TiO2等杂质含量高㊁转化率不稳定,容易引起产品气孔率高㊁气孔形貌不良等严重缺陷,从而导致产品强度衰减快㊁劣化严重[5-6]㊂而高温氧化铝具有杂质含量低㊁刚玉相转化率高等特点,可以减少坯料杂质,提高刚玉相含量,减少瓷体缺陷,从而提高瓷件的弯曲强度㊁降低强度分散性㊁提升抗老化性能[7-8]㊂随着特高压长距离㊁大容量输变电技术的广泛应用,大吨位高强度瓷绝缘子市场需求量越来越大㊂因此,采用高温氧化铝为高铝原料已经成为瓷绝缘子制造的发展趋势㊂但是,目前我国除少数厂家在部分特定瓷绝缘子产品上采用高温氧化铝作为高铝原料生产瓷件外,大部分厂家都还在沿用煅烧铝矾土作为高铝原料㊂国内厂家对基于高温氧化铝的铝质高强度瓷的结构㊁性能研究还不够系统,高温氧化铝在瓷件中的应用技术尚不成熟㊂而国际上包括NGK㊁比彼西集团(PPC Group)和LAPP Insulators在内的电瓷制造企业都以高温氧化铝作为高铝原料,这些厂家具有成熟㊁完整㊁规范的基于高温氧化铝的电瓷生产技术㊁工艺标准和作业规范㊂对于陶瓷材料,组成㊁结构与性能之间存在紧密联系㊂对于电瓷材料,电瓷的组成(包括化学组成和物相组成)和结构是决定其性能的两个最本质的因素㊂因此本文通过对国内外3个厂家生产的基于工业氧化铝配方的大吨位瓷绝缘子用瓷件的化学组成㊁物相组成㊁显微结构(包括气孔数量㊁尺寸㊁形状和分布,晶相种类,晶粒的形貌㊁大小㊁分布等)进行分析测试比对,研究瓷件化学组成㊁物相组成㊁显微结构之间存在的关系,为国内厂家制备强度分散性小㊁质量稳定性好㊁零值率低的直流特高压用工业氧化铝电瓷提供科学依据㊂1㊀实㊀验1.1㊀试剂与材料选取日本NGK生产的400kN以及国内江西省和山东省的2个厂家生产的420kN直流特高压用工业氧化铝电瓷,共计3个样品,分别记为A㊁B㊁C㊂1.2㊀分析和测试从A㊁B㊁C瓷件头部侧壁上(见图1)分别切取3块不含釉㊁砂的小瓷块,小瓷块的体积密度利用阿基米德排水法测试,每个样品的体积密度测试结果取3块小瓷块体积密度的平均值㊂将测试完体积密度的小瓷块研磨成瓷粉,使之能完全通过63μm标准筛㊂将细粉置于(100ʃ5)ħ烘箱中烘干不少于2h后,放入干燥器中冷却至室温㊂瓷粉的化学成分采用荧光分析法,所用设备为德国BrukerS8TIGER型X-射线荧光光谱仪㊂瓷粉的物相组成采用X射线粉末衍射法测试,所用设备为德国Bruker D8Advance型X-射线衍射仪,试验条件为:Cu Kα(λ=0.15406nm),扫描速度为2(ʎ)/min,加速电压为40kV,电流为40mA,物相定量分析采用优化后的方法[9]㊂瓷粉的真密度ρt采用气体容积法,所用设备为Uatryc1200e全自动真密度分析仪㊂瓷件内部显微结构利用飞纳Phenom Prox型台式扫描电子显微镜观察㊂抛光样品制备:在瓷件头部的侧壁上切割长㊁宽均不小于5mm的瓷片,对瓷片的一个面进行研磨㊁抛光处理,用自来水冲洗后放入超声波清洗机中进行超声清洗,再将瓷片放入(100ʃ5)ħ烘箱中烘干㊂腐蚀样品制备:在ɤ25ħ的环境温度下,将研磨㊁抛光㊁清洗㊁干燥后的瓷片用质量分数为5%的氢氟酸溶液腐蚀5min,用水冲洗干净后放入超声波2208㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷清洗机中超声清洗,以去除样品表面残留的氢氟酸,再将瓷片放入(100ʃ5)ħ烘箱中烘干㊂瓷件的抗弯强度采用三点抗弯测试方法测评㊂试条从瓷件头部侧壁上切取(见图1),试条尺寸为3mm ˑ4mm ˑ50mm㊂试条经研磨抛光后,在力试LD23.104微机控制电子万能试验机上测定三点弯曲强度,测试时跨距为40mm,压头的移动速率为0.5mm /min,具体见‘精细陶瓷弯曲强度试验方法“(GB/T 6569 2006)㊂试验结果取10个试条所测数据的平均值㊂图1㊀测试试样取样位置示意图Fig.1㊀Schematic diagram of sampling location of test specimens 2㊀结果与讨论2.1㊀瓷件内部组成对比2.1.1㊀化学组成对比3个瓷件的化学组成如表1所示㊂从表1可以看出,三个厂家的直流特高压瓷绝缘子用瓷件在化学成分方面存在较大差异㊂瓷件A 中SiO 2含量高达53.41%(质量分数,下同),Al 2O 3含量仅为39.91%;瓷件B 和C 中SiO 2含量分别为46.32%和48.44%,均明显低于瓷件A;而瓷件B 和C 中Al 2O 3含量分别为47.22%和45.24%,均高于瓷件A㊂这是因为NGK 在瓷件配方中加入了较多的石英原料,而瓷件B 和C 中为提高瓷件的机械强度,在瓷件配方中加入了更多的煅烧工业氧化铝[7,10]㊂此外,瓷件A 中K 2O +Na 2O 含量为3.62%,而瓷件B㊁C 中K 2O +Na 2O 含量仅分别为3.21%和3.14%,由于配方中的K 2O +Na 2O 主要由钾长石引入,因此NGK 在瓷件配方中加入了更多的钾长石㊂研究[11]表明,配方中增加工业氧化铝用量或减少钾长石用量,都会使瓷件的烧成温度提高㊂因此为了达到相同的烧成效果,瓷件B 和C 的烧成温度应该要高于瓷件A㊂表1㊀瓷件主要化学组成Table 1㊀Main chemical composition of porcelainsSample No.Mass fraction /%SiO 2Al 2O 3Fe 2O 3TiO 2CaO MgO K 2O Na 2O BaO A 53.4139.910.810.260.120.23 3.110.51 1.52B 46.3247.220.820.380.200.25 2.950.26 1.45C 48.4445.240.920.430.180.29 2.930.21 1.25图2㊀瓷件的XRD 谱Fig.2㊀XRD patterns of porcelains 2.1.2㊀物相组成对比3个瓷件的XRD 谱见图2,物相定量分析结果列于表2㊂从表2中可以看出,3个瓷件均由刚玉㊁石英㊁莫来石和玻璃相组成,但是在物相含量方面存在较大差异㊂瓷件A 内部总晶相含量最低,仅为39.89%,而瓷件B㊁C 内部总晶相含量相对更高,分别为48.10%和49.01%㊂瓷件A 内部刚玉相含量也是最低的,仅为26.58%,而瓷件B㊁C 内部刚玉相含量均在36%以上,这是因为瓷件B 和C 在配方中加入了更多的煅烧工业氧化铝粉㊂瓷件中的刚玉一般由原料直接引入,㊀第6期袁志勇等:三种典型直流特高压用氧化铝电瓷组成㊁结构与力学性能的对比研究2209在电瓷的烧成温度下刚玉在熔体中的溶解度很小,并且在烧成温度较高或保温时间较长时,刚玉晶体可以发生长大[12]㊂虽然瓷件A在配方中加入了较多的石英原料,但是在瓷件的物相组成中其石英相含量仅为3.06%(明显低于瓷件B和C),表明其石英原料在瓷件烧成过程的高温阶段参与反应(熔解)更加充分㊂此外,瓷件A中莫来石相含量最高,高达10.25%,A中莫来石相含量基本为瓷件B㊁C内部莫来石相含量的2倍㊂对于工业氧化铝电瓷,瓷件中的莫来石相包括一次莫来石和二次莫来石,其中一次莫来石是从高岭土的分解产物偏高岭石中生成,呈颗粒㊁鳞片或短柱状;而二次莫来石是从富铝的长石熔体中析出,主要呈针状结构[10]㊂针状莫来石能相互交织构成网状结构,起到用晶须补强玻璃类似的作用,且有利于改善晶体与玻璃体的结合状态,从而提高电瓷的机械强度㊂表2㊀瓷件的物相组成Table2㊀Phase composition of porcelainsSample No.Mass fraction/%Corundum Quartz Mullite Total crystal phase Glass phase A26.58 3.0610.2539.8960.11B36.207.12 4.7848.1051.90C37.28 6.04 5.7249.0450.962.2㊀瓷件显微结构对比2.2.1㊀瓷件内部气孔3个瓷件的表面抛光形貌如图3所示㊂从图3可以看出,3个瓷件内部的气孔在尺寸㊁形状和分布等方面都存在明显差异㊂在气孔尺寸方面,瓷件A中气孔孔径分布较集中,大部分孔径在5~15μm,基本不存在5μm以下的小气孔,偶见球形的超大气孔(见图3(a));而瓷件B和C内部以5μm以下的小气孔为主,气孔孔径明显更小㊂在气孔形状方面,瓷件A中气孔基本呈近球形,气孔边缘比较圆滑,基本不存在扁平㊁细长㊁裂纹状的气孔;而瓷件B和C内部不规则的气孔较多,气孔边缘棱角较多,瓷件B和C中均存在大量扁平㊁细长㊁裂纹状气孔(见图4),尤其是瓷件B中扁平㊁细长㊁裂纹状气孔的长度超过80μm(见图4(a))㊂在气孔分布方面,瓷件A中气孔呈独立状态均匀分布在整个瓷件中,而瓷件B和C中均存在一定数量的聚集型㊁联通型气孔,见图3(f)和图3(i),此外瓷件B和C中还存在环绕颗粒周边的环状气孔和裂纹(见图5)㊂2210㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷图3㊀瓷件表面抛光形貌Fig.3㊀Polished surface morphology ofporcelains 图4㊀瓷件中的扁平㊁细长㊁裂纹状气孔Fig.4㊀Flat,slender and cracked pores inporcelains 图5㊀瓷件中颗粒周边的环状气孔和裂纹Fig.5㊀Annular pores and cracks around particles in porcelains ㊀㊀综上,瓷件B 和C 内部气孔形状不规则,且存在较多扁平㊁细长㊁裂纹状气孔和聚集型㊁联通型气孔,其根本原因是,相较于日本NGK,国内两个厂家在烧成阶段的体系内高温熔体的黏度更低㊂一方面,瓷件B 和C 在配方中钾长石㊁石英更少,即使在烧成过程中钾长石和石英均完全熔解,烧成阶段体系内高温熔体的黏度也更低㊂另一方面,国内两个厂家瓷件配方中工业氧化铝用量更高㊁钾长石用量更低,瓷件的烧成温度需更高,而实际上瓷件B 和C 的烧成温度并没有相应提高(瓷件中残余石英更多),配方中的石英颗粒(包括黏土原料带入的石英和额外添加的石英粉)未能很好地溶解于高温液相中,进一步导致高温熔体的黏度偏低㊂此外,对3个瓷件中的真气孔率进行了分析对比,利用阿基米德排水法分别测试了3个瓷件的体积密度,参照‘耐火材料真密度试验方法“(GB /T 5071 2013)测试3个瓷件的真密度,瓷件的真气孔率按公式(1)进行计算㊂第6期袁志勇等:三种典型直流特高压用氧化铝电瓷组成㊁结构与力学性能的对比研究2211㊀P =1-ρb ρt ()ˑ100%(1)式中:P 为真气孔率,%;ρb 为体积密度,g /cm 3;ρt 为真密度,g /cm 3㊂3个瓷件的体积密度㊁真密度㊁真气孔率结果见表3㊂结果表明,瓷件A 的真气孔率最高,高达8.70%,其次为瓷件B,其真气孔率为5.24%,瓷件C 的真气孔率较低,仅为4.18%,这与扫描电子显微镜观察到的结果相一致(图3(c)中气孔面积占图片总面积的比率明显高于图片3(f)和图3(i))㊂表3㊀瓷件的体积密度㊁真密度㊁真气孔率Table 3㊀Bulk density ,true density and true porosity of porcelainsSample No.Bulk density /(g㊃cm -3)True density /(g㊃cm -3)True porosity /%A 2.52 2.768.70B 2.71 2.86 5.24C 2.76 2.89 4.18按照陶瓷材料断裂力学的观点,材料破坏起因的内部缺陷包括气孔或多孔区㊁混入的杂质㊁异常生长的晶粒以及由于这些缺陷造成的结构不均匀性或各向异性,这些都作为应力集中源而降低了材料的强度,而其中气孔的危害是较大的㊂一般而言,气孔的存在使材料的有效承载面积降低,使裂纹扩展所需的能量降低[13-14]㊂研究表明,直径小于20μm 的球形㊁均匀分布的气孔对电瓷机械强度具有一定积极作用,而瓷件内部存在的扁平㊁细长㊁裂纹状㊁聚集型㊁联通型等异常气孔,在外力作用下会成为应力集中区域,从而成为瓷件断裂的根源[15]㊂瓷件B 和C 内部存在较多的异常气孔,这必然引起瓷件强度的不可预估,导致瓷件强度分散性大㊁质量稳定性和可靠性低㊂2.2.2㊀瓷件内部晶相3个瓷件的瓷件表面腐蚀形貌如图6所示,3个瓷件的瓷件能谱全元素面扫描结果见图7㊂从图6(a)~(c)和7(a)中可以观察到,瓷件A 以刚玉和莫来石为主晶相㊂瓷件A 中刚玉相均匀地分布在瓷件中,刚玉相基本都呈板状,尺寸都在10μm 以上㊂瓷件中存在大量的莫来石晶相,二次莫来石呈棒状或针状穿插在刚玉颗粒之间形成很好的相互交织状态,二次莫来石晶体细长㊁交联程度高㊁发育好,但未发现有二次莫来石异常长大的现象㊂同时可见少量的石英颗粒,石英颗粒都被熔蚀带包围,且粒径均在20μm 以下㊂相较于瓷件A,瓷件B㊁C 中刚玉相明显更多,刚玉相除板状外,还有呈细小的颗粒状,尤其是瓷件C 中存在大量的粒径在3μm 以下的粒状刚玉,见图6(h)㊁6(i)和7(c)㊂瓷件B㊁C 中存在的莫来石晶相更少,并且莫来石相更细更短㊁交联程度更低㊁发育不充分,见图6(f)㊁6(i)和7(b)㊂此外,瓷件B㊁C 中石英颗粒更多更大,甚至存在30μm 以上的大石英颗粒,见图6(d)㊁7(b)和6(g)㊁7(c)㊂电瓷产品中存在的大石英颗粒(大于30μm),一方面,石英与玻璃相的膨胀系数相差较大,在冷却时,两者收缩不一致,最终导致石英颗粒与玻璃相的界面上产生环状气孔和裂纹(见图5);另一方面,在烧成的冷却阶段,石英颗粒会在573ħ发生晶型转变,引起体积突变,产生收缩应力,所以在石英晶体内部也会产生较大尺寸的显微裂纹(见图8),这些显微裂纹在外力作用下会成为瓷件断裂的根源,所以大石英颗粒对产品性能不利[16]㊂由此可见,瓷件B 和C 内部存在的大石英颗粒,会导致瓷件强度分散性大㊁质量稳定差㊁安全性和稳定性低㊂2212㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷图6㊀瓷件腐蚀形貌Fig.6㊀Corrosion morphology ofporcelains 图7㊀瓷件的能谱全元素面扫描Fig.7㊀Energy spectra of all-element surface scanning forporcelains 图8㊀瓷件中的大石英颗粒Fig.8㊀Large quartz particles in porcelains第6期袁志勇等:三种典型直流特高压用氧化铝电瓷组成㊁结构与力学性能的对比研究2213㊀2.3㊀力学性能对比3个瓷件头部试条的三点弯曲强度测试结果见表4㊂结果表明,瓷件A头部试条的三点弯曲强度平均值为140MPa,低于瓷件B和C的155和158MPa㊂这是因为国内两个厂家在瓷件配方中加入了更多的煅烧工业氧化铝,瓷件B和C中存在更多的刚玉相,刚玉相含量增加可提高电瓷的机械强度[10]㊂从表4中还可以看出,瓷件A头部试条三点弯曲强度值均在132~148MPa,强度标准差仅为5.9MPa,分散性小㊂而瓷件B和C头部试条三点弯曲强度的标准差分别达到了13.6和11.7MPa,强度分散性较大,这是因为瓷件B和C内部存在大量的异常气孔和大石英颗粒,它们引起瓷件强度的不可预估,导致瓷件强度分散性大,这与瓷件显微结构分析结果相一致㊂表4㊀瓷件的三点弯曲强度Table4㊀Three-point bending strength of porcelains/MPaSample No.Single value12345678910AveragevalueStandarddeviationA132140134137148143148144136134140 5.9B16313417314716614215916513916615513.6C17715914813816317016215114916615811.7综合分析国内两个厂家与国外NGK瓷件在内部组成和微观结构方面存在的差异,为提高国内两个厂家瓷件的性能,主要应降低强度分散性㊁提高质量稳定性和可靠性,关键要降低瓷件中的石英相含量(尤其是大石英颗粒的含量)和避免异常气孔(包括扁平㊁细长㊁裂纹状气孔以及聚集型㊁联通型气孔)的形成,具体可以采取以下几方面措施:1)在瓷件配方上,应适当减少高温氧化铝的用量,并增加石英和钾长石的用量㊂通过减少高温氧化铝的用量和增加钾长石的用量,降低体系的理论烧成温度;增加配方中的石英用量,提高烧成过程中高温状态下液相的黏度,有效避免异常气孔的产生㊂此外增加钾长石的用量,还能增加高温状态下体系中液相总量,有利于针状二次莫来石的析出㊂2)在原料选用上,要严格控制石英原料的粒度,尽可能选用63μm及更细的石英原料,可以单独对石英原料进行研磨以达到更细的粒度㊂3)在烧成工艺上,采用高温快烧工艺㊂在瓷件的烧成温度范围之内,应尽量选用较高的烧成温度,因为体系温度越高,越有利于石英颗粒的熔融,最后在瓷件中残留下来的石英颗粒越少,石英颗粒尺寸越小;采用快烧工艺,尽可能缩短高温阶段的保温时间,这样可以有效避免瓷件中聚集型㊁联通型等异常气孔的产生㊂3㊀结㊀论本文对比分析了三种典型的直流特高压用工业氧化铝电瓷在组成㊁结构与力学性能方面存在的差异㊂结果表明,与国外NGK相比,国内两个厂家由于在瓷件配方中加入了更多的煅烧工业氧化铝㊁更少的钾长石和石英,并且没有根据配方适当提高烧结温度,因此国内两个厂家的瓷件有如下几方面特点:1)在化学组成上,Al2O3含量高5%~8%,SiO2含量低5%~7%,K2O+Na2O总含量低0.4%~0.5%㊂2)在物相组成上,内部总晶相含量高近10%,其中刚玉相高10%,石英相高4%㊁莫来石相少4%;莫来石相更细更短㊁交联程度更低㊂3)在微观结构上,真气孔率低5%~8%,内部气孔以5μm以下的小气孔为主,孔径普遍更小,存在较多扁平㊁细长㊁裂纹状的异形气孔,以及环绕颗粒周边的环状气孔和裂纹;莫来石相更细更短㊁交联程度更低,存在30μm以上的大石英颗粒,且大石英颗粒内部存在裂纹㊂4)在力学性能上,瓷件头部试条三点弯曲强度平均值高于NGK,但是两者的强度分散性更大㊂参考文献[1]㊀郭㊀雁,肖汉宁,胡文华.研磨方式对电瓷坯料粒度分布及性能的影响[J].硅酸盐通报,2011,30(5):1131-1135.GUO Y,XIAO H N,HU W H.Effect of grinding methods on the particle size distribution and performances of electrical porcelain[J].Bulletin2214㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷of the Chinese Ceramic Society,2011,30(5):1131-1135(in Chinese).[2]㊀张㊀锐,吴光亚.高压架空输电线路绝缘子行业发展和技术综述[J].电力技术,2009,18(2):47-53.ZHANG R,WU G Y.Development and technical summary of insulator industry for high voltage overhead transmission lines[J].Electric Power Standardization&Construction Cost Control Information,2009,18(2):47-53(in Chinese).[3]㊀TOUZIN M,GOEURIOT D,GUERRET-PIÉCOURT C,et al.Alumina based ceramics for high-voltage insulation[J].Journal of the EuropeanCeramic Society,2010,30(4):805-817.[4]㊀LIEBERMANN J,SCHULLE W.Bauxitic porcelain-a new high-tech product for high-voltage insulation[J].Key Engineering Materials,2001,206/207/208/209/210/211/212/213:1727-1730.[5]㊀MENG Y,GONG G H,WEI D T,et parative microstructure study of high strength alumina and bauxite insulator[J].CeramicsInternational,2014,40(7):10677-10684.[6]㊀MENG Y,GONG G H,WU Z P,et al.Fabrication and microstructure investigation of ultra-high-strength porcelain insulator[J].Journal of theEuropean Ceramic Society,2012,32(12):3043-3049.[7]㊀唐建新,王艳艳,邢政鹏.干法瓷绝缘子用微晶氧化铝配方的应用研究[J].现代技术陶瓷,2021,42(4):273-280.TANG J X,WANG Y Y,XING Z P.Study on the application of microcrystalline alumina formula for dry-process porcelain insulator[J].Advanced Ceramics,2021,42(4):273-280(in Chinese).[8]㊀范㊀刚.英国球土和工业氧化铝在高强度电瓷坯料配方中的应用[J].电瓷避雷器,2014(6):29-33+37.FAN G.Application of British clay and commercial alumina in the high strength electroceramic composition[J].Insulators and Surge Arresters, 2014(6):29-33+37(in Chinese).[9]㊀高㊀博,闫振华,赵禹来,等.瓷件微观结构对悬式瓷绝缘子机械性能和抗温度变化性能的影响研究[J].电瓷避雷器,2019(5):255-261.GAO B,YAN Z H,ZHAO Y L,et al.Study on the effect of ceramic microstructure on mechanical properties and temperature resistance of suspended porcelain insulators[J].Insulators and Surge Arresters,2019(5):255-261(in Chinese).[10]㊀SEDGHI A,RIAHI-NOORI N,HAMIDNEZHAD N,et al.Effect of chemical composition and alumina content on structure and properties ofceramic insulators[J].Bulletin of Materials Science,2014,37(2):321-325.[11]㊀成㊀茵,肖汉宁,汪㊀洋.高铝瓷中二次莫来石对电瓷强度的影响[J].电瓷避雷器,2003(6):3-5+8.CHENG Y,XIAO H N,WANG Y.Effects of secondary mullite on the mechanical strength of aluminous porcelain[J].Insulators and Surge Arresters,2003(6):3-5+8(in Chinese).[12]㊀李玉书,吴落义,李㊀瑛.电瓷工艺与技术[M].北京:化学工业出版社,2007:13-14.LI Y S,WU L Y,LI Y.Process and technology of electric porcelain[M].Beijing:Chemical Industry Press,2007:13-14(in Chinese).[13]㊀谢清云,周敬恩,沈光汉.电瓷材料显微结构与力学性能的研究[J].电瓷避雷器,2000(3):12-18+33.XIE Q Y,ZHOU J E,SHEN G H.Study on texture of electrical porcelain and its mechanical property[J].Insulators and Surge Arresters,2000(3):12-18+33(in Chinese).[14]㊀陈㊀平,王㊀芬.高强度铝质电瓷材料显微结构的研究[J].陕西科技大学学报,2004,22(5):116-121.CHEN P,WANG F.Study on the microstructure of alumina electric porcelain with high strength[J].Journal of Shaanxi University of Science& Technology,2004,22(5):116-121(in Chinese).[15]㊀ARMAN S,NASER H,NASTARAN R N.The effect of fluxes on alumina silicate porcelain insulator properties and structure[C]//Jan:International Conference on Ecological,Environmental and Biological Sciences(ICEEBS2012),2012:7-8.[16]㊀马丛丛,范发贵,符㊀成,等.高压电瓷微观结构形成机理及对宏观性能的影响[J].电瓷避雷器,2015(2):1-4+9.MA C C,FAN F G,FU C,et al.Microstructure formation mechanism of high voltage electric porcelain and influence on properties[J].Insulators and Surge Arresters,2015(2):1-4+9(in Chinese).。

claybodybuilding-AIC

78may 2015 Developing and building the best clay body for your studio practice requires knowledge of materials, ratios, limits, and of course a whole lot of testing.techno fileclay body buildingby Jonathan KaplanDefining the TermsClay Body: A combination of component clays, glass formers,fluxes, and other materials formulated for a specific firing temperature, forming process, or function.DTA Test: A thermophysicalmeasurement of the expansion of a sample material upon heating and cooling that graphically illustrates quartz inversions.Dunting: Cracking during cooling of ceramic ware due to an excess of crystobalite. Dunting also occurs when ware cools too quickly regardless of cristobalite?Eutectic: Combination of two or more ceramic materials whose melting point is lower than that of any one of the materials used alone.PCE: Pyrometric cone equivalent, a measurement of how refractory a material is.Permeability: The property that allows the flow of liquids through a body, especially in slip casting.Primary Clay: Clay that is formed but not transported from its parent site.Refractory: The resistance to heat. Also refers to the ability of a material to withstand a certain temperature without deforming.Tramp Materials: Non-clay materials that contribute to defects, e.g. lime or coal in fireclays.Secondary or Sedimentary Clay: Clay moved by geologic processes over time from its parent site.Vitrification: Firing ware to maturity where a glassy matrix is produced.Materials, Variable, Ratios, and LimitsTypical clay bodies are built with three main ingredients: clay, feldspar, and silica. Depending on the firing temperature, the ratios between plastic materials (clays) and the non-plastic materials (feldspar, silica) change to produce bodies of excellent workability (1), proper vitrification, and glaze fit. Clay bodies can generally be divided into three groups by temperature ranges: low fire (terra cotta, white talc bodies) firing to maturity between cones 06–04), mid range (white ware, stoneware, porcelain) firing to maturity between cones 4–6, and high fire (stoneware, porcelain) firing to maturity between cones 8–11.Building or designing a successful clay body is the result of an understanding of how the many materials can be combined in specific proportions to produce a desired result. The types of clays and other materials that comprise the clay body, the firing temperature, the atmosphere, and the forming method need to be considered. In addition, the shrinkage of the materials, their fired absorption, what burns off during the firing (the loss on ignition or LOI), and the thermal expansion of the clay body relative to glaze fit (a most critical criteria) must also be considered.Plastic MaterialsBoth primary and secondary (or sedimentary clays) are used to build strong clay bodies at any temperature range. Primary clays are mostly light or white burning kaolins such as Grolleg, EPK K aolin, 6 Tile , Helmer, Kingsley, McNamee, Albion, and Opticast, among others. Kaolins can lighten the color of a stoneware body and are the building blocks for porcelain clay bodies. When using kaolins as a main source of clay, understand that there are plastic kaolins as well as non-plastic kaolins. Kaolins are refractory and require adjusting the feldspar and silica amounts to achieve a vitrified body. Secondary clays are ball clays, fireclays, stoneware clays, and red clays. By varying the distribution of materials, there is less of a chance for potential problems. Fireclays are refractory and generally add strength to a mix. Coarse fireclays will contribute some textural elements to the clay while finer mesh fireclays will not. Stoneware clays are generally plastic and smooth and help with vitrification. Ball clays promote plasticity, some can be high in silica and that must be taken into considerationwith the total quartz content of the mix. Ball clays fire from buff to dark color.Some ball clays that are familiar are Old Mine #4 (OM4), Ti 21, Foundry Hill Creme, Kentucky Stone, FC 340, C & C Ball Clay, Coppen Light, and Champion. Common available fireclays are Hawthorne, Greenstripe, Plainsman Red, Cedar Heights fireclay, Lincoln, Green Ribbon, C Red, and Imco. Familiar stoneware clays include, but are not limited to; Cedar Heights Goldart, and Roseville. Common available red clays are Cedar Heights Redart, and Newman Red.Some clays are specifically processed and blended for a particular forming method. For example, Opticast and FC 340 clays are blended specifically for slip casting. Their particle size is optimized for permeability.Non-Plastic MaterialsAdding silica to a clay body promotes vitrification and glass forming within the ceramic matrix. According to Digitalfire (), the terms “flint, quartz, and silica have come to be used interchangeable in ceramics and you will see them all employed in recipes; they are all the same thing. However, most correctly, the material used in ceramics is silica. Quartz refers to the macro-crystalline mineral we find in nature.”Silica melts at a temperature much higher than we fire to in the studio, 2912–3137°F (1600–1725°C). So in order to lower the meting point to get our clay bodies to vitrify, we need to add melters. Clays contain small amounts of auxiliary fluxes, which can include iron, titanium dioxide, calcium oxide, magnesium oxide, sodium oxide, calcium, magnesium, and lithium, and potassium oxide. When combined with feldspars, which have some of the same fluxes (calcium, magnesium, and lithium), viable eutectics at temperature ranges used by potters and ceramic artists can be achieved. The melting of quartz by these fluxes promotes vitrification. As a result it also lessens the amount of excess quartz, known as free silica, which is highly undesirable because it causes problems and defects as the fired ware cools. Feldspars also contribute alumina and silica.Feldspars such as Custer, G-200, Nepheline Syenite, and Primas are used in mid-range and high-temperature clay bodies. In low-temperature clay bodies, talc is used. Small amounts of frit can be added at the lower temperature range but particular attention needs to be paid to the upper limit temperature of the bisque firing to avoid beginning to vitrify the ware, making it more difficult to glaze. Pyrophyllite, a hydrated alumina silicate, increases thermal shock resistance and promotes the fired strength of clay bodies by developing mullite. It can be substituted for an equal amount of silica.Add grog, which is usually ground up firebrick, to a clay body to provide texture, also known as tooth, to promote even drying and contribute to the fired strength. Refractory calcines such as Mulcoa or Molochite are blended commercial grogs. They are available in many different mesh sizes and are a white-firing alternative to the darker firebrick-type grog, although they are more expensive.Other materials that can be added to clay bodies are bentonites to increase plasticity and workability in bodies that are short; wollastonite to add calcium and silica, help to lower thermal expansion, and increase fired strength; mullite to promote strength and vitrification; and vinegar or inorganic materials such as Pro Bond, Additive A, and other compositions to increase plasticity.Limits for MaterialsRatios of Plastics to Non-PlasticsListed below is a fairly broad range for material limits. These providea good starting point from which to begin developing a clay body (2). The proportion of plastics to non-plastics at the three major temperature ranges is important in developing clay bodies that have good workability and the necessary vitrification (1). These ratios work very well but can certainly be adjusted to suit specific needs. Note that in all temperature ranges, the casting ratio is the same.Shrinkage and AbsorptionAll clays shrink after the physical water evaporates and the body becomes vitreous and dense. The finer the particle size, the greaterthe shrinkage. Fine particle-sized ball clays shrink the most and the coarser particle-size clays such as fireclays, shrink less (3). Shrinkage and absorption at low-fire temperatures are greater than at higher temperatures. Porcelain clay bodies shrink more than stoneware clay bodies. Proper material selection and making accurate shrinkage test bars of the component clays as well as the clay bodies that use themwill provide a great deal of information. The same thinking should also be applied to the absorption or porosity of a clay body. A clay body with a low absorption will be more durable and vitreous. It will be less prone to fail under use, less prone to delayed moisturecrazing, and provide lasting use. may 201579techno fileLoss on IgnitionThe loss on ignition, or LOI, represents the amount of carbonaceous and tramp material that burns off during the initial stages of a firing. By mixing a small amount of ball clay or fireclay with water to form a slurry and passing it through a 60-mesh screen, we can see the residue of sand, small pebbles, bits of coal, limestone, twigs, etc. that need to burn off. As these materials outgas during bisque firing, the importance of a slow bisque firing that reaches a temperature high enough to burn off these materials is critical in preventing glaze defects such as pinholing in glazes or white spots in majolica. Whether one uses commercially prepared clay or mixes their own clay body, a proper bisque firing is an important step to producing defect-free ware.Thermal ExpansionEvery clay company will provide a material analysis of what oxides are present in their clay. Consideration should be paid to the amount of silica. These numbers will play a big role in the expansion of the clay body, which is pivotal in avoiding problemsat the major quartz inversions. Clay and glazecalculation software is very useful in providingmetrics on the amount of silica present as wellas the thermal expansion of the clay body.Putting It All TogetherWe can glean a great deal of informationfrom simple testing of component clays forabsorption and shrinkage and then develop aseries of clay body tests using a spreadsheet (4).Procedure: Dry mix 10 pounds of each testbody in a plastic bag. Mix the dry clay bodyto slip consistency so that it can be screenedthrough a 60-mesh sieve. Note the screenresidue. Dry the slip mixture on a plaster slabto throwing consistency and wedge. Maketest bars for shrinkage/absorption, and thenuse the balance of the clay to throw test-tilerings and small bowls and plates for furtherglaze testing. Should a suitable clay body beobtained, also run a freeze/thaw test, thenput a slice of lemon on a small glazed plateovernight to check for acid attack. Then runa few pieces through a dishwasher cycleto check for alkali attack. Then test again,and again. Finally, send an appropriatesample to a laboratory for a DTA test. Thesemethodological tests of your clay body willinsure that your formula has the necessaryproperties to be successfully used in yourstudio practice.A final note: materials change over time andalso go out of production. Ongoing testingand modification over time assures that yourclay body retains its necessary characteristics.the author Jonathan Kaplan’s ceramic career has spanned over 40 years. He earned an MFA from Southern Illinois University at Ed-wardsville and has worked as a production potter, university educator, ceramic artist, professional mold and modelmaker, ceramic designer, and manufacturer. He serves on the board of Studio Potter and curates Plinth Gallery in Denver, Colorado.ReferencesClay and Glazes for the Potter by Daniel RhodesStoneware and Porcelain by Daniel RhodesCeramic Science for the Potter by William LawrenceCushing’s Handbook by Val CushingClay: A Studio Handbook by Vince PitelkaCeramic Technology for Potters and Sculptors by Yvonne CuffClay Bodies by Robert TichaneThe Potters’ Dictionary of Materials and Techniques by Frank and Janet Hamer Clay and Glaze Handbook by Jeff ZamekCeramics: A Potter’s Handbook by Glenn C. NelsonOut of the Earth and Into the Fire by Mimi ObstlerChemistry and Physics of Clay by Rex GrimshawDigitalfire:/4sight/education/index.htmlThe above clay bodies and their component clays are a starting point for investigation. Thethrowing bodies are meant to be smooth and the casting bodies are set to be a 50/50 ratio. Whatworks for one person may not work for another. Remember to always test and record your results.80may 2015 。

Al—Si—O系陶瓷基电子封装材料粉体的制备

目录摘要 (1)Abstract (3)1 绪论 (5)1.1 陶瓷封装材料 (6)1.2 金属基电子封装材料简介 (6)1.3 氧化铝陶瓷电子封装基片特点 (7)1.4 Al-Si-O复合粉体制备方法 (7)1.5 喷雾干燥原理及特点 (10)1.6 粉体包覆技术 (11)1.7 研究目的和意义及研究内容 (12)1.7.1 研究目的及意义 (12)1.7.2 研究的内容 (12)2 实验部分 (13)2.1 实验原料 (13)2.2 实验步骤 (13)2.2.1 溶胶凝胶法制备二氧化硅 (13)2.2.2 SiO2表面包覆Al2O3的研究 (14)2.3 溶胶凝胶法制备二氧化硅粉体 (15)2.4 样品的表征方法 (15)2.4.1 XRD分析 (15)2.4.2 扫描电镜分析（SEM） (16)2.4.3 透射电镜分析（TEM） (16)2.4.4 X射线能量色散谱分析 (17)2.5 实验用主要仪器设备 (17)3 实验结果与分析 (19)3.1 溶胶凝胶法制备二氧化硅 (19)3.1.1 溶胶凝胶-喷雾干燥制备二氧化硅的XRD谱 (19)3.1.2溶胶凝胶-喷雾干燥制备二氧化硅的SEM图 (19)3.2 非均相成核法制备氧化铝包覆二氧化硅复合粉体 (21)3.2.1 不同浓度硝酸铝滴加生成的粉体TEM照片 (21)3.2.2不同浓度硝酸铝滴定生成的粉体SEM照片 (23)3.2.3 烧结后粉体的SEM图、XRD和能谱分析图 (24)结论 (27)致谢............................................. 错误！未定义书签。

[参考文献]......................................... 错误！未定义书签。

Al-Si-O系陶瓷基电子封装材料粉体的制备摘要：电子封装材料是指用作基片、底板、外壳等来支撑和保护半导体芯片和电子电路等，同时又起到散热和/或导电的作用的一类材料的总称。

溶胶凝胶法制备mullite过程中的拉曼光谱和红外光谱研究

溶胶凝胶法制备mullite过程中的拉曼光谱和红外光谱研究谭小平;梁叔全;秦利平【摘要】The formation of Si and Al sols, and the micro-structural changes of the dual phase Si-Al gel in the in-situ crystallizing process were investigated by Laser Raman(LRa), Fourier Transform infrared spectroscopy (FT-IR) and X-ray diffraction (XRD). The results show that the disordered network structure is composed of Si—O…H and Al—O…H groups in the Si and Al sols, respectively. After the thermal treatment at 600℃, the dual phase Si-Al gel is obtained with non organic impurity phase. At 800 ℃, the amorphous structure of the Si-Al double gel apparently reorganizes, resulting in the formation of a small amount of AlO4 tetrahedron and AlO6 octahedral units. The Si-Al spinel was crystallized from the sample at 1 000 ℃. With the increase of temperature, the peak of Al-Si spinel enhances and then disappears in the XRD pattern. At 1 200 ℃, the characteristic FT-IR bands of mullite at 1 130 cm−1 and 1 170 cm−1 indicates that the Al—O—Si bond is formed, and the Raman bands at 310, 410, 600, 960, 1 037 and 1 129 cm−1 typical for mullite occur. When the temperature increases to 1 450℃, the main crystalline phases are identified as mullite with a small amount of Alumina, the spinel phase disappears. These results indicate mullite is formed from the Si-Al spinel, which is metastable phase at 1 200−1 400℃in the heating process of the dual phase gel.%借助Raman，FT-IR和XRD等技术探讨Si和Al溶胶的形成及Si-Al双相凝胶在原位受控晶化过程中微结构变化。