Annealing effect on the crystal structure and exchange bias in Heusler

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单斜相阿利特的晶体结构研究(英文)

T1 ←⎯620⎯℃→ T2 ←⎯920⎯℃→ T3 ←⎯980⎯℃→ M1 ←⎯990⎯℃→ M2 ←⎯1 06⎯0 ℃→ M3 ←⎯1 07⎯0 ℃→ R
Recent work have concerned the identification of stabilized modifications at room temperature due to the importance in the quality control of Portland cement. Differential thermal analysis (DTA) showes the T1–T2 transi-
The objective of this paper was to describe the superstructures of the monoclinic modification of C3S. Some indices for describing the reflections due to the superstructures were proposed to express the orientations and
1 Experimental
The specimens were provided by China Building Materials Academy (CBMA). A selected-area electron diffraction (SAED) technique was applied to record the reflections. A TEM device (Model JEM–2010UHR, JEOL, Tokyo, Japan) equipped with a double tilt goniometry was used to record the HRTEM images at an accelerating voltage of 200 kV. The preparation of specimens for the TEM was performed by a conventional method applied for the observation of ceramic powders except for the dispersion of the specimens in ethanol. Table 1 shows the chemical compositions of the specimens, which was provided by the CBMA. The three specimens are remarked as A, B, and C.

The Properties of Crystals and Crystal Structures

The Properties of Crystals and CrystalStructuresCrystal structures have always fascinated science lovers and researchers alike. The beautiful and intricate patterns that crystals exhibit are not just aesthetically pleasing; they also provide important information about the physical and chemical properties of the crystals themselves. In this article, we will explore the properties of crystals and crystal structures and how they impact various scientific fields.What are crystals?Crystals are solids that have highly ordered structures, meaning that their atoms or molecules are arranged in a repeating pattern. This pattern is what gives crystals their characteristic geometric shape. Crystals can be formed from a wide variety of materials, including minerals, metals, and organic compounds.One of the defining features of crystals is that they have repeating units called unit cells. The unit cell is the smallest part of a crystal that still exhibits the same structural pattern as the whole crystal. By analyzing the unit cell, scientists can determine the basic structure of a crystal.What are the properties of crystals?One of the most important properties of crystals is their symmetry. Because crystals have an ordered structure, their symmetry is also highly organized. This symmetry is what gives crystals their characteristic shapes and also affects their physical properties, such as their melting point and conductivity.Another important property of crystals is their cleavage. Cleavage refers to the way in which a crystal breaks along certain planes. This property is determined by the arrangement of atoms within the crystal structure and can be used to identify different types of crystals.Crystal structures and their importance in scienceCrystal structures play an important role in various scientific fields, including chemistry, physics, and materials science. By understanding the structure of crystals, scientists can predict their physical and chemical properties, which can be used to develop new materials for various applications.For example, the development of new drugs often relies on an understanding of the crystal structure of the active ingredient. By analyzing the crystal structure, scientists can determine how the drug interacts with its targets and how it can be modified to increase its effectiveness.Crystal structures are also important in the field of materials science. By studying the crystal structure of materials, scientists can determine their mechanical and electrical properties. This information can be used to develop new materials with specific properties, such as advanced ceramics for use in electronics or stronger metals for use in aerospace applications.ConclusionIn conclusion, crystals and crystal structures are fascinating objects that provide important information about the physical and chemical properties of materials. The highly ordered structure of crystals gives them unique properties that can be harnessed for a variety of scientific and practical applications. By continuing to study crystals and their structures, scientists can unlock new insights into the world around us and develop new materials that will shape our future.。

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单位：河北工程大学
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Annu.Rev.Mater.Res.31_1_2001

Annu.Rev.Mater.Res.2001.31:1–23Copyright c2001by Annual Reviews.All rights reserved S YNTHESIS AND D ESIGN OF S UPERHARDM ATERIALSJ Haines,JM L´e ger,and G BocquillonLaboratoire de Physico-Chimie des Mat´e riaux,Centre National de la Recherche Scientiﬁque,1place Aristide Briand,92190Meudon,France;e-mail:haines@cnrs-bellevue.fr;leger@cnrs-bellevue.frKey Words diamond,cubic boron nitride,carbon nitride,high pressure,stishovite s Abstract The synthesis of the two currently used superhard materials,diamond and cubic boron nitride,is brieﬂy described with indications of the factors inﬂuencing the quality of the crystals obtained.The physics of hardness is discussed and the importance of covalent bonding and ﬁxed atomic positions in the crystal structure,which determine high hardness values,is outlined.The materials investigated to date are described and new potentially superhard materials are presented.No material that is thermodynamically stable under ambient conditions and composed of light (small)atoms will have a hardness greater than that of diamond.Materials with hardness values similar to that of cubic boron nitride (cBN)can be obtained.However,increasing the capabilities of the high-pressure devices could lead to the production of better quality cBN compacts without binders.INTRODUCTIONDiamond has always fascinated humans.It is the hardest substance known,based on its ability to scratch any other material.Its optical properties,with the highest refraction index known,have made it the most prized stone in jewelry.Furthermore,diamond exhibits high thermal conductivity,which is nearly ﬁve times that of the best metallic thermal conductors (copper or silver)at room temperature and,at the same time,is an excellent electrical insulator,even at high temperature.In industry,the hardness of diamond makes it an irreplaceable material for grinding tools,and diamond is used on a large scale for drilling rocks for oil wells,cutting concrete,polishing stones,machining,and honing.The diamonds used for industry are now mostly man-made because their cutting edges are much sharper than those of natural diamonds,which have been eroded by geological time.The synthesis of diamond has been a goal of science from Moissant at the end of the nineteenth century to the successful synthesis under high pressures in 1955(1).However,diamond has a major drawback in that it reacts with iron and cannot be used for machining steel.This has prompted the synthesis of a second superhard0084-6600/01/0801-0001$14.001A n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r gb y C h i n e s e Ac ade m y of S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .2HAINESL ´EGERBOCQUILLONmaterial,cubic boron nitride (cBN),whose structure is derived from that of dia-mond with half the carbon atoms being replaced by boron and the other half by nitrogen atoms.The resulting compound is half as hard as diamond,but it does not react with iron and can be used for machining steel.Cubic boron nitride does not exist in nature and is prepared under high-pressure high-temperature conditions,as is synthetic diamond.However,its synthesis is more difﬁcult,and it has not been possible to prepare large crystals.Industry is thus looking for new superhard ma-terials that will need to be much harder than present ceramics (Si 3N 4,Al 2O 3,TiC).Hardness is a quality less well deﬁned than many other physical properties.Hardness was ﬁrst deﬁned as the ability of one material to scratch another;this corresponds to the Mohs scale.This scale is highly nonlinear (talc =1,diamond =10);however,this deﬁnition of hardness is not reliable because materials of similar hardness can scratch each other and the resulting value depends on the speciﬁc details of the contact between the two materials.It is well known (2)that at room temperature copper can scratch magnesium oxide and at high temperatures cBN can scratch diamond (principle of soft indenter).Another,more accurate,way of deﬁning and measuring hardness is by the indentation of the material by a hard indenter.According to the nature and shape of the indenter,different scales are used:Brinell,Rockwell,Vickers,and Knoop.The last two are the most frequently used.The indenter is made of a pyramidal-shaped diamond with a square base (Vickers),or elongated lozenge (Knoop).The hardness is deduced from the size of the indentation produced using a deﬁned load;the unit is the same as that for pressure,the gigapascal (GPa).Superhard materials are deﬁned as having a hardness of above 40GPa.The hardness of diamond is 90GPa;the second hardest material is cBN,with a hardness of 50GPa.The design of new materials with a hardness comparable to diamond is a great challenge to scientists.We ﬁrst describe the current status of the two known super-hard materials,diamond and cBN.We then describe the search for new bulk super-hard materials,discuss the possibility of making materials harder than diamond,and comment on the new potentially superhard materials and their preparation.DIAMOND AND CUBIC BORON NITRIDE DiamondThe synthesis of diamond is performed under high pressure (5.5–6GPa)and high temperature (1500–1900K).Carbon,usually in the form of graphite,and a transi-tion metal,e.g.iron,cobalt,nickel,or alloys of these metals [called solvent-catalyst (SC)],are treated under high-pressure high-temperature conditions;upon heating,graphite dissolves in the metal and if the pressure and temperature conditions are in the thermodynamic stability ﬁeld of diamond,carbon can crystallize as dia-mond because the solubility of diamond in the molten metal is less than that of graphite.Some details about the synthesis and qualities of diamond obtained by this spontaneous nucleation method are given below,but we do not describe the growthA n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r gb y C h i n e s e Ac ade m y of S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .SUPERHARD MATERIALS 3of single-crystal diamond under high pressure,which is necessary in order to ob-tain large single crystals with dimensions greater than 1mm.Crystals of this size are expensive and represent only a very minor proportion of the diamonds used for machining;they are principally used for their thermal properties.It is well known that making diamonds is relatively straightforward,but control-ling the quality of the diamonds produced is much more difﬁcult.Improvements in the method of synthesis since 1955have greatly extended the size range and the mechanical properties and purity of the synthetic diamond crystals.Depending on the exact pressure (P )and temperature (T )of synthesis,the form and the nature of the carbon,the metal solvent used,the time (t )of synthesis,and the pathways in P-T-t space,diamond crystals (3,4)varying greatly in shape (5),size,and fri-ability are produced.These three characteristics are used to classify diamonds;the required properties differ depending on the industrial application.Friability is related to impact strength.It is the most important mechanical property for the practical use of superhard materials,and low friability is required in order for tools to have a long lifetime.In commercial literature,the various types of diamonds are classed as a function of their uses,which depend mainly on their friabilities,but the numerical values are not given,so it is difﬁcult to compare the qualities of diamonds from various sources.The friability,which is deﬁned by the percentage of diamonds destroyed in a speciﬁc grinding process,is obtained by subjecting a deﬁned quantity of diamonds to repeated impacts by grinding in a ball-mill or by the action of a load falling on them.The friability values depend strongly on the experimental conditions used,and only values for crystals measured under the same conditions can be compared.The effect of various synthesis parameters on their quality can be evaluated by considering the total mass of diamond obtained in one experiment,the distribution size of these diamonds,and the friability of the diamonds of a deﬁned size.A ﬁrst parameter is the source of the carbon.Most carbon-based substance can be used to make diamonds (6),but the nature of the carbon source has an effect on the quantity and the quality of synthetic diamonds.The best carbon source for diamond synthesis is graphite,and its characteristics are important.The effect of the density,gas permeability,and purity of graphite on the diamond yield have been investigated using cobalt as the SC (7).Variations of the density and gas permeability have no effect on the diamond yield,but carbon purity is important.The main impurity in synthetic graphite is CaO.If good quality diamonds are required,the calcium content should be kept below 1000ppm in order to avoid excessive nucleation on the calcium oxide particles.A second factor that alters the quality of diamonds is the nature of the SC.The friability and the size distribution are better with CoFe (alloy of cobalt with a small quantity of iron)than with invar,an iron-nickel alloy (Table 1:Ia,Ib;Figure 1a ).Another parameter is how the mixture of carbon and SC is prepared.When ﬁne or coarse powders of intimately mixed graphite and SC are used,a high yield of diamonds with high friabilities is obtained (8).These diamonds are very small,A n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .4HAINESL ´EGERBOCQUILLONTABLE 1Friabilities of some diamonds as a function of the details of the synthesis process fora selected size of 200–250µmSDA a MBS a 100+70Ia Ib IIa IIb IIc IIIa IIIb Synthesis CoFeInvar 1C-2SC 1C-1SC 2C-2SC Cycle A Cycle B detail stacking stacking stacking Friability 74121395053637250(%)aThe SDA 100+is among De Beers best diamond with a very high strength that is recommended for sawing the hardest stones and cast refractories.MBS 70is in the middle of General Electric’s range of diamonds for sawing and drilling.Other diamonds were obtained in the laboratory using a belt–type apparatus with a working chamber of 40mm diameter.(C,graphite;SC,solvent-catalyst.)with metal inclusions,and they are linked together with numerous cavities ﬁlled with SC.A favorable geometry in order to obtain well-formed diamonds is to stack disks of graphite and SC.The effect of local concentration has been exam-ined by changing the stacking of these disks (Table 1:IIa,IIb,IIc;Figure 1b ).The method of stacking modiﬁes the local oversaturation of dissolved carbon and thus the local spontaneous diamond germination.For the synthesis of dia-mond,the heating current goes directly through the graphite-SC mixture.Because the electrical resistivity of the graphite is much greater than that of the SC,the temperature of the graphite is raised by the Joule effect,whereas that of the SC increases mainly because of thermal conduction.Upon increasing the thickness of the SC disk,the local thermal gradient increases and the dissolved atoms of carbon cannot move as easily;the local carbon oversaturation then enhances the spontaneous diamond germination.This enables one to work at lower tempera-tures and pressures,which results in slower growth and therefore better quality diamonds.Another important factor for the yield and the quality of the diamonds is the pathway followed in P-T-t space.The results of two cycles with the same ﬁnal pressure and temperature are shown.In cycle A (Figure 1d ),the graphite-SC mixture reaches the eutectic melting temperature while it is still far from the equilibrium line between diamond and graphite;as a result spontaneous nucleation is very high and the seeds grow very quickly.These two effects explain the high yield and the poor quality and small size and high friability of the diamonds compared with those obtained in cycle B (Figure 1d ;see Table 1IIIa and IIIb and Figure 1c ).Large crystals (over 400µm)of good quality are obtained when the degree of spontaneous nucleation is limited.The pathway in P-T-t space must then remain near the graphite-diamond phase boundary (Figure 1d ),and the time of the treatment must be extended in the ﬁnal P-T-t conditions.Usually,friability increases with the size of the diamonds.Nucleation takes place at the beginning of the synthesis when the carbon oversaturation is important,and the carbon in solution is then absorbed by the existing nuclei,which grow larger.A n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .SUPERHARD MATERIALS5Figure 1Size distribution of diamonds in one laboratory run for different synthesis pro-cesses;effects of (panel a )the nature of the metal,(panel b )the stacking of graphite and metal disks,(panel c )the P-T pathway.(Panel d )P-T pathways for synthesis.1:graphite-diamond boundary and 2:melting temperature of the carbon-eutectic.The diamond synthesis occurs between the boundaries 1and 2.The growth time is about the same for all the crystals,thus those that can grow more quickly owing to a greater local thermal gradient become the largest.Owing to their rapid growth rate,they trap more impurities and have more defects,and therefore their friability is higher.Similarly,friability increases with the diamond yield.The diamonds produced by the spontaneous nucleation method range in size up to 800–1000µm.The best conditions for diamond synthesis correspond to a compromise between the quantity and the quality of the diamonds.A n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .6HAINESL ´EGERBOCQUILLONCubic Boron NitrideCubic boron nitride (cBN)is the second hardest material.The synthesis of cBN isperformed in the same pressure range as that for diamond,but at higher tempera-tures,i.e.above 1950K.The general process is the same;dissolution of hexagonal boron nitride (hBN)in a solvent-catalyst (SC),followed by spontaneous nucle-ation of cBN.However,the synthesis is much more complicated.The usual SCs are alkali or alkaline-earth metals and their nitrides (9):Mg,Ca,Li 3N,Mg 3N 2,and Ca 3N 2.All these SCs are hygroscopic,and water or oxygen are poisons for the synthesis.Thus great care must be taken,which requires dehydration of the materials and preparation in glove boxes,to avoid the presence of water in the high-pressure cell.Furthermore,the above compounds react ﬁrst with hBN to form inter-mediate compounds,Li 3BN 2,Mg 3B 2N 4,or Ca 3B 2N 4,which become the true SC.These compounds and the hBN source are electrical insulators,thus an internal furnace must be used,which makes fabrication of the high pressure cell more complicated and reduces the available volume for the samples.In addition,the chemical reaction involved is complicated by this intermediate step,and in gen-eral the yield of cBN is lower than for diamond.Work is in progress to determine in situ which intermediate compounds are involved in the synthesis process.The crystals of cBN obtained from these processes are of lower quality (Figure 2)and size than for diamond.Depending on the exact conditions,orange-yellow or dark crystals are obtained;the color difference comes from a defect or an excess of boron (less than 1%);the dark crystals,which have an excess of born,are harder.As in the synthesis of diamond,the initial forms of the SC source,hBN,play important roles,but the number of parameters is larger.For the source of BN,it is better to use pressed pellets of hBN powder rather than sintered hBN products,as the latter contain additives (oxides);a very ﬁne powder yields a better reactivity.Doping of Li,Ca,or Mg nitrides with Al,B,Ti,or Si induces a change in the morphology and color of cBN crystals,which are dark instead of orange,are larger (500µm),have better shapes and,in addition,gives a higher yield (10).Use of supercritical ﬂuids enables cBN to be synthesized at lower pressures and temperatures (2GPa,800K),but the resulting crystal size is small (11).Diamond and cBN crystals are produced on a large scale,and the main problem is how to use them for making viable tools for industry.Different compacts of these materials are made (12)for various pacts of diamonds are made using cobalt as the SC.The mixture is treated under high-pressure high-temperature conditions,at which superﬁcial graphitization of the diamonds takes place,and then under the P-T-t diamond synthesis conditions so as to transform the graphite into diamond and induce intergranular growth of diamonds.The diamond compacts produced in this way still contain some cobalt as a binder,but their hardness is close to that of single-crystal pacts of cBN cannot be made in the same way because the SCs are compounds that decompose in air.Sintering without binders (13)is possible at higher pressures of about 7.5–8GPa and temperatures higher than 2200K,but these conditions are currently outside the range of thoseA n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y.SUPERHARD MATERIALS7Figure 2SEM photographs of diamond (top )and cBN (bottom )crystals of different qualities depending on the synthesis conditions (the long vertical bar corresponds to a distance of 100µm).Top left :good quality mid-sized diamonds of cubo-octahedral shape with well-deﬁned faces and sharp edges;top right :lower quality diamonds;bottom left :orange cBN crystals;bottom right :very large black cBN crystals of better shapes.A n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .8HAINESL ´EGERBOCQUILLONused in industrial pacts of cBN with TiC or TaN binders are ofmarkedly lower hardness because there is no direct bonding between the superhard crystals,in contrast to diamond compacts.In addition,they are expensive,and this has motivated the search for other superhard materials.SEARCH FOR NEW SUPERHARD MATERIALSOne approach for increasing hardness of known materials is to manipulate the nanostructure.For instance,the effect of particle size on the hardness of materials has been investigated.It is well known that high-purity metals have very low shear strengths;this arises from the low energy required for nucleation and motion of dislocations in metals.The introduction of barriers by the addition of impurities or grain size effects may thus enhance the hardness of the starting phase.In this case,intragranular and intergranular mechanisms are activated and compete with each other.As each mechanism has a different dependency on grain size,there can be a maximum in hardness as the function of the grain size.This effect of increasing the hardness with respect to the single-crystal value does not exist in the case of ceramic materials.In alumina,which has been thoroughly studied,the hardness (14)of ﬁne-grained compacts is at most the hardness of the single crystal.When considering superhard materials,any hardness enhancement would have to come from the intergranular material,which would be by deﬁnition of lower hard-ness.In the case of thin ﬁlms,it has been reported that it is possible to increase the hardness by repeating a layered structure of two materials with nanometer scale dimensions,which are deposited onto a surface (15).This effect arises from the repulsive barrier to the movement of dislocations across the interface between the two materials and is only valid in one direction for nanometer scale defor-mations.This could be suitable for coatings,but having bulk superhard materials would further enhance the unidirectional hardness of such coatings.In addition,hardness in these cases is determined from tests at a nanometer scale with very small loads,and results vary critically (up to a factor of three)with the nature of the substrate and the theoretical models necessary to estimate quantitatively the substrate’s inﬂuence (16).We now discuss the search for bulk superhard materials.Physics of HardnessThere is a direct relation between bulk modulus and hardness for nonmetallic ma-terials (Figure 3)(17–24),and here we discuss the fundamental physical properties upon which hardness depends.Hardness is deduced from the size of the inden-tation after an indenter has deformed a material.This process infers that many phenomena are involved.Hardness is related to the elastic and plastic properties of a material.The size of the permanent deformation produced depends on the resistance to the volume compression produced by the pressure created by the indenter,the resistance to shear deformation,and the resistance to the creation and motion of dislocations.These various types of resistance to deformation indicateA n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .SUPERHARD MATERIALS 9Figure 3Hardness as a function of the bulk modulus for selected materials (r-,rutile-type;c,cubic;m,monoclinic).WC and RuO 2do not ﬁll all the requirements to be superhard (see text).which properties a material must have to exhibit the smallest indentation possible and consequently the highest hardness.There are three conditions that must be met in order for a material to be hard:The material must support the volume decrease created by the applied pressure,therefore it must have a high bulk modulus;the material must not deform in a direction different from the applied load,it must have a high shear modulus;the material must not deform plastically,the creation and motion of the dislocations must be as small as possible.These conditions give indications of which materials may be superhard.We ﬁrst consider the two elastic properties,bulk modulus (B)and shear modulus (G),which are related by Poisson’s ratio (ν).We consider only isotropic materials;a superhard material should preferably be isotropic,otherwise it would deform preferentially in a given direction (the crystal structure of diamond is isotropic,but the mechanical properties of a single crystal are not fully isotropic because cleavage may occur).In the case of isotropic materials,G =(3/2)B (1−2ν)/(1+ν);In order for G to be high,νmust be small,and the above expression reduces thenA n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .10HAINESL ´EGERBOCQUILLONto G =(3/2)B (1−3ν).The value of νis small for covalent materials (typicallyν=0.1),and there is little difference between G and B:G =1.1B.A typical value of νfor ionic materials is 0.25and G =0.6B;for metallic materials νis typically 0.33and G =0.4B;in the extreme case where νis 0.5,G is zero.The bulk and shear moduli can be obtained from elastic constants:B =(c 11+2c 12)/3,G =(c 11−c 12+3c 44)/5.Assuming isotropy c 11−c 12=2c 44,it follows that G =c 44;Actually G is always close to c 44.In order to have high values of B and G,then c 11and c 44must be high with c 12low.This is the opposite of the central forces model in which c 12=c 44(Cauchy relation).The two conditions,that νbe small and that central forces be absent,indicate that bonding must be highly directional and that covalent bonding is the most suitable.This requirement for high bulk moduli and covalent or ionic bonding has been previously established (17–19,21–24)and theoretical calculations (19,25,26)over the last two decades have aimed at ﬁnding materials with high values of B (Figure 3).The bulk modulus was used primarily for the reason that it is cheaper to calculate considering the efﬁcient use of computer time,and an effort was made to identify hypothetical materials with bulk moduli exceeding 250–300GPa.At the present time with the power of modern computers,elastic constants can be obtained theoretically and the shear modulus calculated (27).The requirement for having directional bonds arises from the relationship be-tween the shear modulus G and bond bending (28).Materials that exhibit lim-ited bond bending are those with directional bonds in a high symmetry,three-dimensional lattice,with ﬁxed atomic positions.Covalent materials are much better candidates for high hardness than ionic compounds because electrostatic interac-tions are omnidirectional and yield low bond-bending force constants,which result in low shear moduli.The ratio of bond-bending to bond-stretching force constants decreases linearly from about 0.3for a covalent material to essentially zero for a purely ionic compound (29,30).The result of this is that the bulk modulus has very little dependence on ionicity,whereas the shear modulus will exhibit a relative de-crease by a factor of more than three owing entirely to the change in bond character.Thus for a given value of the bulk modulus,an ionic compound will have a lower shear modulus than a covalent material and consequently a lower hardness.There is an added enhancement in the case of ﬁrst row atoms because s-p hybridization is much more complete than for heavier atoms.The electronic structure also plays an important role in the strength of the bonds.In transition metal carbonitrides,for example,which have the rock-salt structure,the hardness and c 44go through a maximum for a valence electron concentration of about 8.4per unit cell (31).The exact nature of the crystal and electronic structures is thus important for determin-ing the shear modulus,whereas the bulk modulus depends mainly on the molar volume and is less directly related to ﬁne details of the structure.This difference is due to the fact that the bulk modulus is related to the stretching of bonds,which are governed by central forces.Materials with high bulk moduli will thus be based on densely packed three-dimensional networks,and examples can be found among covalent,ionic,and metallic materials.In ionic compounds,the overall structure isA n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .principally deﬁned by the anion sublattice,with the cations occupying interstitial sites,and compounds with high bulk moduli will thus have dense anion packing with short anion-anion distances.The shear modulus,which is related to bond bending,depends on the nature of the bond and decreases dramatically as a func-tion of ionicity.In order for the compound to have a high shear modulus and high hardness (Figure 4),directional (covalent)bonding and a rigid structural topology are necessary in addition to a high bulk modulus.A superhard material will have a high bulk modulus and a three-dimensional isotropic structure with ﬁxed atomic positions and covalent or partially covalent ionic bonds.Hardness also depends strongly on plastic deformation,which is related to the creation and motion of dislocations.This is not controlled by the shear modulus but by the shear strength τ,which varies as much as a factor of 10for different materials with similar shear moduli.It has been theoretically shown that τ/G is of the order of 0.03–0.04for a face-centered cubic metal,0.02for a layer structure such as graphite,0.15for an ionic compound such as sodium chloride,and 0.25for a purely covalent material such as diamond (32).Detailed calculationsmustFigure 4Hardness as a function of the shear modulus for selected materials (r-,rutile-type;c,cubic).A n n u . R e v . M a t e r . R e s . 2001.31:1-23. D o w n l o a d e d f r o m a r j o u r n a l s .a n n u a l r e v i e w s .o r g b y C h i n e s e A c a d e m y o f S c i e n c e s - L i b r a r y o n 05/16/09. F o r p e r s o n a l u s e o n l y .。

新型太阳能电池学术研讨会墙报录用列表

32
仇鑫灿 1-Efficient, stable and flexible perovskite solar cells using two-step solution processed SnO2 layers as electron-transport-material.
33
宋晶
1-Enhancing Performance of PSCs Using Nb-Doped SnO2 Nanosheets as
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蒋玉荣 CH3NH3PbI3-xClx Plannar Hybrid Perovskite Solar Cells: Role of Ionic
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杨颖
1-Effect of_doping of_NaI monovalent cation halide on_the_structural， morphological, optical and optoelectronic properties of MAPbI3 perovskite.
47
冯昱霖 1-Interfacial Negative Capacitance in Planar Perovskite Solar Cells An Interpretation based on Band Theory.
48
尹学文 1-Inverted Perovskite Solar Cells with Efficient Mixed-Fullerene Derivative
ቤተ መጻሕፍቲ ባይዱ30

Yb∶Ca3(NbGa)5O12晶体的坩埚下降法生长及光学性能研究

第53卷第4期2024年4月人㊀工㊀晶㊀体㊀学㊀报JOURNAL OF SYNTHETIC CRYSTALS Vol.53㊀No.4April,2024YbʒCa 3(NbGa )5O 12晶体的坩埚下降法生长及光学性能研究赵㊀涛,艾㊀蕾,梁团结,钱慧宇,孙志刚,潘建国(宁波大学材料科学与化学工程学院,浙江省光电探测材料及器件重点实验室,宁波㊀315211)摘要:使用坩埚下降法成功生长出了镱离子掺杂钙铌镓石榴石晶体(YbʒCa 3(NbGa)5O 12)㊂通过XRD 测试分析了晶体的结构,该晶体为立方晶系,晶胞参数a =b =c =12.471Å㊂对该晶体进行了拉曼光谱㊁透过光谱㊁吸收和发射光谱㊁荧光寿命等测试,计算了该晶体的吸收截面㊁发射截面㊁增益截面等㊂研究了在空气中退火对该晶体吸收光谱㊁发射光谱㊁荧光寿命的影响,退火前在935nm 处吸收截面为1.82ˑ10-20cm 2,退火后降低为1.40ˑ10-20cm 2,退火前在1031nm 处的发射截面为0.56ˑ10-20cm 2,退火后降低为0.40ˑ10-20cm 2,退火前荧光衰减时间为1.42ms,退火后为1.32ms㊂结果表明,YbʒCa 3(NbGa)5O 12单晶在空气中退火会对晶体的激光性能造成不利影响㊂关键词:YbʒCa 3(NbGa)5O 12晶体;坩埚下降法;吸收光谱;发射光谱;荧光衰减;退火中图分类号:O782㊀㊀文献标志码:A ㊀㊀文章编号:1000-985X (2024)04-0620-07Growth and Optical Properties of YbʒCa 3(NbGa )5O 12Crystals by Bridgman MethodZHAO Tao ,AI Lei ,LIANG Tuanjie ,QIAN Huiyu ,SUN Zhigang ,PAN Jianguo(Key Laboratory of Photoelectric Detection Materials and Devices of Zhejiang Province,School of Materials Science and Chemical Engineering,Ningbo University,Ningbo 315211,China)Abstract :Ytterbium ion doped calcium niobium gallium garnet crystal (Yb ʒCa 3(NbGa)5O 12)was successfully grown by Bridgman method.The structure of the crystal was analyzed by XRD.The crystal is cubic crystal system,and the unit cell parameter a =b =c =12.471Å.The crystal was tested by Raman spectroscopy,transmission spectroscopy,absorption and emission spectroscopy,and fluorescence lifetime.The absorption cross section,emission cross section,and gain cross section of the crystal were calculated.The effects of annealing in air on the absorption spectrum,emission spectrum and fluorescence lifetime of the crystal were studied.The absorption cross section at 935nm before annealing is 1.82ˑ10-20cm 2,and it decreases to 1.40ˑ10-20cm 2after annealing.The emission cross section at 1031nm before annealing is 0.56ˑ10-20cm 2,and it decreases to 0.40ˑ10-20cm 2after annealing.The fluorescence decay time before annealing is 1.42ms,and it is 1.32ms after annealing.The results demonstrate that the annealing of YbʒCa 3(NbGa)5O 12single crystal in air will adversely affect the laser performance of the crystal.Key words :YbʒCa 3(NbGa)5O 12crystal;Bridgman method;absorption spectrum;emission spectrum;fluorescence decay;annealing㊀㊀㊀收稿日期:2023-12-08㊀㊀基金项目:国家自然科学基金(51832009,512302300)㊀㊀作者简介:赵㊀涛(1997 ),男,山西省人,硕士研究生㊂E-mail:1254983331@ ㊀㊀通信作者:孙志刚,博士,助理研究员㊂E-mail:sunzhigang@0㊀引㊀㊀言钙铌镓石榴石(CNGG)晶体是一类无序激光晶体,结构介于激光玻璃的无序结构和激光晶体的有序结构之间㊂无序结构的激光玻璃,是一类典型的非均匀加宽的激光增益介质,但玻璃具有长程无序结构,限制㊀第4期赵㊀涛等:YbʒCa3(NbGa)5O12晶体的坩埚下降法生长及光学性能研究621㊀了声子的平均自由程,导致其热学性能相对较差,限制了高效㊁高功率密度激光的获得[1]㊂而传统的激光晶体如钇铝石榴石(YAG)晶体,具有很好的热学性质,但长程有序的特点使其具有相对单一的激活离子取代位置,导致其配位单一,激活离子的光谱较窄[2]㊂无序的钙铌镓石榴石晶体兼具两者的优点,具有光谱的非均匀加宽特性和较高的热导率,使得其在激光领域中具有潜在的应用价值㊂NdʒCNGG晶体的具有较宽的吸收与发射光谱,Pan等[3]采用直拉法生长了无序的NdʒCNGG晶体,InGaAs LD泵浦的峰值吸收截面约为4.1ˑ10-20cm2,在808nm LD激发的发射荧光谱中,4F3/2ң4I11/2的半峰全宽(full width at half maximum, FWHM)为15nm,4F3/2ң4I13/2半峰全宽为27nm,在超快激光脉冲产生方面展示出巨大的潜力㊂目前,研究人员对NdʒCNGG晶体的连续波㊁调Q及锁模超短脉冲激光特性已做了大量㊁系统的研究[4-6]㊂20世纪90年代初,随着体积小㊁效率高㊁寿命长的LD泵浦源的出现,Yb3+作为激光基质激活离子的研究迅猛发展㊂Yb3+具有最简单的能级结构,与Nd3+相比,具有本征量子缺陷低,辐射量子效率高,能级寿命长,吸收和发射光谱宽等特点㊂特别是Yb3+的吸收峰位于900~1000nm,能与目前商用的InGaAs半导体激光二极管泵浦源有效耦合,并且不需要严格控制温度㊂YbʒCa3(NbGa)5O12晶体(YbʒCNGG)已有相关报道,可获得连续激光输出,并通过锁模和调Q获得脉冲激光输出[7-9],证明了YbʒCNGG在激光领域的潜在价值㊂目前报道的YbʒCNGG晶体都是使用提拉法生长,该晶体的坩埚下降法生长还没有报道㊂坩埚下降法生长晶体是在密闭环境中进行,能有效防止原料Ga2O3的挥发;此外,与提拉炉相比较,坩埚下降炉价格低廉,设备维护简单,使用坩埚下降法生长晶体能够极大地降低生产成本,因此YbʒCNGG晶体可能更适合使用坩埚下降法生长㊂本文成功使用坩埚下降法生长出较大尺寸的YbʒCNGG晶体,并开展了其光学性能研究㊂1㊀实㊀㊀验1.1㊀原料制备和晶体生长YbʒCNGG晶体在1450ħ左右一致熔融,但在高温下Ga2O3原料会挥发,因此本实验采用坩埚下降法,在密闭环境中生长该晶体㊂使用的原料为Yb2O3(纯99.99%),CaCO3(纯99.99%),Nb2O5(纯99.99%), Ga2O3(纯99.999%),采用Ca3Nb1.6875Ga3.1875O12成分配比,按照以下的化学反应式进行多晶料的合成㊂2.892CaCO3+0.813375Nb2O5+1.626375Ga2O3+0.054Yb2O3=0.964Ca3Nb1.6875Ga3.1875O12㊃0.036Yb3Ga5O12+2.892CO2(1)按上述配比称量原料,进行充分研磨,放入混料机中混合24h,再进行液压机压块,随后放入马弗炉进行第一次烧结,烧结温度1000ħ,保温10h;取出后再次研磨㊁压块,进行第二次烧结,烧结温度1250ħ,保温时间30h,得到YbʒCNGG的多晶料㊂将多晶料放进装有YAG[111]籽晶的铂金坩埚,放入坩埚下降炉中进行晶体生长㊂接种温度为1450ħ,下降速度8mm/d㊂晶体生长结束后,以20ħ/h左右的速率使炉温降至室温,以消除晶体生长过程中所产生的热应力㊂众所周知,激光晶体在高温环境中工作一段时间后,性能会有所降低㊂在高温㊁富氧或贫氧环境中工作一段时间后某些单晶会改变颜色,导致其光学吸收带发生变化,这种现象已经在硅酸铋[10]㊁铌酸盐[11-12]㊁磷酸盐[13]和碱金属钼酸盐[14-16]等氧化物中发现㊂因此,本文在空气中对YbʒCNGG晶体进行了热退火,以此来探究高温环境工作后晶体的光学性能变化㊂将加工好的一块晶片切成两块,其中一块放进马弗炉中,在空气氛围下进行退火,退火温度为1000ħ,保温时间10h㊂1.2㊀性能测试使用德国Bruker XRD D8Advance型X射线粉末衍射仪对YbʒCNGG晶体的粉末样品进行XRD测试,辐射源为Cu靶X射线管,工作电压和电流分别为40kV和40mA,扫描范围10ʎ~70ʎ,步幅为0.02ʎ㊂使用DXR3Raman Microscope光谱仪记录了晶体在295K下的拉曼光谱,激发源为532nm波长的激光㊂使用美国Lambda950型紫外可见近红外分光光度计测量了晶体的吸收和透过光谱㊂使用法国FL3-111型荧光光谱仪测试了晶体的发射光谱,激发源为980nm激光㊂采用英国FLS980荧光光谱仪测试了晶体的荧光衰减曲线,激发波长980nm,监测波长1031nm㊂622㊀研究论文人工晶体学报㊀㊀㊀㊀㊀㊀第53卷2㊀结果与讨论2.1㊀晶体生长图1(a)为采用坩埚下降法生长得到的YbʒCNGG晶体,晶体直径为25mm,接种后生长部分长度约为80mm,其中偏析层部分约为25mm㊂晶体呈现咖啡色,透明,内部有少量裂纹,晶体开裂与晶体自身性质以及生长工艺有关㊂图1(b)为加工后的YbʒCNGG晶片,晶片直径25mm,厚度为1mm,属于(111)晶面,晶片中横向裂纹是加工所致㊂图1㊀坩埚下降法生长的YbʒCNGG晶体Fig.1㊀YbʒCNGG crystals grown by Bridgman method2.2㊀XRD分析图2为YbʒCNGG晶体单晶部分和顶部偏析层部分的粉末XRD图谱,将单晶部分的XRD数据导入Jade 中,通过拟合得出该晶体是Ia3d空间群,属于立方晶系,晶胞参数a=b=c=12.471Å,α=β=γ=90ʎ,比已报道的CNGG晶体晶胞参数(12.51Å)略小[17],原因是掺杂的Yb3+半径小于被取代的Ca2+半径,导致晶体晶格收缩㊂通过Jade分析,顶部偏析层的杂质成分大部分是立方焦火成岩(Ca2Nb2O7),这与文献[18]中得出结论一致,原因是掺入Yb3+后,生成了镱镓石榴石(Yb3Ga5O12),导致Ca2+与Nb5+的过量,从而生成了不属于石榴石相的Ca2Nb2O7㊂2.3㊀拉曼光谱图3是室温下YbʒCNGG退火前后晶体样品的拉曼图谱对比,孤立金属氧四面体基团[MO4](M代表Ga 和Nb)在700~900cm-1存在对称伸缩振动,这些[MO4]基团是石榴石晶格的结构单元,M阳离子进入到石榴石结构的d位[19]㊂在700~900cm-1看到两个密集的振动峰C1和C3,分别是[GaO4]和[NbO4]基团群的对称伸缩振动造成的,C1和C2峰下降明显,C3和C4变化较小的可能原因是晶体中部分Ga3+挥发,改变了晶体的结构和振动特性,影响了振动模式的活性㊂Ga3+挥发会对晶体中[GaO4]基团的对称伸缩振动产生影响㊂通常情况下,Ga O键连接可能会中断或减弱,这种情况可能导致对称伸缩振动变弱,在拉曼光谱中可能会表现为C1和C2峰强度下降㊂C2和C4分别是C1和C3的伴峰,此处出现峰,则代表[GaO4]和[NbO4]附近出现阳离子空位,峰强度越高,则代表阳离子空位浓度越高㊂从图中可以看出,退火后C2和C4处都出现了微弱的伴峰,表明在退火后的晶体中,阳离子空位浓度增加了,主要原因是高温退火后晶体表面的Ga3+浓度降低,但是幅度较小[20]㊂2.4㊀透过和吸收光谱退火前后晶体样品的透过图谱如图4(a)所示,600~2500nm的整体透过率接近80%,说明晶体质量较高,退火后晶体颜色变化不明显㊂图4(b)是YbʒCNGG晶体的吸收截面图,吸收峰对应Yb3+的2F7/2(基态)ң2F5/2(激发态)跃迁㊂基态2F7/2和激发态2F5/2分别被晶体场劈裂为4个和3个Stark能级,从基态多重态的几个Stark能级到激发态多重态2F7/2(0㊁1㊁2㊁3)ң2F5/2(0ᶄ㊁1ᶄ㊁2ᶄ)的电子跃迁大多数是声子辅助的,从而产生了相当宽的谱带㊂晶体退火前在935nm处吸收截面为1.82ˑ10-20cm2,退火后为1.40ˑ10-20cm2;退火前在971nm处吸收截面为1.22ˑ10-20cm2,退火后为1.03ˑ10-20cm2,退火后吸收截面明显降低㊂此外,㊀第4期赵㊀涛等:YbʒCa 3(NbGa)5O 12晶体的坩埚下降法生长及光学性能研究623㊀从图4(c)和4(d)可以计算得出,晶体退火前在935nm 处FWHM 为47.46nm,退火后为44.60nm;退火前在971nm 处FWHM 为23.47nm,退火后为23.86nm㊂退火后在935nm 处的FWHM 比退火前小了2.86nm㊂图2㊀YbʒCNGG 晶体中部单晶部分及顶部偏析层部分的粉末XRD 图谱Fig.2㊀Powder XRD patterns of the middle single crystal part and the top segregation layer of YbʒCNGGcrystal 图3㊀室温下退火前后YbʒCNGG 晶体样品的拉曼图谱Fig.3㊀Raman spectra of YbʒCNGG crystal samples before and post annealing at roomtemperature图4㊀室温下退火前后YbʒCNGG 晶体样品的性能测试㊂(a)透过光谱;(b)吸收光谱;(c)退火前晶体样品吸收光谱的高斯拟合图;(d)退火后晶体样品吸收光谱的高斯拟合图Fig.4㊀Performance testing of YbʒCNGG crystal samples before and post annealing at room temperature.(a)Transmission spectra;(b)absorption spectra;(c)Gaussian fitting of absorption spectra of crystal sample before annealing;(d)Gaussian fitting of the absorption spectrum of crystal sample post annealing 2.5㊀发射光谱关于YbʒCNGG 晶体的发射截面σem (λ)计算,本文使用互易法(reciprocity method),用下列公式进行计算㊂624㊀研究论文人工晶体学报㊀㊀㊀㊀㊀㊀第53卷σem (λ)=σαbsZ l Z u exp E zl -hc λkT ()(2)式中:σabs 为吸收截面,h 为普朗克常数,k 为玻耳兹曼常数,c 为光速,λ为波长,T 为实验温度,Z l /Z u 为下㊁上能级的配分函数比,E zl 为零声子线㊂如图5(a)所示,计算得出退火前975nm 处的发射截面为1.28ˑ10-20cm 2,退火后为1.11ˑ10-20cm 2,退火前1031nm 处的发射截面为0.56ˑ10-20cm 2,退火后为0.40ˑ10-20cm 2㊂退火后975㊁1031nm 处的发射截面均低于退火前㊂图5(b)是在980nm 激光激发下得到的发射光谱,发射峰位于1031nm 处,在相同测试条件下,退火后该晶体的发射强度明显低于退火前,这与计算得出的结果相一致,表明YbʒCNGG 晶体在空气中退火后,对其激光性能有不利影响㊂原因是空气中的高温退火可能会对材料的物理和化学性质产生影响,包括晶格结构的变化和缺陷的生成㊂退火过程中晶格结构的变化和缺陷的形成可能对透过谱和发射谱性能产生影响㊂晶格结构变化:高温退火可能引起晶格结构的重新排列㊂在退火过程中,原子或分子在晶体中重新定位以达到更低的能量状态㊂这可能导致晶格略微变化,晶格参数可能发生微小的变化,如晶胞参数㊁晶体取向等㊂这种微小的结构变化可能会影响透过谱和发射谱的特性㊂缺陷的生成:高温退火也可能导致缺陷的生成㊂例如,点缺陷(Ga 3+的挥发)㊁位错或晶界等缺陷的产生㊂这些缺陷可能导致电子状态的变化㊁局部晶格畸变或者在晶体中引入能级㊂这些缺陷可能会影响材料的光学性质,包括透过谱和发射谱㊂图5㊀室温下退火前后YbʒCNGG 晶体样品的发射截面曲线(a)和980nm 激光激发下得到的发射光谱(b)Fig.5㊀Emission cross-section curves (a)and emission spectra at 980nm excitation (b)of YbʒCNGG crystal samples before and post annealing at room temperature2.6㊀增益截面根据上述吸收和发射截面光谱,增益截面σg (λ)可由下式计算:σg (λ)=βσem (λ)-(1-β)σabs (λ)(3)式中:β为激发态离子反转分数㊂图6所示为退火前后的YbʒCNGG 晶体样品在不同β值(0,0.25,0.50,0.75,1.00)下的增益截面曲线㊂如图6(a)所示,在1010~1040nm 处,当布居反转分数达到25%时,增益截面变为正值㊂如此低的反转比例意味着1031nm 波长的YbʒCNGG 激光器将具有较低的泵浦阈值,这表明YbʒCNGG 晶体是1031nm 激光器的理想候选材料㊂在高抽运情况下,增益截面谱也较宽,表现出良好的可协调性㊂而退火后该晶体增益截面曲线如图6(b)所示,并且在布居反转比例达到50%时,在1031nm 附近的增益带宽明显低于退火前,因此理论上通过被动锁模达到最小脉冲也将会受到影响[21],也就是说,在高温下工作会对该晶体超快激光的产生造成不利影响㊂2.7㊀荧光衰减室温下对退火前后的YbʒCNGG 晶体样品进行荧光衰减测试㊂如图7所示,激发波长980nm,监测波长1031nm,采用单指数函数拟合,如公式(4)所示㊂y =A 1e -x t +y 0(4)㊀第4期赵㊀涛等:YbʒCa3(NbGa)5O12晶体的坩埚下降法生长及光学性能研究625㊀式中:A1为前因子,y0为初始强度,t为时间,x㊁y为测试的横纵坐标,对应波长㊁强度㊂通过拟合得到退火前的荧光衰减时间为1.42ms,退火后的荧光衰减时间为1.32ms,观察到退火后Yb3+的寿命减少,表明这种退火在晶体中引入了进一步的缺陷,很可能是由表面Ga3+的挥发造成的,与文献中采用提拉法生长的YbʒCNGG晶体τ=816μs相比较,结果相差很大,可能是该晶体有很强的重吸收,造成直接测量荧光寿命不准确,但是与文献中退火后Yb3+的寿命会减少的结论是一致的[20]㊂图6㊀室温下退火前后YbʒCNGG晶体样品增益截面曲线Fig.6㊀YbʒCNGG crystal samples gain cross-section curves before and post annealing at room temperature图7㊀室温下退火前后YbʒCNGG晶体样品荧光衰减曲线Fig.7㊀YbʒCNGG crystal samples fluorescence decay curves before and post annealing at room temperature3㊀结㊀㊀论采用坩埚下降法,生长出尺寸为ϕ25mmˑ80mm的YbʒCNGG透明单晶,通过XRD粉末衍射,得出了偏析层的主要杂质成分为Ca2Nb2O7㊂通过透过和吸收光谱得出该晶体退火前在935和971nm处有很宽的吸收带宽,分别为47.46和23.47nm,退火后935nm处吸收带宽变窄㊂尽管常规情况下退火有助于提高晶体的均匀性和激光性能,但在本文中通过对YbʒCNGG晶体退火前后晶体发射截面和增益截面的计算,以及发射光谱和荧光衰减的测量,发现采用高温退火可能会引入缺陷并导致激光性能下降㊂这可能暗示着退火温度需要重新评估或者退火周期需要调整以更好地维持晶体性能,后续本团队会继续研究不同退火条件对YbʒCNGG晶体激光性能的影响㊂参考文献[1]㊀于浩海,潘忠奔,张怀金,等.无序激光晶体及其超快激光研究进展[J].人工晶体学报,2021,50(4):648-668+583.YU H H,PAN Z B,ZHANG H J,et al.Development of disordered laser crystals and their ultrafast lasers[J].Journal of Synthetic Crystals, 2021,50(4):648-668+583(in Chinese).[2]㊀KANCHANAVALEERAT E,COCHET-MUCHY D,KOKTA M,et al.Crystal growth of high doped NdʒYAG[J].Optical Materials,2004,26626㊀研究论文人工晶体学报㊀㊀㊀㊀㊀㊀第53卷(4):337-341.[3]㊀PAN H,PAN Z B,CHU H W,et al.GaAs Q-switched NdʒCNGG lasers:operating at4F3/2ң2I11/2and4F3/2ң2I13/2transitions[J].OpticsExpress,2019,27(11):15426-15432.[4]㊀SHI Z B,FANG X,ZHANG H J,et al.Continuous-wave laser operation at1.33μm of NdʒCNGG and NdʒCLNGG crystals[J].Laser 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退火温度对纯钛TA1_织构及各向异性的影响

第50卷第4期中南大学学报(自然科学版) V ol.50No.4 2019年4月Journal of Central South University (Science and Technology)Apr. 2019 DOI: 10.11817/j.issn.1672−7207.2019.04.007退火温度对纯钛TA1织构及各向异性的影响张贵华，江海涛，吴波，杨永刚，田世伟，郭文启(北京科技大学工程技术研究院，北京，100083)摘要：通过X线衍射(XRD)和电子背散射衍射(EBSD)等分析技术，研究退火温度对冷轧态TA1钛板显微组织及织构的影响规律。

研究结果表明：TA1钛板冷轧退火后，微观组织发生再结晶并形成典型的双峰分裂基面织构特征。

在退火温度不大于700 ℃时，组织变化主要以回复与再结晶的形核生长为主，生成>011(和)3<0231 )22111(类型再结晶织构组分，此时轧制织构组分逐渐消失；当退火温度达到800 ℃时，晶粒变化以合并1><00长大为主，再结晶织构组分>1)2(的强度也继续增强。

同时，织构组分对板材的各<0011213(和>1<001132向异性有着直接影响，由于棱锥型织构>11)2<00112(再结晶织构组分特征的作用，可开动3(和>1<0011)32的滑移系统分别为易激活的柱面<a>滑移和较难开动的基面<a>滑移或棱锥面<c+a>滑移，从而导致板面内TD方向的拉伸强度比RD方向的拉伸强度大，而45°方向强度最低，从而产生较大的板面各向异性。

关键词：TA1钛板；织构；退火；再结晶；各向异性；电子背散射衍射(EBSD)中图分类号：TG146.23 文献标志码：A 文章编号：1672−7207(2019)04−0806−08 Effect of annealing temperature on texture and anisotropy ofmechanical properties of pure titanium(TA1) sheetZHANG Guihua, JIANG Haitao, WU Bo, YANG Yonggang, TIAN Shiwei, GUO Wenqi (Institute of Engineering Technology, University of Science and Technology Beijing, Beijing 100083, China)Abstract: The effect of evolution of microstructure and texture of commercially pure titanium (TA1) annealed at different temperatures was investigated by X-ray diffraction (XRD), and electron backscattered diffraction (EBSD). The results show that recovery and recrystallization of the cold rolled TA1 titanium sheet occur during the annealing process, and typical TD-split basal texture was formed. When the annealing temperature is below 700 ℃, the microstructure is characterized by recovery and recrystallization, and recrystallization texture components are presented. The as-rolled texture component is gradually weakened and disappears with the increase of the heat treatment temperature. When the annealing temperature reaches 800 ℃, the grain growth is dominated by merged-growth and the intensity of11)2(recrystallized texture component continue to increase. In addition, anisotropy11<00<03(and>1>011)32of mechanical properties of TA1 sheet is related to the texture components. Due to pyramid textures>011(3<0312 and>11(recrystallization textures, the cylinder <a> slip is respectively easier to be activated and the base <00)2211<a> slip or pyramidal plane <c+a> slip becomes more difficult to be activated respectively, which leads to greater tensilestrength in the TD direction than the RD direction of the sheet. As a result, the anisotropy of mechanical properties of TA1 sheet is caused.Key words: TA1 titanium sheet; texture; annealing; recrystallization; anisotropy; electron backscattered diffraction (EBSD)收稿日期：2018−05−15；修回日期：2018−08−27基金项目(Foundation item)：国家重点研发计划项目(2016YFB0101605) (Project(2016YFB0101605) supported by the National Key Research and Development Program of China)通信作者：江海涛，博士，教授，从事金属材料方面研究；E-mail：****************.cn第4期张贵华，等：退火温度对纯钛TA1织构及各向异性的影响807工业纯钛在航空航天、舰船、核能等高科技领域均有广泛的用途[1−4]，在实际的应用中，除了固有的腐蚀性能外，其机械性能也是设计的重要标准。

Growth and applications of group III-nitrides

J.Phys.D:Appl.Phys.31(1998)2653–2710.Printed in the UK PII:S0022-3727(98)68952-XREVIEW ARTICLEGrowth and applications ofGroup III-nitridesO AmbacherWalter Schottky Institute,Technical University Munich,Am Coulombwall,D-85748Garching,GermanyReceived18February1997,inﬁnal form15June1998Abstract.Recent research results pertaining to InN,GaN and AlN are reviewed,focusing on the different growth techniques of Group III-nitride crystals andepitaxialﬁlms,heterostructures and devices.The chemical and thermal stability ofepitaxial nitrideﬁlms is discussed in relation to the problems of depositionprocesses and the advantages for applications in high-power and high-temperaturedevices.The development of growth methods like metalorganic chemical vapourdeposition and plasma-induced molecular beam epitaxy has resulted in remarkableimprovements in the structural,optical and electrical properties.New developmentsin precursor chemistry,plasma-based nitrogen sources,substrates,the growth ofnucleation layers and selective growth are covered.Deposition conditions andmethods used to grow alloys for optical bandgap and lattice engineering areintroduced.The review is concluded with a description of recent Group III-nitridesemiconductor devices such as bright blue and white light-emitting diodes,theﬁrstblue-emitting laser,high-power transistors,and a discussion of further applicationsin surface acoustic wave devices and sensors.1.IntroductionGroup III-nitrides have been considered a promising system for semiconductor devices applications since1970, especially for the development of blue-and UV-light-emitting diodes.The III–V nitrides,aluminium nitride (AlN),gallium nitride(GaN)and indium nitride(InN), are candidate materials for optoelectrical applications at such photon energies,because they form a continuous alloy system(InGaN,InAlN,and AlGaN)whose direct optical bandgaps for the hexagonal wurtzite phase range from 1.9eV forα-InN and 3.4eV forα-GaN to 6.2eV forα-AlN.The cubic modiﬁcations have bandgaps in the range from 1.7eV forβ-InN and 3.2eV for β-GaN to 4.9eV forβ-AlN(ﬁgures1and2)[1–6]. Other advantageous properties include high mechanical and thermal stability,large piezoelectric constants and the possibility of passivation by forming thin layers of Ga2O3 or Al2O3with bandgaps of approximately 4.2eV and 9eV.The spontaneous and piezoelectric polarization(in the wurtzite materials)and the high electron drift velocities (2×105m s−1[7])of GaN can be used to fabricate high-power transistors based on AlGaN/GaN heterostructures. In addition,AlN is an important material with a variety of applications such as passive barrier layers,high-frequency acoustic wave devices,high-temperature windows,and dielectric optical enhancement layers in magneto-optic multilayer structures[8,9].Very informative reviews of the growth techniques and structural,optical and electrical properties of Group III-nitrides and their alloys have been presented by Strite et al[10,11].A good overview of applications of Group III-nitride based heterostructures for UV emitters and high-temperature,high-power electronic devices is provided in[12]and[13].This review focuses on the development of the different growth techniques successfully applied to the deposition of Group III-nitride epitaxial ﬁlms and heterostructures,such as chemical transport and metalorganic chemical vapour deposition(MOCVD), sputtering and molecular beam epitaxy(MBE).The quality of state-of-the-art material and its application for optical and electronic devices are discussed in detail in order to point out possible limitations,promising developments and future trends.Theﬁrst systematic effort to grow InN,GaN and AlN by chemical vapour deposition or sputtering processes took place in the1970s in order to characterize the optical and structural properties of thinﬁlms.At that time, neither metalorganic precursors containing In or Al with electronic grade purity,plasma sources for nitrogen radicals compatible with MBE systems,nor substrate material with reasonably good thermal and lattice matches to the nitrides were available.The InN and GaN material had large concentrations of free electrons,presumed to result from oxygen impurities and intrinsic defects,and the structural quality of the AlNﬁlms was not good enough for optical0022-3727/98/202653+58$19.50c 1998IOP Publishing Ltd2653OAmbacherFigure 1.Bandgap and bowing parameters of hexagonal (α-phase)and cubic (β-phase)InN,GaN,AlN and their alloys versus lattice constant a 0[1–6].or electronic applications.Primarily,the development of MOCVD and plasma-induced molecular beam epitaxy (PIMBE)over the last eight years has led to a number of recent advances and important improvements in structural properties.2.Crystal structure,polarity and polarization of InN,GaN and AlNIn contrast to cubic III–V semiconductors like GaAs and InP with the zincblende structure,the thermodynamically stable phase of InN,GaN and AlN,is the hexagonal wurtzite structure (α-phase).Beside the α-phase,a metastable β-phase with zincblende structure exists and a cubic high-pressure modiﬁcation with NiAs structure was observed for pressures above 25kbar in the case of AlN [14].Because the α-and β-phases of Group III-nitrides only differ in the stacking sequence of nitrogen and metal atoms (polytypes),the coexistence of hexagonal and cubic phases is possible in epitaxial layers,for example due to stacking faults.The hexagonal crystal structure of Group III-nitrides can be described by the edge length a 0of the basal hexagon,the height c 0of the hexagonal prism and an internal parameter u deﬁned as the anion–cation bond length along the (0001)axis.Because of the differentcations and ionic radii (Al 3+:0.39˚A,Ga 3+:0.47˚A,In 3+:0.79˚A[15]),InN,GaN and AlN have different lattice constants,bandgaps and binding energies as shown in table 1[16,17].Both wurtzite and zincblende structures have polar axes (lack of inversion symmetry).In particular,the bondsinFigure 2.Experimental results of bandgaps of hexagonal Group III-nitrides versus lattice constant c 0at room temperature [1–6].Table ttice constants,bandgaps and binding energies of hexagonal InN,GaN and AlN.Wurtzite,300K AlN GaN InN a 0(˚A)b 3.112 3.189 3.54c 0(˚A)b 4.982 5.185 5.705c 0/a 0(exp.)b 1.6010 1.6259 1.6116c 0/a 0(calc.)a 1.6190 1.6336 1.6270u 0a0.3800.3760.377a Bohr (˚A)a 5.814 6.04 6.66E B (M–N)c (eV)b2.882.201.98a From [16].b From [17].cM =In,Ga or Al,N =nitride.the 0001 direction for wurtzite and 111 direction for zincblende are all faced by nitrogen in the same direction and by the cation in the opposite direction.Both bulk and surface properties can depend signiﬁcantly on whether the surface is faced by nitrogen or metal atoms [18,19].The most common growth direction of hexagonal GaN is normal to the {0001}basal plane,where the atoms are arranged in bilayers consisting of two closely spaced hexagonal layers,one with cations and the other with anions,so that the bilayers have polar faces.Thus,in the case of GaN a basal surface should be either Ga-or N-faced.By Ga-faced we mean Ga on the top position of the {0001}bilayer,corresponding to [0001]polarity (ﬁgure 3).Ga-faced does not mean Ga-terminated;termination should only be used2654Growth and applications of GroupIII-nitridesFigure3.Different polarities(Ga-and N-faced)of wurtzite GaN.to describe a surface property.A Ga-face surface mightbe N-terminated if it is covered with nitrogen atoms,butwithoutﬂipping the crystal it will never be N-faced.Itis,however,important to note that the(0001)and(000¯1) surfaces of GaN are inequivalent(by convention,the[0001]direction is given by a vector pointing from a Ga atom toa nearest-neighbour N atom).It has been reported that high-quality epitaxial GaNﬁlms deposited by MOCVD on c-plane sapphire substratesgrow in the(0001)direction with Ga-faced surfaces,whileMBE growth commonly occurs in the(000¯1)direction, yielding an N-facedﬁlm[20–22].Polar faces are known to have very marked effectson growth in binary cubic semiconductors.For example,growth along the Ga-faced{111}direction of GaAs isknown to be slow and has the tendency to produce planarsurfaces,whereas growth of the As-face is fast and rough[23].Similarly,Ponce et al found that the smooth sideof bulk single crystal platelets corresponds to the Ga-face(0001)whereas the N-face(000¯1)is much rougher[20].In the following we will discuss the inﬂuence ofspontaneous and piezoelectric polarization on the physicalproperties of Group III-nitrides.This class of polarization-related properties is obviously important for devices(section9)because the electricﬁelds inﬂuence the shapeof the band edges and the carrier distribution insidenitride-based heterostructures.Therefore spontaneousand piezoelectric polarization can inﬂuence the radiativerecombination in light-emitting devices as well as theelectrical properties of the transistor structures discussedin detail later.Wurtzite is the structure with highest symmetrycompatible with the existence of spontaneous polarization[16,24,25]and the piezoelectric tensor of wurtzite hasthree independent nonvanishing components.Therefore,polarization in these material systems will have both aspontaneous and a piezoelectric component.Becauseof the sensitive dependence of spontaneous polarizationon the structural parameters,there are some quantitativedifferences in polarization of the three nitrides studied here.The increasing nonideality of the crystal structure going Table2.Spontaneous polarization,piezoelectric and dielectric constants of AlN,GaN and InN.Wurtzite AlN GaN InNP SP(C m−2)−0.081−0.029−0.032e33(C m−2) 1.46a0.73a0.97a1.55b1c0.65d0.44ee31(C m−2)−0.60a−0.49a−0.57a−0.58b−0.36c−0.33d−0.22ee15(C m−2)−0.48b−0.3c−0.33d−0.22eε119.0b9.5fε3310.7b10.4fa From[16].b From[26].c From[27].d From[28].e From[29].f From[30].from GaN to InN to AlN(u0increases and c0/a0decreases (table1)),corresponds to an increase in spontaneous polarization.In the absence of external electricﬁelds,the total macroscopic polarization P of a solid is the sum of the spontaneous polarization P SP in the equilibrium lattice and the strain-induced or piezoelectric polarization P P E.Here we consider polarizations along the(0001)axis, because this is the direction along which standard bulk materials,epitaxialﬁlms and heterostructures are grown. Spontaneous polarization along the c-axis is P SP=P SP z (the direction of spontaneous polarization is determined by the polarity;the direction of the piezoelectric polarization depends on the polarity and whether the material is under tensile or compressive stress)and piezoelectric polarization can be calculated by using the piezoelectric coefﬁcients e33 and e13(table2)asP P E=e33εz+e31(εx+εy)(1) where a0and c0are the equilibrium values of the lattice parameters,εz=(c−c0)/c0is the strain along the c-axis, and the in-plane strainεx=εy=(a−a0)/a0is assumed to be isotropic.The third independent component of the piezoelectric tensor,e15,is related to the polarization induced by shear strain and will not be discussed.To give an example of the possible inﬂuence of polarization on the physical properties of nitride-based heterostructures,we calculate the electricﬁeld caused by polarization inside a Ga-faced Al x Ga1−x N/GaN/Al x Ga1−x N quantum well.We assume that the GaN is grown pseudomorphically on the AlGaN(a(GaN)=a(AlGaN))2655O AmbacherTable3.Experimental and calculated values of the piezoelectric constants and bulk modulus for wurtzite and zincblende Group III-nitrides.AlN GaN InNGPawurtzite exp.a calc.b exp.c calc.b exp.d calc.bC11345396374367190223C12125137106135104115C131201087010312192C33395373379405182224C44118116101951048B201207180202139141zincblende calc.e calc.b calc.e calc.b calc.e calc.bC11304304296293184187C12152160154159116125C4419919320615517786a From[31].b From[32].c From[33].d From[34].e From[35].and that screening effects due to free carriers and surface states can be neglected.The lattice constants a and c of the GaN layer are decreased and increased respectively,due to the biaxial compressive stress which becomes larger with increasing Al content of the AlGaNﬁlm.The relation between the lattice constants of the hexagonal GaN is given byc−c0 c0=−2C13C33a−a0a0(2)where C13and C33are elastic constants(table3). Using equations(1)and(2)the amount of piezoelectric polarization in the direction of the c-axis can be determined byP P E=2a−a0a0e31−e33C13C33.(3)The strain of the pseudomorphically grown GaN can be calculated using Vegard’s law(linear interpolation of the lattice constants of relaxed Al x Ga1−x N from the values for GaN and AlN:a(x)=(−0.077x+3.189)˚A(table1)), leading toP P E(GaN)=0.0163x C m−2(4) and a total polarization ofP(GaN)=P SP(GaN)+P P E(GaN)=(−0.029+0.0163x)C m−2.(5) The polarization generates an electricﬁeld E(GaN)inside the GaN layer:E(GaN)=−P(GaN)ε(GaN)ε0=(3.6×106−2.1×106x)V cm−1(6) whereε(GaN)(table2)andε0are the dielectric constants of GaN andvacuum.Figure4.Polarization(spontaneous,piezoelectric and total polarization)of a relaxed Al x Ga1−x N and a pseudomorphic on top of Al x Ga1−x N grown GaN layer versus Al content x. The interface chargeσis caused by the different total polarizations of the GaN and the AlGaNﬁlm.The AlGaN is assumed free of strain and therefore the piezoelectric polarization equals zero.The total polarization of the AlGaN can be described by a linear approximation between the spontaneous polarization of GaN and AlN:P(AlGaN)=P SP(AlGaN)=(−0.029−0.052x)C m−2.(7)A charge density at the GaN/AlGaN interfaces,σ(GaN/ AlGaN),is caused by the different polarizations of GaN and AlGaN:±σ(GaN/AlGaN)=P(GaN)−P(AlGaN)=±0.068x C m−2.(8) The spontaneous polarization,piezoelectric polarization and interface charge density of GaN embedded in two Al0.15Ga0.85N layers are determined to be−0.029,0.0025 and±0.0025C m−2respectively.(For AlGaN/GaN/AlGaN heterostructures with different Al content x,seeﬁgure4.) The electricﬁeld caused by polarization effects can reach a strength of3×106V cm−1.The modiﬁcation of the band edges due to spontaneous polarization and piezoelectricﬁelds inside the GaN layer can have a signiﬁcant inﬂuence on the optical properties (ﬁgure5).Due to the Stark and Franz–Keldysh effects, the effective bandgap of GaN will be red-shifted and the recombination probability of electron hole pairs will be decreased because of the spatial separation of electrons and holes[36,37].These physical effects thus change the energy of the electroluminescence out of GaN or InGaN2656Growth and applications of GroupIII-nitridesFigure5.Conduction and valence band edges of a pseudomorphic grown AlGaN/GaN/AlGaN(x=0.15)andGaN/InGaN/GaN(x=0.06)heterostructure.The arrows indicate schematically the radiative recombination of an electron and a hole,which is red-shifted in comparison to the bandgap energy due to the Stark effect.quantum wells and the recombination rates of carriers inside a Group III-nitride based laser structure(section9.5).The strong electricﬁelds can also enhance electron or hole accumulation at AlGaN/GaN interfaces(ﬁgure5).This effect can be used in heterostructureﬁeld effect transistors, as discussed later in section9.3.At which interface (lower or upper)of a AlGaN/GaN/AlGaN heterostructure electrons or holes are conﬁned depends on the polarity of the material.In respect of polarization effects,the Group III-nitrides exhibit unusual properties.The piezoelectric constants have the same sign as in II–VI compounds,and opposite to those of III–V compounds.The absolute values of the piezoelectric constants are up to ten times larger than in conventional III–V and II–VI compounds.In particular the constants e33and e31of AlN are larger than those of ZnO and BeO[38],and are therefore the largest known so far among the tetrahedrally bonded semiconductors.The spontaneous polarization(the polarization at zero strain) is also very large in the nitrides.That of AlN is only about three toﬁve times smaller than in typical ferroelectric perovskites[39].For these reasons,the spontaneous and piezoelectric polarization of hexagonal Group III-nitrides can have a much larger inﬂuence on the electrical and optical properties of devices than in other III–V compounds. Finally it should be mentioned that free carriers with a concentration above1018cm−3,charged defects or compensation of surface charges by adsorbates can reduce the polarization-induced electricﬁelds and have to be considered in a detailed analysis of polarization-related effects.3.Thermal properties and stabilityThe primary methods of obtaining crystal material rely on growing epitaxial layers on different substrates at high temperatures.Unfortunately,the different coefﬁcients of thermal expansion between substrate and nitride introduce residual stress upon cooling.These induced stresses can cause additional structural defects and piezoelectricﬁelds and will inﬂuence the optical and electrical properties of ﬁlms and devices.The determination of thermal expansivity is not only related to other thermal properties(thermal conductivity, speciﬁc heat)but can also yield parameters pertinent to2657OAmbacherFigure ttice constant and c /a ratio versus temperature.other basic properties,like the temperature dependence of the band gap [2].The value of the thermal expansion coefﬁcient depends on many parameters,such as defect concentration,free carriers,and strains,and the published values are somewhat scattered.The thermal expansivities perpendicular and parallel to the c -axis in hexagonal material are usually different.The lattice parameters and the thermal expansion coefﬁcient have been measured to intermediate temperatures for AlN,GaN and once for InN [30,40–46].The increase of the lattice constant a and the thermal expansion coefﬁcients of hexagonal InN,GaN and AlN with increasing temperature measured by different groups are shown for comparison in ﬁgures 6and 7.As the lattice constants a and c increase,the c/a ratio of the lattice constants becomes smaller with increasing temperature.The experimental data of the lattice constants and thermal expansion coefﬁcients for AlN and GaN are in good agreement with the theoretical calculations of Wang and Reeber (ﬁgure 7)[47].The calculated thermal expansion coefﬁcients of AlN,GaN and InN at 100K are 1.3×10−8K −1,1.2×10−6K −1and 2.4×10−6K −1for a 0and −5×10−8K −1,1.1×10−6K −1and 2.8×10−6K −1for c 0.At 600K,these values become 5.3×10−6K −1,5×10−6K −1and 5.7×10−6K −1for a 0and 4.4×10−6K −1,4.4×10−6K −1and 3.7×10−6K −1for c 0.Below 100K the thermal expansion coefﬁcient of AlN was calculated to be negative.Above 600K up to the decomposition temperature (discussed below),the thermal expansion coefﬁcients gradually increase by up to 25%.The lattice constants,binding energy and decomposi-tion temperature of Group III-nitrides have important con-sequences for the thermal stability of nitride-baseddevices.Figure 7.Thermal expansion coefﬁcients parallel (α(c ))and perpendicular (α(a ))to the c -axis versus temperature.The relatively large range of uncertainty and the limited number of experimental data concerning nitrogen diffusion,the temperature dependence of the nitrogen ﬂux from a ni-tride crystal surface and the nitrogen pressures necessary to stabilize a GaN melt are due in part to the very high melt-ing points T M and N 2equilibrium pressures of the Group III-nitrides.Recently,the thermal stability of InN was investigated at N 2pressures extending up to 18.5kbar [48].It was shown by differential thermal analysis that,over the whole investigated pressure range (0.1–18.5kbar),rapid decomposition of InN occurs above (710±10)◦C.Trainor and Rose [49]observed dissociation of thin InN ﬁlms at 500◦C and 1bar N 2.Guo and Kato [50]observed a change of the reﬂective high-energy electron diffraction (RHEED)pattern of InN single crystal ﬁlms when the temperature was raised above 550◦C.They concluded that the crystals decomposed to In and N 2above that temperature.MacChesney et al observed the equilibrium partial pressure of N 2over InN to be 1bar at 800K,increasing exponentially with 1/T to 105bar at 1100K [51](ﬁgure 8).The ﬁrst study of the thermal stability of GaN was made by Johnson et al [52].More recently Sime and Margrave [53]investigated the evaporation of GaN and Ga metal in the temperature range 900to 1150◦C at 1atm pressure of N 2,NH 3and H 2,while studying the formation and decomposition equilibria.They determined the heat of evaporation and proposed the existence of [GaN]x polymers in the gas phase.Thurmond and Logan [54]chemically measured the partial pressure ratios which2658Growth and applications of GroupIII-nitridesFigure8.Equilibrium N2pressure over the MN(s)+M(l) systems(M=In,Ga or Al),and melting points T M from high-pressure experiments and theoretical calculations [48–60].exist in a(H2+NH3)gas mixture streaming over Ga and GaN.They determined the equilibria both in the case of formation and decomposition of GaN.Lorenz and Binkowski[55]observed the decomposition at given temperatures by measuring the time dependence of the increase in N2pressure.An extensive study on the thermal stability of GaN at high temperatures and pressures up to60kbar has been performed by Karpinsky et al[56], using a gas pressure technique and a tungsten carbide anvil cell.In the high-pressure range,the p(1/T)curve strongly deviates from the linear dependence proposed by Thurmond and Logan(ﬁgure8),but there is good agreement in the Gibbs free enthalpy with G0=(32.43T−3.77×104)cal/mol for GaN synthesis.The value of enthalpy H0of−37.7kcal/mol is in good agreement with the value estimated by Madar et al[57].Resistively heated graphiteﬁlaments were used by Class[58]to evaluate the melting behaviour and temperature of AlN.The melting of AlN was observed at 2750–2850◦C,at nitrogen pressures of100and200bar. Slack and McNelly[59]calculated N2pressures over AlN in equilibrium with liquid Al to be1,10and100bar at 2563◦C,2815◦C and3117◦C respectively.The melting temperatures for different nitrides were evaluated by Van Vechten[60]with the use of a semi-empirical theory for electronegativity,concluding that the melting point of AlN is close to3487K.Figure8summarizes the results of the theoretical calculations and experiments described above, including the meltingtemperatures.Figure9.The partial pressure of mass28amu(N+2,CO+, C2H+4)versus effusion or decomposition temperature T E. The decomposition of InN,GaN and AlN is observed above 630◦C,850◦C and1040◦C respectively[61].To quantitatively determine and compare the thermal stability of thinﬁlms,the desorption and decomposition of polycrystalline InN deposited at550◦C,and epitaxial GaN and AlN grown at950◦C and1050◦C were measured by heating the samples in vacuum and recording the partial pressure of relevant gases using a quadrupole mass spectrometer[61].For a known heating routine,the desorption spectra can be analysed toﬁnd the binding energies of various desorbed species as well as the thermal stability of the sample[62].Figure9shows the partialpressure of mass28amu(N+2,CO+,C2H+4)versus effusion or decomposition temperature T E.The nitrogen partial pressure increases exponentially above T E=630◦C, 850◦C and1040◦C for InN,GaN and AlN respectively, illustrating that the decomposition temperature in vacuum is much lower than the melting point.To determine the effective decomposition activation energy more precisely, the nitrogenﬂux was calculated from the measured nitrogen pressure.The rate of nitrogen evolution (N)is equal to the rate of decomposition,and the slope of ln[ (N)] versus1/T(ﬁgure10)gives the effective activation energy of the decomposition in vacuum E MN.The decomposition rate equals the desorption of one monolayer every second( (N)=1.5×1015cm−2s−1)at795,970 and1050◦C,and the activation energy of the thermally induced decomposition is determined to be E MN=3.5eV (336kJ mol−1), 3.9eV(379kJ mol−1)and 4.3eV (414kJ mol−1)for InN,GaN and AlN respectively(table4). This indicates temperature limits for high-temperature or high-power devices.2659OAmbacherFigure 10.Nitrogen ﬂux or desorption rate for InN,GaN and AlN in the temperature range of decomposition.The rate of N evolution is equal to the rate of decomposition and the slope of ln[ (N)]versus 1/T gives the effective activation energy E MN of the decomposition in vacuum [61].Table 4.Density ρ,melting point T M ,decomposition temperature T E and activation energy E MN of thedecomposition of InN,GaN and AlN (p <10−6mbar).ρaT M a T E b E MN b(g cm −3)(K)(◦C)(kJ mol −1)AlN 3.2934871040414GaN 6.072791850379InN6.812146630336a From [17].bFrom [61].In connection with the thermal stability of hexagonal nitrides,the thermally activated nitrogen self-diffusion should be mentioned.Nitrogen diffusion is a fundamental transport process in Group III-nitrides.In general,self-diffusion processes in wide-bandgap semiconductors or insulators are more complex than in metals [63].This is due to the large variety of native defects,different possible charge states of defects,and to the much larger effects of small concentrations of defects on the Fermi level position [64].Diffusion processes also play an important role in device fabrication and the thermal stability of high-power devices [65].Diffusion of dopants is utilized to engineer p–n junctions and transistor structures [66].In other cases,diffusion can be destructive to delicate structures due to the transport of dopants from thin layers into the adjacent layers or thesubstrate.Figure 11.Arrhenius plot of the measured nitrogen self-diffusion coefﬁcient in hexagonal GaN,obtained at temperatures between 770and 970◦C.The diffusion coefﬁcients were calculated from the concentration depth proﬁles determined from SIMS signals of 14N (opensquares),15N (open circles),14N 2(full squares)and 15N 2(full circles).The line represents a least-squares ﬁt to the measured data.In III–V compounds,diffusion measurements are difﬁcult to perform because of the high partial vapour pressure of the Group V elements and the dependence of native defect species and concentrations on stoichiometry.Goldstein [67]and Palfrey et al [68]diffused radioactive 72Ga into bulk GaAs to study Ga self-diffusion.The measurement of the depth proﬁle of 72Ga was realized by mechanical sectioning and determination of the radioactivity.Within the temperature range investigated,the authors reported activation enthalpies varying between 2.6and 5.6eV.Wang et al [69]grew and studied an isotopically tagged 69GaAs/71GaAs heterostructure.Upon heating between 800and 1225◦C under an As-rich condition,the Ga diffusion coefﬁcient D was determined by secondary ion mass spectrometry (SIMS)to be D(Ga )=(43±25)cm 2s −1exp[(−4.24±0.06)eV k −1BT −1]over six orders of magnitude in D .In analogy to that experiment,nitrogen self-diffusion was studied by using Ga 14N/Ga 15N/Ga 14N (500/500/500nm)isotopic heterostructures grown on c -plane sapphire by PIMBE [70].Concentration proﬁles of nitrogen isotopes after annealing were measured using SIMS.The activation enthalpy and entropy of nitrogen self-diffusion were obtained by analysing the diffusion length measured for annealing temperatures between 7702660。

退火工艺对高强细晶IF钢的显微组织与性能的影响

第23卷第2期2011年2月钢铁研究学报Jour nal of Ir on and Steel ResearchV ol.23,N o.2February 2011基金项目:国家自然科学基金资助项目(50734002)作者简介:乔立峰(1968 ),男,博士生,教授级高级工程师; E mail:lilyz hm68@; 收稿日期:2009 12 14退火工艺对高强细晶IF 钢的显微组织与性能的影响乔立峰1,2, 刘振宇1, 刘相华1, 王国栋1(1.东北大学轧制及连轧自动化国家重点实验室,辽宁沈阳110004; 2.鞍钢集团鞍凌公司钢轧厂,辽宁凌源122500)摘要:以新型的含铌高强细晶IF 钢为研究对象,在实验室进行了热轧、冷轧以及轧后模拟连续退火试验。

通过微观组织观察可以发现化学成分的改善、轧制及退火工艺的控制不仅可以使这种钢具有细小的晶粒,而且存在大量细小的析出物Nb(C 、N);同时晶界附近析出物非常稀少,称之为P FZ 带(晶界无析出物区),且仅存在于晶界的一侧。

试验结果表明由于铌系析出物非常细小以及晶粒细化作用使试验钢具有较高强度和良好的伸长率;而PF Z 带的存在,这种钢具有较低的屈服强度。

与传统的I F 钢相比,试验钢具有晶粒细小、屈强比低、伸长率良好且塑性应变比r 值较高的特点。

关键词:高强细晶IF 钢;铌碳氮化合物;PF Z 带;力学性能文献标志码:A 文章编号:1001 0963(2011)02 0043 05Effect of Annealing Process on Microstructure and MechanicalProperties of Super Fine Grain IF SteelQIAO Li feng 1,2, LIU Zhen y u 1, LIU Xiang hua 1, WAN G Guo dong1(1.T he St ate Key Labor ator y o f Rolling and A ut omatic,No rtheastern U niver sity,Shenyang 110004,L iaoning ,China;2.A nling Iro n and St eel Co L td Steel M aking and Ro lling P lant of Ansteel G no up,L ing yuan 122500,Liaoning,China)Abstract:T he hot ro lling ,co ld ro lling and simulative continuous annealing ex per iments w ere carr ied out in the la bo rato ry on the base of new t ype SF G H SS (super f ine g rain,high str eng th steel sheet).T he micro structur e ob servation results show t hat t he micr ostr ucture of this new type steel contains not only very fine ferr ite grain but also N b(C,N )pr ecipitates by im pr oving chemical composit ion,contro lled ro lling and co nt rolled annealing.T he P FZ zone is free of precipit ate called P recipitated F ree Z one on t he one side o f the g rain bo undary.T he results sho w that the SFG steel has ver y high tensile str eng th and go od tensile elo ng atio n by fined Nb,T i(C,N)pr ecipitates and ver y fined ferr ite g rain.O n the o ther hand,it also has v ery low yield streng th by the for mation of the PFZ.Con trast to the co nv entional IF steel,t he SFG steel have the character s of super f ine g rain,hig h tensile str eng th,lo w y ield strength/tensile strength r ate,g ood elo ng atio n and high r value.Key words:super f ine hig h st rength IF steel;N b(C,N )pr ecipitate;P FZ;mechanical pr operty汽车减重是降低油耗的主要途径,因而也是减少二氧化碳排放的最有效对策。

用快速光热退火制备多晶硅薄膜的研究

第34卷第2期人　工　晶　体　学　报 Vol.34　No.2 2005年4月 JOURNAL OF SY NTHETI C CRYST ALS Ap ril,2005用快速光热退火制备多晶硅薄膜的研究张宇翔,王海燕,陈永生,杨仕娥,郜小勇,卢景霄,冯团辉,李　瑞,郭　敏(郑州大学物理工程学院,郑州450052)摘要:用等离子体增强型化学气相沉积先得到非晶硅(a2Si:H)薄膜,再用卤钨灯照射的方法对其进行快速光热退火(RPT A),得到了多晶硅薄膜。

然后,进行XRD衍射谱、暗电导率和拉曼光谱等的测量。

结果发现,a2Si:H薄膜在RPT A退火中,退火温度在750℃以上,晶化时间需要2m in,退火温度在650℃以下,晶化时间则需要2.5h;晶化后,晶粒的优先取向是(111)晶向;退火温度850℃时,得到的晶粒最大,暗电导率也最大;退火温度越高,晶化程度越好;退火时间越长,晶粒尺寸越大;光子激励在RPT A退火中起着重要作用。

关键词:多晶硅薄膜;快速光热退火;固相晶化中图分类号:O484 文献标识码:A 文章编号:10002985X(2005)022*******Study on Polycryst a lli n e S ili con F il m s Preparedby Rap i d Photo2ther ma l Annea li n gZHAN G Yu2xiang,WAN G Hai2yan,CHEN Yong2sheng,YAN G Sh i2e,G AO X iao2yong,LU J ing2xiao,FEN G Tuan2hui,L I R ui,G UO M in(College of Physical Science and Technol ogy,Zhengzhou University,Zhengzhou450052,China)(R eceived3O ctober2004)Abstract:An a mor phous silicon fil m was deposited by p las ma enhanced che m ical vapor depositi on(PECVD)and annealed by tungsten hal ogen la mp s heating.The p r ocess can crystallize an a mor phoussilicon fil m in less durati on with l ower ther mal budget than traditi onal furnace annealing.The structure ofthe obtained polycrystalline Si fil m was studied by XRD,Ra man s pectr oscopy and dark conductivity.Theresults show that the annealing te mperature and annealing ti m e have a great effect on the crystallizati on ofa mor phous silicon thin fil m.Key words:polycrystalline silicon fil m;rap id phot o2ther mal annealing;s olid2phase crystallizati on1　引　言多晶硅薄膜是综合了晶体硅材料和非晶硅合金薄膜的优点,在能源科学、信息科学的微电子技术中有着广泛应用的一种新型功能薄膜材料[1]。

Effect of annealing temperature on microstructural evolution and mechanical property of Ti alloy

Short CommunicationEffect of the annealing temperature on the microstructural evolution and mechanical properties of TiZrAlValloyR.Jing a ,⇑,S.X.Liang a ,b ,C.Y.Liu a ,M.Z.Ma a ,R.P.Liu a ,⇑a State Key Laboratory of Metastable Materials Science and Technology,Yanshan University,Qinhuangdao 066004,China bCollege of Equipment Manufacture,Hebei University of Engineering,Handan 056038,Chinaa r t i c l e i n f o Article history:Received 13December 2012Accepted 15June 2013Available online 27June 2013a b s t r a c tThis study aimed to evaluate the effects of the annealing temperature on the structural evolution and mechanical properties of TiZrAlV alloy.The microstructural evolution and mechanical properties of the alloy were investigated by X-ray diffraction,metallographic analysis,tensile testing,and microhardness testing.The results showed that the thickness of the a phase that precipitated from the parent phase was sensitive to the annealing temperature.With increased annealing temperature,the a -phase tended to exhibit equiaxed grains,except for the specimen annealed at 1050°C.The tensile strength of the equi-axed a grains were also demonstrated to have higher tensile strength than those of the lamellar a phase.The optimal mechanical properties of the alloy was obtained after annealing at 850°C,i.e.,r b =1245MPa,r 0.2=1006MPa,and e =16.89%.Ó2013Elsevier Ltd.All rights reserved.1.IntroductionTitanium-based alloys are increasingly being used as structural materials in the aerospace and automotive industries because of their remarkable advantages,such as exceptional strength-to-weight ratio,good hardenability,good elevated temperature performance,excellent fatigue/crack-propagation behavior,and corrosion resistance [1,2].Compared with other conventional stainless steel or structural materials,the mechanical properties of Ti alloys enable their weight to be reduced to about 40%in aerospace and automotive applications [3,4].Currently,the main-stream Ti structural material is the a +b phase Ti–6Al–4V alloy because of its better physical and mechanical properties than com-mercial-purity Ti and other Ti alloys.The a +b phase Ti–6Al–4V alloy is often used in aerospace applications,pressure vessels,blades and discs of aircraft turbines and compressors,surgical implants,etc.[5–8].The mechanical properties of dual-phase Ti alloys are closely related to their microstructure.The metallurgical processes such as thermo-mechanical processing and different heat treatment methods,which bring modiﬁcations in the micro-structure,can strongly inﬂuence their mechanical properties of these alloys [9].The majority of commercially used dual-phase Ti alloys are usually thermo-mechanically processed and subjected to different heat treatments to obtain the ideal microstructure for the desired application.In general,these alloys exist as two typical microstructures,namely,Widmanstätten lath precipitateof the hexagonal close-packed a phase distributed in a matrix of body-centered cubic b phase,and the combination of some equi-axed a -phase grains distributed in a transformed b phase.In general,Ti alloys have low hardness (HV 300–320)and yield strength (880–900MPa)[10].Previous studies [11]have used zir-conium,which has similar chemical properties to Ti,as an alloying element to strengthen Ti–6Al–4V alloy,even though zirconium is considered a neutral element [12,13].The addition of 20%(by mass)Zr to Ti–6Al–4V alloy has been experimentally found to in-crease the alloy strength and microhardness with acceptable elon-gation.In this work,different microstructures of the alloy were obtained by controlling the annealing process.The mechanical properties of the alloy were found to be very sensitive to the annealing temperature.2.Experimental procedureThe alloy used for this study is prepared by electromagnetic induction melting the mixture of sponge Ti (99.7wt%),sponge Zr (Zr +Hf P 99.5wt%),industrially pure Al (99.5wt%)and V (99.9wt%)under an argon atmosphere.Table 1shows the chemical composition of the studied alloy.The alloy was then ﬂipped and re-melted three times to ensure a homogeneous chemical compo-sition.The ingot used in the experiment was homogenized at 1000°C for 12h,followed by cooling to room temperature.Then the ingot underwent multiple breakdowns after being held at 1000°C above the b transus temperature for 90min to completely break the coarse grains.The ingot was held at 900°C for 90min and then subjected to the ﬁnal heat forging in the a +b phase0261-3069/$-see front matter Ó2013Elsevier Ltd.All rights reserved./10.1016/j.matdes.2013.06.039Corresponding authors.Tel.:+863358074723;fax:+863358074545.E-mail addresses:qwe_jr@ (R.Jing),riping@ (R.P.Liu).982R.Jing et al./Materials and Design52(2013)981–986Fig.1.DSC curve of TiZrAlV alloy.region,and the ingot was lathed into bar40mm in diameter.Thesamples(approximately10mmÂ10mmÂ70cm)were cut fromthe bar using wire-electrode cutting and used for subsequentannealing trials.Differential scanning calorimetry(DSC)was used to determinethe phase transition temperatures with a heating rate°C/min which was adopted the standard of ASTM:F2004–05(2010).The nominal a?a+b transus temperature andb?b transus temperature for TiZrAlV alloy are about789and946°C,respectively,as shown in Fig.1.Heat treatment wasperformed in a tubular vacuum furnace under a protective argonpatterns of TiZrAlV alloy:(a)forging,(b)annealing treatment at different temperatures,and(c)detail of33–43°of forging and annealingsigniﬁed that the alloy only formed the solid solution phase and that no other intermetallic compound and/or phase existed (Fig.2a).A comparison of the XRD patterns at different annealing temperatures (Fig.2b)revealed that the phase composition of all annealed alloys consisted of a and b phases.With increased annealing temperature,the b phase (110)reﬂection peak near 38°gradually broadened and the intensity of the (110)diffraction peak increased.However,at 1050°C annealing temperature,the b phase (200)peak disappeared.The XRD patterns also showed that the proportion of a and b phases evidently changed with the chan-ged in annealing temperature.This phenomenon may be caused by the difference of the migration rate of the atom under the high temperature.Generally,with the temperature increasing,the fre-quencies of the atoms migration are also increased gradually.In the insulation process,the moving distancesof Al atom (which is a -stabilized element)and V atom (which is b -stabilized element)were different,and in the subsequent cooling process,the b phase transformed into the a phase which caused the Al atom enriched in the b phase lattice and changed the b phase lattice parameters.Therefore,it may make the intensity of the b phase (110)reﬂection peak increase and the (002)reﬂection peak decrease when the annealing temperature was heated to 1050°C.Furthermore,the annealing holding time was shorter (30min),in this process,the a phase transformed into the b phase may be incomplete at an-nealed treatment at 1000°C,while annealed temperature was in-creased to 1050°C,the a phase may be completely transformed into the b phase,therefore,in the specimen annealed at 1000°C the initial a phase was also existed,but the specimen which was annealed at 1050°C did not exist the initial a phase.This may re-sult that the differences of a phase between 1000°C and 1050°C is obviously.Fig.3shows the microstructure of the annealing temperatures.The specimens ited Widmanstätten morphology (Fig.3a),i.e.,chaotic arrangement of slender a lath and b annealing temperature to 1000°C,the b peared and the a lath gradually (Fig.3b–e).In this process,the alloy axed trend with increased annealing have caused the increased equiaxed a phase the annealing process.First,the lamellar a ‘‘interleaved,’’which restricts the other a longitudinal direction.Consequently,a only along the transverse direction,which promotes the thickening of the a lamellar.Second,the new a phase that precipitates from the parent phase grows along a speciﬁc habit plane and has a cer-tain orientation relationship with the primary a phase.Thus,the new precipitated a phase growing along the longitudinal direction is hindered such that the equiaxed degree is increased.Obasi [14]also indicated that the phase transformation in Ti alloys during heating (a ?b )and cooling (b ?a )is governed by the so-called Burgers orientation relationship {0002}a ||{110}b and h 11À20i a ||h 111i b with 6possible b -orientations during the a ?b phase transformation and 12possible a orientations that can transform from a single parent b grain during b ?a phase transformation.However,when the annealing temperature reached 1050°C,the alloy revealed the typical basketweave mor-phology (Fig.3f),i.e.,a crisscross slender a lath.b grain boundaries and some parallel lamellar the grain boundaries (the Widmanstätten microstructure)observed in this process.In most diffusion phase and precipitation processes,the nucleations of the heterogeneously occurs at some preferential nucleation the matrix such as the grain boundary,dislocation,phase [15].When the annealing temperature (e.g.,Optical microstructure of TiZrAlV alloy under different annealing temperature:(a)800°C,(b)850°C,(c)900°C,(d)950°C,(e)1000°C,and Fig.4.True stress–strain curve of the studied alloy under different conditions.the thickness of the a lath became limited.Therefore,the thicknessof the new precipitated a phase after annealing at1050°C was smaller than that after annealing at1000°C.The mechanical properties of the alloy were evaluated through uniaxial tensile tests.Fig.4and Table2show the true stress–strain curves and mechanical properties of the specimens at different annealing temperatures.The mechanical properties of the ZrTiAlV alloy evidently depended on the annealing temperature and micro-structure.When the annealing temperatures were between800 and1000°C,the yield strength r0.2and ultimate strength r b de-creased from1009and1290MPa to978and1181MPa,respec-tively.The elongation only slightly changed after annealing at of the residual b phase during the annealing process as well as the thickness of the a lath.The main factors inﬂuencing the mechanical properties of an-nealed samples in which only the a and b phases exist are the phase content,size,and morphology of the a phase[16–18].Because of the limited number of independent slips modes,the hcp structure of Ti exhibits a vary strong grain-boundary,or Hall–Petch strength-ening at room temperature.The thickness of the a grain boundary directly inﬂuences the strength mismatch between the a+b matrix and the grain boundary[19].Consider the case of b processed microstructures.Some of the microstructural features involved with progressively increasing length scales are width of the a-laths, the colony size,and the b grain size(feature sizes may range from sub-micron to millimeters).Depending on the thermo-mechanical treatment the alloy is subjected to,such as cooling rates from +b dual phase region or above the b-transus,these features can vary signiﬁcantly.Quantifying them over the diverse range length scales involved becomes rather important.Thus,to investi-gate the effect of the annealing temperature on the microstructure and mechanical properties,the specimens prepared at different annealing temperatures were subjected to SEM analysis,as shown Fig.6.The measured thicknesses of the a lath from the SEM images are shown in Fig.7a.With increased annealing temperature Fig.5.Microhardness of annealed specimens under different conditions.SEM images of TiZrAlV alloy under different annealing temperature:(a)800°C,(b)850°C,(c)900°C,(d)950°C,(e)1000°C,and(f)from 800°C to 1000°C,the thickness of the a lath in the annealed samples increased from 1.07l m to 4.22l m.When the annealing temperature reached 1050°C,the thickness was reduced to 1.12l m.According to the Hall–Petch equation,(i.e.,r =r 0+kd À1/2,where d is the thickness of the a lath),the strength of an annealed alloy is related to the a lath thickness,as shown in Fig.7b.On one hand,the change of the a lath thickness moved the distance of dis-location to the phase boundary,which resulted in increased num-ber of dislocations piling up such that the stress concentration was more severe.On the other hand,reducing the a lath thickness increased the density of the phase boundary in the same cross-sectional area.Consequently,the movement of the dislocation obstacle increased.Thus,based on the OM images,SEM images,and true stress–strain curve,the slender a lath obtained at 800°C and 1050°C increased the strength of the specimens and made dis-location movement difﬁcultly.Moreover,with increased equiaxed a -phase degree,the strength of the annealed specimens gradually decreased.This result implied that the strength of the processed alloy lamellar a phase microstructure was higher than that of processed equiaxed a -phase microstructure.The magnitude of the titanium alloy tensile elongation is con-nected with the non-uniform degree of the tensile micro deforma-tion zone,as well as the length and the spacing of slip bands.With the spacing of slip bands decreasing,the plastic deformability in-crease before the material fracture [20].Compared with the lamel-lae microstructure,the slip bands spacing of the duplex microstructure is smaller,thus this microstructure possess a high-er ability of deformation.When sample was annealed at 850°C,the feature of microstructure presented the duplex microstructure (Fig.3b),therefore,the elongation reached the largest value in this experiment i.e.16.89%.4.ConclusionThe phase transition,microstructure evolution,and their effects on the mechanical properties of TiZrAlV alloy were investigated.The conclusions were as follows:(1)TiZrAlV alloy exhibited an a +b phase after high-tempera-ture annealing.The intensity of the b (110)diffraction peak increased with increased annealing temperature.However,the intensity of the b (200)diffraction peak gradually decreased with increased annealing temperature.When the temperature reached 1050°C,the b (200)diffraction peak completely disappeared.(2)The thickness of the a phase was sensitive to the annealingtemperature.With increased annealing temperature,the a phase tended to exhibit equiaxed grains,except for speci-mens annealed at 1050°C.After annealing at 1000°C,the maximum thickness was 4.22l m.(3)The mechanical properties of the annealed specimens weresensitive to the morphology of the precipitated a phase and yo the annealing temperature.The optimal mechanical properties of the alloy were obtained after annealing at 850°C,i.e.,r b =1245MPa,r 0.2=1006MPa,and e =16.89%.AcknowledgmentsThis work was supported by the SKPBRC (Grant No.2010CB731600),NSFC (Grant No.51121061/51171160/51171163).References[1]Eylon D,Vassel A,Combres Y,Boyer RR,Bania PJ,Schutz RW.Issues in thedevelopment of beta titanium alloys.JOM 1994;46:14–5.[2]Ivasishin OM,Markovsky PE,Matviychuk YuV,Semiatin SL.Precipitation andrecrystallization behavior of beta titanium alloys during continuous heat treatment.Metall Mater Trans A 2003;34(1):147–58.[3]Lütjering G,Williams JC.Titanium,Springer,B.Derby,Ed.,Springer-Verlag,Berlin,Heidelberg,Germany,2003,p.27–260.[4]Schauerte O.Titanium in automotive production.Adv Eng Mater2003;6:411–8.[5]Okazaki Y,Rao S,Ito Y,Tateishi T.Corrosion resistance,mechanical properties,corrosion fatigue strength and cytocompatibility of new Ti alloys without Al and V.Biomaterials 1998;19:1197.[6]Schutz RW,Watkins HB.Recent developments in titanium alloy application inthe energy industry.Mater Sci Eng A 1998;243:305–15.[7]Gorynin IV.Modeling of the motion of particles in a rotary crusher.Mater SciEng A 1999;263:112.[8]Cheng WW,Chern Lin JH,Ju CP.Bismuth effect on castability and mechanicalproperties of Ti–6Al–4V alloy cast in copper mold.Mater Lett 2003;57(16–17):2591–6.[9]Sieniawski J,Filip R,Ziaja W.The effect of 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nucleation on defectsin diffusional phase transformation and precipitation.Mater Sci Eng A 2001;312:145–54.[16]Tiley J,Searles T,Lee E,Kar S,Banerjee R,Russ JC,et al.Quantiﬁcation ofmicrostructural features in a/b titanium alloys.Mater Sci Eng A 2004;372:191–8.[17]Kong FT,Chen Y,Yang F.Effect of heat treatment on microstructures andtensile properties of as-forged Ti–45Al–5Nb–0.3Y alloy.Intermetallics 2011;19(2):212–6.[18]Rack HJ,Qazi JI.Titanium alloys for biomedical applications.Mater Sci Eng C2006;26(8):1269–77.[19]Lütjering G,Albrecht J.Inﬂuence of cooling rate and b grain size on the tensileproperties of(a+b)Ti alloys.In:Proceedings of the8th world titanium conference;1995.[20]Terlinde G,Luetjering G.Inﬂuence of grain size and age hardening ondislocation pile-ups and tensile fracture for a Ti–Al alloy.Metall Trans 1982;13(7):1283–92.986R.Jing et al./Materials and Design52(2013)981–986。

大角度晶界的英语

大角度晶界的英语Abstract:Macroscopic grain boundaries (MGBs) are a critical feature in polycrystalline materials, significantly influencing mechanical properties, thermal conductivity, and electrical conductivity. This paper delves into the characteristics, formation mechanisms, and the impact of MGBs on the performance of materials, with a focus on their rolein various applications.1. IntroductionGrain boundaries are interfaces between two crystalline grains in a polycrystalline solid. When the misorientation between grains is significant, these boundaries can be considered macroscopic, exhibiting distinct properties that differ from those of the bulk material. The study of MGBs is essential for understanding material behavior and optimizing performance in engineering applications.2. Characteristics of Macroscopic Grain BoundariesMacroscopic grain boundaries are characterized by their misorientation angles, which are typically greater than 15 degrees. They can be classified into several types based on the crystallographic relationship between the grains they separate:- Twin boundaries: Where the misorientation is a mirror reflection across the boundary plane.- Coincidence site lattice (CSL) boundaries: Where a high density of lattice points from both grains coincide at the boundary.- General boundaries: With no specific crystallographic relationship, these are the most common type.3. Formation MechanismsMGBs can form during various material processing techniques:- Recrystallization: After severe deformation, grains can grow, leading to the formation of MGBs.- Grain growth: During annealing, larger grains can consume smaller ones, resulting in increased misorientations at the boundaries.- Phase transformations: Changes in crystal structure during phase transitions can create MGBs.4. Impact on Material PropertiesThe presence of MGBs has a profound effect on the properties of polycrystalline materials:- Strength: MGBs can impede dislocation motion, increasing the material's strength.- Ductility: They can act as sites for crack initiation, affecting ductility.- Conductivity: MGBs can scatter electrons and phonons, reducing thermal and electrical conductivity.5. ApplicationsUnderstanding MGBs is crucial for optimizing materials in various applications:- Metalworking: Controlling grain size and boundary characteristics can enhance the mechanical properties of metals.- Electronics: In semiconductor devices, MGBs can influence carrier mobility and device performance.- Ceramics: MGBs in ceramics can affect fracture toughness and thermal shock resistance.6. Experimental Techniques for Studying MGBsSeveral experimental methods are used to study MGBs:- Scanning electron microscopy (SEM): Can reveal the morphology of MGBs.- Transmission electron microscopy (TEM): Provides detailed information on the atomic structure of boundaries.- Electron backscatter diffraction (EBSD): Allows for the determination of grain orientation and the identification of MGBs.7. Computational ModelingComputational techniques, such as molecular dynamics and phase-field modeling, are used to simulate MGB formation and behavior:- Molecular Dynamics (MD): Offers insights into atomic-scale processes at MGBs.- Phase-Field Modeling: Can predict the evolution of grain structures and boundary characteristics during processing.8. ConclusionMacroscopic grain boundaries play a critical role in determining the properties of polycrystalline materials. Understanding their characteristics, formation, and impact is essential for the development of advanced materials with tailored properties for specific applications. Future research should focus on developing new techniques to control MGBs and on multiscale modeling to predict their effects on material behavior.References1. Hull, D., & Bacon, D. J. (2011). Introduction to Dislocations (5th ed.). Butterworth-Heinemann.2. Humphreys, F. J., & Hatherly, M. (2004). Recrystallization and Related Annealing Phenomena (2nd ed.). Elsevier.3. Gottstein, G. (2002). Physical Foundations ofMaterials Science. Springer.4. Randle, V. (2017). Grain Boundary CharacterDistribution and its Applications. CRC Press.This document provides a comprehensive overview of macroscopic grain boundaries, discussing theircharacteristics, formation, and impact on material properties, as well as the experimental and computational techniques used to study them. It concludes with the significance of MGBs in material applications and the importance of future researchin this area.。

黄永刚单晶塑性有限元umat子程序

SUBROUTINE UMAT(stress,statev,ddsdde,sse,spd,scd,1 rpl, ddsddt, drplde, drpldt,2 stran,dstran,time,dtime,temp,dtemp,predef,dpred,cmname,3 ndi,nshr,ntens,nstatv,props,nprops,coords,drot,pnewdt,4 celent,dfgrd0,dfgrd1,noel,npt,layer,kspt,kstep,kinc)c WRITE (6,*) 'c NOTE: MODIFICATIONS TO *UMAT FOR ABAQUS VERSION 5.3 (14 APR '94) cc (1) The list of variables above defining the *UMAT subroutine,c and the first (standard) block of variables dimensioned below,c have variable names added compared to earlier ABAQUS versions.cc (2) The statement: include 'aba_param.inc' must be added as below.cc (3) As of version 5.3, ABAQUS files use double precision only.c The file aba_param.inc has a line "implicit real*8" and, sincec it is included in the main subroutine, it will define the variablesc there as double precision. But other subroutines still need thec definition "implicit real*8" since there may be variables that arec not passed to them through the list or common block.cc (4) This is current as of version 5.6 of ABAQUS.cc (5) Note added by J. W. Kysar (4 November 1997). This UMAT has beenc modified to keep track of the cumulative shear strain in eachc individual slip system. This information is needed to correct anc error in the implementation of the Bassani and Wu hardening law.c Any line of code which has been added or modified is precededc immediately by a line beginning CFIXA and succeeded by a linec beginning CFIXB. Any comment line added or modified will beginc with CFIX.cc The hardening law by Bassani and Wu was implemented incorrectly.c This law is a function of both hyperbolic secant squared and hyperbolicc tangent. However, the arguments of sech and tanh are related to the *total*c slip on individual slip systems. Formerly, the UMAT implemented thisc hardening law by using the *current* slip on each slip system. Thereinc lay the problem. The UMAT did not restrict the current slip to be ac positive value. So when a slip with a negative sign was encountered, thec term containing tanh led to a negative hardening rate (since tanh is anc odd function).c The UMA T has been fixed by adding state variables to keep track of thec *total* slip on each slip system by integrating up the absolute valuec of slip rates for each individual slip system. These "solution dependentc variables" are available for postprocessing. The only required changec in the input file is that the DEPV AR command must be changed.cC----- Use single precision on Cray byC (1) deleting the statement "IMPLICIT*8 (A-H,O-Z)";C (2) changing "REAL*8 FUNCTION" to "FUNCTION";C (3) changing double precision functions DSIGN to SIGN.CC----- Subroutines:CC ROTATION -- forming rotation matrix, i.e. the directionC cosines of cubic crystal [100], [010] and [001]C directions in global system at the initialC stateCC SLIPSYS -- calculating number of slip systems, unitC vectors in slip directions and unit normals toC slip planes in a cubic crystal at the initialC stateCC GSLPINIT -- calculating initial value of current strengthsC at initial stateCC STRAINRA TE -- based on current values of resolved shearC stresses and current strength, calculatingC shear strain-rates in slip systemsCC LATENTHARDEN -- forming self- and latent-hardening matrixCC ITERATION -- generating arrays for the Newton-RhapsonC iterationCC LUDCMP -- LU decompositionCC LUBKSB -- linear equation solver based on LUC decomposition method (must call LUDCMP first) C----- Function subprogram:C F -- shear strain-rates in slip systemsC----- Variables:CC STRESS -- stresses (INPUT & OUTPUT)C Cauchy stresses for finite deformationC STATEV -- solution dependent state variables (INPUT & OUTPUT) C DDSDDE -- Jacobian matrix (OUTPUT)C----- Variables passed in for information:CC STRAN -- strainsC logarithmic strain for finite deformationC (actually, integral of the symmetric part of velocityC gradient with respect to time)C DSTRAN -- increments of strainsC CMNAME -- name given in the *MATERIAL optionC NDI -- number of direct stress componentsC NSHR -- number of engineering shear stress componentsC NTENS -- NDI+NSHRC NSTATV -- number of solution dependent state variables (asC defined in the *DEPVAR option)C PROPS -- material constants entered in the *USER MA TERIAL C optionC NPROPS -- number of material constantsCC----- This subroutine provides the plastic constitutive relation ofC single crystals for finite element code ABAQUS. The plastic slipC of single crystal obeys the Schmid law. The program gives theC choice of small deformation theory and theory of finite rotationC and finite strain.C The strain increment is composed of elastic part and plasticC part. The elastic strain increment corresponds to latticeC stretching, the plastic part is the sum over all slip systems ofC plastic slip. The shear strain increment for each slip system isC assumed a function of the ratio of corresponding resolved shearC stress over current strength, and of the time step. The resolvedC shear stress is the double product of stress tensor with the slipC deformation tensor (Schmid factor), and the increment of currentC strength is related to shear strain increments over all slipC systems through self- and latent-hardening functions.C----- The implicit integration method proposed by Peirce, Shih andC Needleman (1984) is used here. The subroutine provides an option C of iteration to solve stresses and solution dependent stateC variables within each increment.C----- The present program is for a single CUBIC crystal. However,C this code can be generalized for other crystals (e.g. HCP,C Tetragonal, Orthotropic, etc.). Only subroutines ROTATION andC SLIPSYS need to be modified to include the effect of crystalC aspect ratio.CC----- Important notice:CC (1) The number of state variables NSTATV must be larger than (or CFIX equal to) TEN (10) times the total number of slip systems inC all sets, NSLPTL, plus FIVE (5)CFIX NSTATV >= 10 * NSLPTL + 5C Denote s as a slip direction and m as normal to a slip plane.C Here (s,-m), (-s,m) and (-s,-m) are NOT consideredC independent of (s,m). The number of slip systems in each setC could be either 6, 12, 24 or 48 for a cubic crystal, e.g. 12C for {110}<111>.CC Users who need more parameters to characterize theC constitutive law of single crystal, e.g. the frameworkC proposed by Zarka, should make NSTATV larger than (or equal C to) the number of those parameters NPARMT plus nine timesC the total number of slip systems, NSLPTL, plus fiveCFIX NSTATV >= NPARMT + 10 * NSLPTL + 5CC (2) The tangent stiffness matrix in general is not symmetric ifC latent hardening is considered. Users must declare "UNSYMM"C in the input file, at the *USER MATERIAL card.CPARAMETER (ND=150)C----- The parameter ND determines the dimensions of the arrays inC this subroutine. The current choice 150 is a upper bound for aC cubic crystal with up to three sets of slip systems activated.C Users may reduce the parameter ND to any number as long as larger C than or equal to the total number of slip systems in all sets.C For example, if {110}<111> is the only set of slip systemC potentially activated, ND could be taken as twelve (12).cinclude 'aba_param.inc'cCHARACTER*8 CMNAMEEXTERNAL Fdimension stress(ntens),statev(nstatv),1 ddsdde(ntens,ntens),ddsddt(ntens),drplde(ntens),2 stran(ntens),dstran(ntens),time(2),predef(1),dpred(1),3 props(nprops),coords(3),drot(3,3),dfgrd0(3,3),dfgrd1(3,3)DIMENSION ISPDIR(3), ISPNOR(3), NSLIP(3),2 SLPDIR(3,ND), SLPNOR(3,ND), SLPDEF(6,ND),3 SLPSPN(3,ND), DSPDIR(3,ND), DSPNOR(3,ND),4 DLOCAL(6,6), D(6,6), ROTD(6,6), ROTATE(3,3),5 FSLIP(ND), DFDXSP(ND), DDEMSD(6,ND),6 H(ND,ND), DDGDDE(ND,6),7 DSTRES(6), DELATS(6), DSPIN(3), DVGRAD(3,3),8 DGAMMA(ND), DTAUSP(ND), DGSLIP(ND),9 WORKST(ND,ND), INDX(ND), TERM(3,3), TRM0(3,3), ITRM(3)DIMENSION FSLIP1(ND), STRES1(6), GAMMA1(ND), TAUSP1(ND),2 GSLP1(ND), SPNOR1(3,ND), SPDIR1(3,ND), DDSDE1(6,6),3 DSOLD(6), DGAMOD(ND), DTAUOD(ND), DGSPOD(ND),4 DSPNRO(3,ND), DSPDRO(3,ND),5 DHDGDG(ND,ND)C----- NSLIP -- number of slip systems in each setC----- SLPDIR -- slip directions (unit vectors in the initial state)C----- SLPNOR -- normals to slip planes (unit normals in the initialC state)C----- SLPDEF -- slip deformation tensors (Schmid factors)C SLPDEF(1,i) -- SLPDIR(1,i)*SLPNOR(1,i)C SLPDEF(2,i) -- SLPDIR(2,i)*SLPNOR(2,i)C SLPDEF(3,i) -- SLPDIR(3,i)*SLPNOR(3,i)C SLPDEF(4,i) -- SLPDIR(1,i)*SLPNOR(2,i)+C SLPDIR(2,i)*SLPNOR(1,i)C SLPDEF(5,i) -- SLPDIR(1,i)*SLPNOR(3,i)+C SLPDIR(3,i)*SLPNOR(1,i)C SLPDEF(6,i) -- SLPDIR(2,i)*SLPNOR(3,i)+C SLPDIR(3,i)*SLPNOR(2,i)C where index i corresponds to the ith slip systemC----- SLPSPN -- slip spin tensors (only needed for finite rotation)C SLPSPN(1,i) -- [SLPDIR(1,i)*SLPNOR(2,i)-C SLPDIR(2,i)*SLPNOR(1,i)]/2C SLPSPN(2,i) -- [SLPDIR(3,i)*SLPNOR(1,i)-C SLPDIR(1,i)*SLPNOR(3,i)]/2C SLPSPN(3,i) -- [SLPDIR(2,i)*SLPNOR(3,i)-C SLPDIR(3,i)*SLPNOR(2,i)]/2C where index i corresponds to the ith slip systemC----- DSPDIR -- increments of slip directionsC----- DSPNOR -- increments of normals to slip planesCC----- DLOCAL -- elastic matrix in local cubic crystal systemC----- D -- elastic matrix in global systemC----- ROTD -- rotation matrix transforming DLOCAL to DCC----- ROTATE -- rotation matrix, direction cosines of [100], [010]C and [001] of cubic crystal in global systemCC----- FSLIP -- shear strain-rates in slip systemsC----- DFDXSP -- derivatives of FSLIP w.r.t x=TAUSLP/GSLIP, whereC TAUSLP is the resolved shear stress and GSLIP is the C current strengthCC----- DDEMSD -- double dot product of the elastic moduli tensor withC the slip deformation tensor plus, only for finiteC rotation, the dot product of slip spin tensor withC the stressCC----- H -- self- and latent-hardening matrixC H(i,i) -- self hardening modulus of the ith slipC system (no sum over i)C H(i,j) -- latent hardening molulus of the ith slipC system due to a slip in the jth slip system C (i not equal j)CC----- DDGDDE -- derivatice of the shear strain increments in slipC systems w.r.t. the increment of strainsCC----- DSTRES -- Jaumann increments of stresses, i.e. corotationalC stress-increments formed on axes spinning with theC materialC----- DELATS -- strain-increments associated with lattice stretchingC DELATS(1) - DELATS(3) -- normal strain increments C DELATS(4) - DELATS(6) -- engineering shear strain C incrementsC----- DSPIN -- spin-increments associated with the material elementC DSPIN(1) -- component 12 of the spin tensorC DSPIN(2) -- component 31 of the spin tensorC DSPIN(3) -- component 23 of the spin tensorCC----- DVGRAD -- increments of deformation gradient in the currentC state, i.e. velocity gradient times the increment ofC timeCC----- DGAMMA -- increment of shear strains in slip systemsC----- DTAUSP -- increment of resolved shear stresses in slip systemsC----- DGSLIP -- increment of current strengths in slip systemsCCC----- Arrays for iteration:CC FSLIP1, STRES1, GAMMA1, TAUSP1, GSLP1 , SPNOR1, SPDIR1,C DDSDE1, DSOLD , DGAMOD, DTAUOD, DGSPOD, DSPNRO, DSPDRO, C DHDGDGCCC----- Solution dependent state variable STATEV:C Denote the number of total slip systems by NSLPTL, whichC will be calculated in this code.CC Array STATEV:C 1 - NSLPTL : current strength in slip systemsC NSLPTL+1 - 2*NSLPTL : shear strain in slip systemsC 2*NSLPTL+1 - 3*NSLPTL : resolved shear stress in slip systemsCC 3*NSLPTL+1 - 6*NSLPTL : current components of normals to slipC slip planesC 6*NSLPTL+1 - 9*NSLPTL : current components of slip directionsCCFIX 9*NSLPTL+1 - 10*NSLPTL : total cumulative shear strain on eachCFIX slip system (sum of the absoluteCFIX values of shear strains in each slipCFIX system individually)CFIXCFIX 10*NSLPTL+1 : total cumulative shear strain on allC slip systems (sum of the absoluteC values of shear strains in all slipC systems)CCFIX 10*NSLPTL+2 - NSTA TV-4 : additional parameters users may needC to characterize the constitutive lawC of a single crystal (if there areC any).CC NSTATV-3 : number of slip systems in the 1st setC NSTATV-2 : number of slip systems in the 2nd setC NSTATV-1 : number of slip systems in the 3rd setC NSTATV : total number of slip systems in allC setsCCC----- Material constants PROPS:CC PROPS(1) - PROPS(21) -- elastic constants for a general elasticC anisotropic materialCC isotropic : PROPS(i)=0 for i>2C PROPS(1) -- Young's modulusC PROPS(2) -- Poisson's ratioCC cubic : PROPS(i)=0 for i>3C PROPS(1) -- c11C PROPS(2) -- c12C PROPS(3) -- c44CC orthotropic : PORPS(i)=0 for i>9C PROPS(1) - PROPS(9) are D1111, D1122, D2222, C D1133, D2233, D3333, D1212, D1313, D2323,C respectively, which has the same definitionC as ABAQUS for orthotropic materialsC (see *ELASTIC card)CC anisotropic : PROPS(1) - PROPS(21) are D1111, D1122,C D2222, D1133, D2233, D3333, D1112, D2212,C D3312, D1212, D1113, D2213, D3313, D1213,C D1313, D1123, D2223, D3323, D1223, D1323,C D2323, respectively, which has the sameC definition as ABAQUS for anisotropicC materials (see *ELASTIC card)CCC PROPS(25) - PROPS(56) -- parameters characterizing all slipC systems to be activated in a cubicC crystalCC PROPS(25) -- number of sets of slip systems (maximum 3),C e.g. (110)[1-11] and (101)[11-1] are in theC same set of slip systems, (110)[1-11] andC (121)[1-11] belong to different sets of slipC systemsC (It must be a real number, e.g. 3., not 3 !)CC PROPS(33) - PROPS(35) -- normal to a typical slip plane inC the first set of slip systems,C e.g. (1 1 0)C (They must be real numbers, e.g.C 1. 1. 0., not 1 1 0 !)C PROPS(36) - PROPS(38) -- a typical slip direction in theC first set of slip systems, e.g.C [1 1 1]C (They must be real numbers, e.g.C 1. 1. 1., not 1 1 1 !)CC PROPS(41) - PROPS(43) -- normal to a typical slip plane inC the second set of slip systemsC (real numbers)C PROPS(44) - PROPS(46) -- a typical slip direction in theC second set of slip systemsC (real numbers)CC PROPS(49) - PROPS(51) -- normal to a typical slip plane inC the third set of slip systemsC (real numbers)C PROPS(52) - PROPS(54) -- a typical slip direction in theC third set of slip systemsC (real numbers)CCC PROPS(57) - PROPS(72) -- parameters characterizing the initialC orientation of a single crystal inC global systemC The directions in global system and directions in localC cubic crystal system of two nonparallel vectors are neededC to determine the crystal orientation.CC PROPS(57) - PROPS(59) -- [p1 p2 p3], direction of firstC vector in local cubic crystalC system, e.g. [1 1 0]C (They must be real numbers, e.g.C 1. 1. 0., not 1 1 0 !)C PROPS(60) - PROPS(62) -- [P1 P2 P3], direction of firstC vector in global system, e.g.C [2. 1. 0.]C (It does not have to be a unitC vector)CC PROPS(65) - PROPS(67) -- direction of second vector inC local cubic crystal system (real C numbers)C PROPS(68) - PROPS(70) -- direction of second vector inC global systemCCC PROPS(73) - PROPS(96) -- parameters characterizing the visco-C plastic constitutive law (shearC strain-rate vs. resolved shearC stress), e.g. a power-law relationCC PROPS(73) - PROPS(80) -- parameters for the first set ofC slip systemsC PROPS(81) - PROPS(88) -- parameters for the second set of C slip systemsC PROPS(89) - PROPS(96) -- parameters for the third set ofC slip systemsCCC PROPS(97) - PROPS(144)-- parameters characterizing the self-C and latent-hardening laws of slipC systemsCC PROPS(97) - PROPS(104)-- self-hardening parameters for the C first set of slip systemsC PROPS(105)- PROPS(112)-- latent-hardening parameters for C the first set of slip systems and C interaction with other sets ofC slip systemsCC PROPS(113)- PROPS(120)-- self-hardening parameters for the C second set of slip systemsC PROPS(121)- PROPS(128)-- latent-hardening parameters for C the second set of slip systems C and interaction with other sets C of slip systemsCC PROPS(129)- PROPS(136)-- self-hardening parameters for the C third set of slip systemsC PROPS(137)- PROPS(144)-- latent-hardening parameters for C the third set of slip systems and C interaction with other sets ofC slip systemsCCC PROPS(145)- PROPS(152)-- parameters characterizing forward time C integration scheme and finiteC deformationCC PROPS(145) -- parameter theta controlling the implicitC integration, which is between 0 and 1C 0. : explicit integrationC 0.5 : recommended valueC 1. : fully implicit integrationCC PROPS(146) -- parameter NLGEOM controlling whether the C effect of finite rotation and finite strainC of crystal is considered,C 0. : small deformation theoryC otherwise : theory of finite rotation andC finite strainCCC PROPS(153)- PROPS(160)-- parameters characterizing iterationC methodCC PROPS(153) -- parameter ITRATN controlling whether theC iteration method is used,C 0. : no iterationC otherwise : iterationCC PROPS(154) -- maximum number of iteration ITRMAXCC PROPS(155) -- absolute error of shear strains in slipC systems GAMERRCC----- Elastic matrix in local cubic crystal system: DLOCALDO J=1,6DO I=1,6DLOCAL(I,J)=0.END DOEND DOCHECK=0.DO J=10,21CHECK=CHECK+ABS(PROPS(J))END DOIF (CHECK.EQ.0.) THENDO J=4,9CHECK=CHECK+ABS(PROPS(J))END DOIF (CHECK.EQ.0.) THENIF (PROPS(3).EQ.0.) THENC----- Isotropic materialGSHEAR=PROPS(1)/2./(1.+PROPS(2))E11=2.*GSHEAR*(1.-PROPS(2))/(1.-2.*PROPS(2))E12=2.*GSHEAR*PROPS(2)/(1.-2.*PROPS(2))DO J=1,3DLOCAL(J,J)=E11DO I=1,3IF (I.NE.J) DLOCAL(I,J)=E12END DODLOCAL(J+3,J+3)=GSHEAREND DOELSEC----- Cubic materialDO J=1,3DLOCAL(J,J)=PROPS(1)DO I=1,3IF (I.NE.J) DLOCAL(I,J)=PROPS(2)END DODLOCAL(J+3,J+3)=PROPS(3)END DOEND IFELSEC----- Orthotropic metarialDLOCAL(1,1)=PROPS(1)DLOCAL(1,2)=PROPS(2)DLOCAL(2,1)=PROPS(2)DLOCAL(2,2)=PROPS(3)DLOCAL(1,3)=PROPS(4)DLOCAL(3,1)=PROPS(4)DLOCAL(2,3)=PROPS(5)DLOCAL(3,2)=PROPS(5)DLOCAL(3,3)=PROPS(6)DLOCAL(4,4)=PROPS(7)DLOCAL(5,5)=PROPS(8)DLOCAL(6,6)=PROPS(9)END IFELSEC----- General anisotropic materialID=0DO J=1,6DO I=1,JID=ID+1DLOCAL(I,J)=PROPS(ID)DLOCAL(J,I)=DLOCAL(I,J)END DOEND DOEND IFC----- Rotation matrix: ROTA TE, i.e. direction cosines of [100], [010]C and [001] of a cubic crystal in global systemCCALL ROTATION (PROPS(57), ROTA TE)C----- Rotation matrix: ROTD to transform local elastic matrix DLOCAL C to global elastic matrix DCDO J=1,3J1=1+J/3J2=2+J/2DO I=1,3I1=1+I/3I2=2+I/2ROTD(I,J)=ROTATE(I,J)**2ROTD(I,J+3)=2.*ROTATE(I,J1)*ROTA TE(I,J2)ROTD(I+3,J)=ROTA TE(I1,J)*ROTATE(I2,J)ROTD(I+3,J+3)=ROTA TE(I1,J1)*ROTA TE(I2,J2)+2 ROTA TE(I1,J2)*ROTA TE(I2,J1)END DOEND DOC----- Elastic matrix in global system: DC {D} = {ROTD} * {DLOCAL} * {ROTD}transposeCDO J=1,6DO I=1,6D(I,J)=0.END DOEND DODO J=1,6DO I=1,JDO K=1,6DO L=1,6D(I,J)=D(I,J)+DLOCAL(K,L)*ROTD(I,K)*ROTD(J,L) END DOEND DOD(J,I)=D(I,J)END DOEND DOC----- Total number of sets of slip systems: NSETNSET=NINT(PROPS(25))IF (NSET.LT.1) THENWRITE (6,*) '***ERROR - zero sets of slip systems'STOPELSE IF (NSET.GT.3) THENWRITE (6,*)2 '***ERROR - more than three sets of slip systems'STOPEND IFC----- Implicit integration parameter: THETATHETA=PROPS(145)C----- Finite deformation ?C----- NLGEOM = 0, small deformation theoryC otherwise, theory of finite rotation and finite strain, Users C must declare "NLGEOM" in the input file, at the *STEP card CIF (PROPS(146).EQ.0.) THENNLGEOM=0ELSENLGEOM=1END IFC----- Iteration?C----- ITRATN = 0, no iterationC otherwise, iteration (solving increments of stresses andC solution dependent state variables)CIF (PROPS(153).EQ.0.) THENITRATN=0ELSEITRATN=1END IFITRMAX=NINT(PROPS(154))GAMERR=PROPS(155)NITRTN=-1DO I=1,NTENSDSOLD(I)=0.END DODO J=1,NDDGAMOD(J)=0.DTAUOD(J)=0.DGSPOD(J)=0.DO I=1,3DSPNRO(I,J)=0.DSPDRO(I,J)=0.END DOEND DOC----- Increment of spin associated with the material element: DSPIN C (only needed for finite rotation)CIF (NLGEOM.NE.0) THENDO J=1,3DO I=1,3TERM(I,J)=DROT(J,I)TRM0(I,J)=DROT(J,I)END DOTERM(J,J)=TERM(J,J)+1.D0TRM0(J,J)=TRM0(J,J)-1.D0END DOCALL LUDCMP (TERM, 3, 3, ITRM, DDCMP)DO J=1,3CALL LUBKSB (TERM, 3, 3, ITRM, TRM0(1,J)) END DODSPIN(1)=TRM0(2,1)-TRM0(1,2)DSPIN(2)=TRM0(1,3)-TRM0(3,1)DSPIN(3)=TRM0(3,2)-TRM0(2,3)END IFC----- Increment of dilatational strain: DEVDEV=0.D0DO I=1,NDIDEV=DEV+DSTRAN(I)END DOC----- Iteration starts (only when iteration method is used)1000 CONTINUEC----- Parameter NITRTN: number of iterationsC NITRTN = 0 --- no-iteration solutionCNITRTN=NITRTN+1C----- Check whether the current stress state is the initial stateIF (STATEV(1).EQ.0.) THENC----- Initial stateCC----- Generating the following parameters and variables at initialC state:C Total number of slip systems in all the sets NSLPTLC Number of slip systems in each set NSLIPC Unit vectors in initial slip directions SLPDIRC Unit normals to initial slip planes SLPNORCNSLPTL=0DO I=1,NSETISPNOR(1)=NINT(PROPS(25+8*I))ISPNOR(2)=NINT(PROPS(26+8*I))ISPNOR(3)=NINT(PROPS(27+8*I))ISPDIR(1)=NINT(PROPS(28+8*I))ISPDIR(2)=NINT(PROPS(29+8*I))ISPDIR(3)=NINT(PROPS(30+8*I))CALL SLIPSYS (ISPDIR, ISPNOR, NSLIP(I), SLPDIR(1,NSLPTL+1),2 SLPNOR(1,NSLPTL+1), ROTATE)NSLPTL=NSLPTL+NSLIP(I)END DOIF (ND.LT.NSLPTL) THENWRITE (6,*)2 '***ERROR - parameter ND chosen by the present user is less than3 the total number of slip systems NSLPTL'STOPEND IFC----- Slip deformation tensor: SLPDEF (Schmid factors)DO J=1,NSLPTLSLPDEF(1,J)=SLPDIR(1,J)*SLPNOR(1,J)SLPDEF(2,J)=SLPDIR(2,J)*SLPNOR(2,J)SLPDEF(3,J)=SLPDIR(3,J)*SLPNOR(3,J)SLPDEF(4,J)=SLPDIR(1,J)*SLPNOR(2,J)+SLPDIR(2,J)*SLPNOR(1,J)SLPDEF(5,J)=SLPDIR(1,J)*SLPNOR(3,J)+SLPDIR(3,J)*SLPNOR(1,J)SLPDEF(6,J)=SLPDIR(2,J)*SLPNOR(3,J)+SLPDIR(3,J)*SLPNOR(2,J) END DOC----- Initial value of state variables: unit normal to a slip planeC and unit vector in a slip directionCSTATEV(NSTATV)=FLOAT(NSLPTL)DO I=1,NSETSTA TEV(NSTA TV-4+I)=FLOAT(NSLIP(I))END DOIDNOR=3*NSLPTLIDDIR=6*NSLPTLDO J=1,NSLPTLDO I=1,3IDNOR=IDNOR+1STA TEV(IDNOR)=SLPNOR(I,J)IDDIR=IDDIR+1STA TEV(IDDIR)=SLPDIR(I,J)END DOEND DOC----- Initial value of the current strength for all slip systemsCCALL GSLPINIT (STATEV(1), NSLIP, NSLPTL, NSET, PROPS(97))C----- Initial value of shear strain in slip systemsCFIX-- Initial value of cumulative shear strain in each slip systemsDO I=1,NSLPTLSTA TEV(NSLPTL+I)=0.CFIXASTA TEV(9*NSLPTL+I)=0.CFIXBEND DOCFIXASTATEV(10*NSLPTL+1)=0.CFIXBC----- Initial value of the resolved shear stress in slip systemsDO I=1,NSLPTLTERM1=0.DO J=1,NTENSIF (J.LE.NDI) THENTERM1=TERM1+SLPDEF(J,I)*STRESS(J)ELSETERM1=TERM1+SLPDEF(J-NDI+3,I)*STRESS(J) END IFEND DOSTA TEV(2*NSLPTL+I)=TERM1END DOELSEC----- Current stress stateCC----- Copying from the array of state variables STA TVE the following C parameters and variables at current stress state:C Total number of slip systems in all the sets NSLPTLC Number of slip systems in each set NSLIPC Current slip directions SLPDIRC Normals to current slip planes SLPNORCNSLPTL=NINT(STA TEV(NSTATV))DO I=1,NSETNSLIP(I)=NINT(STATEV(NSTATV-4+I))END DOIDNOR=3*NSLPTLIDDIR=6*NSLPTLDO J=1,NSLPTLDO I=1,3IDNOR=IDNOR+1SLPNOR(I,J)=STATEV(IDNOR)IDDIR=IDDIR+1SLPDIR(I,J)=STA TEV(IDDIR)END DOEND DOC----- Slip deformation tensor: SLPDEF (Schmid factors)DO J=1,NSLPTLSLPDEF(1,J)=SLPDIR(1,J)*SLPNOR(1,J)SLPDEF(2,J)=SLPDIR(2,J)*SLPNOR(2,J)SLPDEF(3,J)=SLPDIR(3,J)*SLPNOR(3,J)SLPDEF(4,J)=SLPDIR(1,J)*SLPNOR(2,J)+SLPDIR(2,J)*SLPNOR(1,J)SLPDEF(5,J)=SLPDIR(1,J)*SLPNOR(3,J)+SLPDIR(3,J)*SLPNOR(1,J)SLPDEF(6,J)=SLPDIR(2,J)*SLPNOR(3,J)+SLPDIR(3,J)*SLPNOR(2,J) END DOEND IFC----- Slip spin tensor: SLPSPN (only needed for finite rotation)IF (NLGEOM.NE.0) THENDO J=1,NSLPTLSLPSPN(1,J)=0.5*(SLPDIR(1,J)*SLPNOR(2,J)-2 SLPDIR(2,J)*SLPNOR(1,J))SLPSPN(2,J)=0.5*(SLPDIR(3,J)*SLPNOR(1,J)-2 SLPDIR(1,J)*SLPNOR(3,J))SLPSPN(3,J)=0.5*(SLPDIR(2,J)*SLPNOR(3,J)-2 SLPDIR(3,J)*SLPNOR(2,J))END DOEND IFC----- Double dot product of elastic moduli tensor with the slipC deformation tensor (Schmid factors) plus, only for finiteC rotation, the dot product of slip spin tensor with the stress:C DDEMSDCDO J=1,NSLPTLDO I=1,6DDEMSD(I,J)=0.DO K=1,6DDEMSD(I,J)=DDEMSD(I,J)+D(K,I)*SLPDEF(K,J)END DOEND DOEND DOIF (NLGEOM.NE.0) THENDO J=1,NSLPTLDDEMSD(4,J)=DDEMSD(4,J)-SLPSPN(1,J)*STRESS(1)DDEMSD(5,J)=DDEMSD(5,J)+SLPSPN(2,J)*STRESS(1)IF (NDI.GT.1) THENDDEMSD(4,J)=DDEMSD(4,J)+SLPSPN(1,J)*STRESS(2)DDEMSD(6,J)=DDEMSD(6,J)-SLPSPN(3,J)*STRESS(2)END IFIF (NDI.GT.2) THENDDEMSD(5,J)=DDEMSD(5,J)-SLPSPN(2,J)*STRESS(3)DDEMSD(6,J)=DDEMSD(6,J)+SLPSPN(3,J)*STRESS(3)END IFIF (NSHR.GE.1) THENDDEMSD(1,J)=DDEMSD(1,J)+SLPSPN(1,J)*STRESS(NDI+1)DDEMSD(2,J)=DDEMSD(2,J)-SLPSPN(1,J)*STRESS(NDI+1)DDEMSD(5,J)=DDEMSD(5,J)-SLPSPN(3,J)*STRESS(NDI+1)DDEMSD(6,J)=DDEMSD(6,J)+SLPSPN(2,J)*STRESS(NDI+1) END IFIF (NSHR.GE.2) THENDDEMSD(1,J)=DDEMSD(1,J)-SLPSPN(2,J)*STRESS(NDI+2)DDEMSD(3,J)=DDEMSD(3,J)+SLPSPN(2,J)*STRESS(NDI+2)DDEMSD(4,J)=DDEMSD(4,J)+SLPSPN(3,J)*STRESS(NDI+2)DDEMSD(6,J)=DDEMSD(6,J)-SLPSPN(1,J)*STRESS(NDI+2) END IFIF (NSHR.EQ.3) THENDDEMSD(2,J)=DDEMSD(2,J)+SLPSPN(3,J)*STRESS(NDI+3)DDEMSD(3,J)=DDEMSD(3,J)-SLPSPN(3,J)*STRESS(NDI+3)DDEMSD(4,J)=DDEMSD(4,J)-SLPSPN(2,J)*STRESS(NDI+3)DDEMSD(5,J)=DDEMSD(5,J)+SLPSPN(1,J)*STRESS(NDI+3) END IFEND DOEND IFC----- Shear strain-rate in a slip system at the start of increment:C FSLIP, and its derivative: DFDXSPCID=1DO I=1,NSETIF (I.GT.1) ID=ID+NSLIP(I-1)CALL STRAINRATE (STATEV(NSLPTL+ID), STATEV(2*NSLPTL+ID),2 STA TEV(ID), NSLIP(I), FSLIP(ID), DFDXSP(ID),3 PROPS(65+8*I))END DOC----- Self- and latent-hardening laws。

金属材料专业英语词汇对照

Material Science 材料科学Material Science Definition 材料科学定义Machinability[məʃi:nə'biliti]加工性能Strength .[streŋθ]强度Corrosion & resistance durability.[kə'rəʊʒən] &[ri'zistəns] .[ 'djʊrə'bɪlətɪ] 抗腐蚀及耐用Special metallic features 金属特性Allergic, re-cycling & environmental protection 抗敏感及环境保护Chemical element 化学元素Atom of Elements 元素的原子序数Atom and solid material 原子及固体物质Atom Constitutes 原子的组织图Periodic Table 周期表Atom Bonding 原子键结合Metal and Alloy 金属与合金Ferrous & Non Ferrous Metal 铁及非铁金属Features of Metal 金属的特性Crystal Pattern 晶体结构Crystal structure, Space lattice & Unit cell 晶体结构，定向格子及单位晶格X – ray crystal analytics method X线结晶分析法Metal space lattice 金属结晶格子Lattice constant 点阵常数Mill's Index 米勒指数Metal Phase and Phase Rule金相及相律Solid solution 固熔体Substitutional type solid solution 置换固熔体Interstitial solid solution 间隙固熔体Intermetallic compound 金属间化合物Transformation 转变Transformation Point 转变点Magnetic Transformation 磁性转变Allotropic Transformation 同素转变Thermal Equilibrium 热平衡Degree of freedom 自由度Critical temperature 临界温度Eutectic 共晶Peritectic [.peri’tekti k] Temperature包晶温度Peritectic Reaction 包晶反应Peritectic Alloy 包晶合金Hypoeutectic Alloy 亚共晶体Hypereutectic Alloy 过共晶体Plastic Deformation 金属塑性Slip Plan 滑动面Distortion 畸变Work Hardening 硬化Annealing 退火Crystal Recovery 回复柔软Recrystallization 再结晶Properties & testing of metal 金属材料的性能及试验Chemical Properties 化学性能Physical Properties 物理性能Magnetism 磁性Specific resistivity & specific resistance 比电阻Specific gravity & specific density比重Specific Heat比热热膨胀系数 Coefficient of thermal expansion导热度 Heat conductivity机械性能 Mechanical properties屈服强度(降伏强度) (Yield strength)弹性限度、杨氏弹性系数及屈服点 elastic limit, Young’s module of elasticity to yield point伸长度 Elongation断面缩率 Reduction of area破坏性检验 destructive inspections渗透探伤法 Penetrate inspection磁粉探伤法 Magnetic particle inspection放射线探伤法 Radiographic inspection超声波探伤法 Ultrasonic inspection显微观察法 Microscopic inspection破坏的检验 Destructive Inspection冲击测试 Impact Test疲劳测试 Fatigue Test蠕变试验Creep Test潜变强度 Creeps Strength第一潜变期 Primary Creep第二潜变期 Secondary Creep第三潜变期 Tertiary Creep主要金属元素之物理性质 Physical properties of major Metal Elements工业标准及规格–铁及非铁金属 Industrial Standard – Ferrous & Non – ferrous Metal磁力 Magnetic简介 General软磁 Soft Magnetic硬磁 Hard Magnetic磁场 Magnetic Field磁性感应 Magnetic Induction导磁率[系数,性] Magnetic Permeability磁化率 Magnetic Susceptibility (Xm)磁力(Magnetic Force)及磁场 (Magnetic Field)是因物料里的电子 (Electron)活动而产生抗磁体、顺磁体、铁磁体、反铁磁体及亚铁磁体 Diamagnetism, Paramagnetic, Ferromagnetisms, Antiferromagnetism & Ferrimagnetisms抗磁体 Diamagnetism磁偶极子 Dipole负磁力效应 Negative effect顺磁体 Paramagnetic正磁化率 Positive magnetic susceptibility铁磁体 Ferromagnetism转变元素 Transition element交换能量 Positive energy exchange外价电子 Outer valence electrons化学结合 Chemical bond自发上磁 Spontaneous magnetization磁畴 Magnetic domain相反旋转 Opposite span比较抗磁体、顺磁体及铁磁体 Comparison of Diamagnetism, Paramagnetic & Ferromagnetism反铁磁体 Antiferromagnetism亚铁磁体 Ferrimagnetism磁矩 magnetic moment净磁矩 Net magnetic moment钢铁的主要成份 The major element of steel钢铁用"碳"之含量来分类 Classification of Steel according to Carbon contents铁相 Steel Phases钢铁的名称 Name of steel铁素体Ferrite渗碳体 Cementitle奥氏体 Austenite珠光体及共析钢 Pearlite &Eutectoid奥氏体碳钢 Austenite Carbon Steel单相金属 Single Phase Metal共释变态 Eutectoid Transformation珠光体 Pearlite亚铁释体 Hyppo-Eutectoid初释纯铁体 Pro-entectoid ferrite过共释钢 Hype-eutectoid粗珠光体 Coarse pearlite中珠光体 Medium Pearlite幼珠光体 Fine pearlite磁性变态点 Magnetic Transformation钢铁的制造 Manufacturing of Steel连续铸造法 Continuous casting process电炉 Electric furnace均热炉 Soaking pit全静钢 Killed steel半静钢 Semi-killed steel沸腾钢(未净钢) Rimmed steel钢铁生产流程 Steel Production Flow Chart钢材的熔铸、锻造、挤压及延轧 The Casting, Fogging, Extrusion, Rolling & Steel熔铸 Casting锻造 Fogging挤压 Extrusion延轧Rolling冲剪 Drawing & stamping特殊钢以元素分类Classification of Special Steel according to Element特殊钢以用途来分类 Classification of Special Steel according to End Usage 易车(快削)不锈钢 Free Cutting Stainless Steel含铅易车钢 Leaded Free Cutting Steel含硫易车钢 Sulphuric Free Cutting Steel硬化性能 Hardenability钢的脆性 Brittleness of Steel低温脆性 Cold brittleness回火脆性 Temper brittleness日工标准下的特殊钢材 Specail Steel according to JIS Standard铬钢–日工标准 JIS G4104 Chrome steel to JIS G4104铬钼钢钢材–日工标准 G4105 62 Chrome Molybdenum steel to JIS G4105镍铬–日工标准 G4102 63 Chrome Nickel steel to JIS G4102镍铬钼钢–日工标准 G4103 64 Nickel, Chrome & Molybdenum Steel to JIS G4103高锰钢铸–日工标准 High manganese steel to JIS standard片及板材 Chapter Four-Strip, Steel & Plate冷辘低碳钢片(双单光片)(日工标准 JIS G3141) 73 - 95 Cold Rolled (Low carbon) Steel Strip (to JIS G 3141)简介 General美材试标准的冷辘低碳钢片 Cold Rolled Steel Strip American Standard – American Society for testing and materials (ASTM)日工标准 JIS G3141冷辘低碳钢片 (双单光片)的编号浅释 Decoding of cold rolled(Low carbon)steel strip JIS G3141材料的加工性能 Drawing ability硬度 Hardness表面处理 Surface finish冷辘钢捆片及张片制作流程图表 Production flow chart cold rolled steel coil sheet冷辘钢捆片及张片的电镀和印刷方法 Cold rolled steel coil & sheet electro-plating & painting method冷辘(低碳)钢片的分类用途、工业标准、品质、加热状态及硬度表 End usages, industrial standard, quality, condition and hardness of cold rolled steel strip硬度及拉力 Hardness & Tensile strength test拉伸测试(顺纹测试) Elongation test杯突测试(厚度: 0.4公厘至 1.6公厘，准确至 0.1公厘 3个试片平均数 ) Erichsen test (Thickness: 0.4mm to 1.6mm, figure round up to 0.1mm)曲面(假曲率) Camber厚度及阔度公差 Tolerance on Thickness & Width平坦度(阔度大于 500公厘，标准回火 ) Flatness (width>500mm, temper: standard)弯度 Camber冷辘钢片储存与处理提示 General advice on handling & storage of cold rolled steel coil & sheet 防止生锈 Rust Protection生锈速度表 Speed of rusting焊接 Welding气焊 Gas Welding埋弧焊 Submerged-arc Welding电阻焊 Resistance Welding冷辘钢片(拉力: 30-32公斤/平方米)在没有表面处理状态下的焊接状况 Spot welding conditions for bared (free from paint, oxides etc) Cold rolled mild steel sheets(T/S:30-32 Kgf/ µ m2)时间效应(老化)及拉伸应变 Aging & Stretcher Strains日工标准(JIS G3141)冷辘钢片化学成份 Chemical composition – cold rolled steel sheet to JIS G3141冷辘钢片的"理论重量"计算方程式 Cold Rolled Steel Sheet – Theoretical mass 日工标准(JIS G3141)冷辘钢片重量列表 Mass of Cold-Rolled Steel Sheet to JIS G3141冷辘钢片订货需知Ordering of cold rolled steel strip/sheet其它日工标准冷轧钢片(用途及编号) JIS standard & application of other cold Rolled Special Steel电镀锌钢片或电解钢片Electro-galvanized Steel Sheet/Electrolytic Zinc Coated Steel Sheet电解/电镀锌大大增强钢片的防锈能力Galvanic Action improving Weather & Corrosion Resistance of the Base Steel Sheet上漆能力 Paint Adhesion电镀锌钢片的焊接 Welding of Electro-galvanized steel sheet点焊 Spot welding滚焊 Seam welding电镀锌(电解)钢片 Electro-galvanized Steel Sheet生产流程 Production Flow Chart常用的镀锌钢片(电解片)的基层金属、用途、日工标准、美材标准及一般厚度 Base metal, application, JIS & ASTM standard, and Normal thickness of galvanized steel sheet锌镀层质量 Zinc Coating Mass表面处理 Surface Treatment冷轧钢片 Cold-Rolled Steel Sheet/Strip热轧钢片 Hot-Rolled Sheet/Strip电解冷轧钢片厚度公差 Thickness Tolerance of Electrolytic Cold-rolled sheet热轧钢片厚度公差 Thickness Tolerance of Hot-rolled sheet冷轧或热轧钢片阔度公差 Width Tolerance of Cold or Hot-rolled sheet长度公差 Length Tolerance理论质量 Theoretical Mass锌镀层质量(两个相同锌镀层厚度) Mass Calculation of coating (For equal coating)/MM锌镀层质量(两个不同锌镀层厚度) Mass Calculation of coating (For differential coating)/MM镀锡薄铁片(白铁皮/马口铁) (日工标准 JIS G3303)简介 General镀锡薄铁片的构造 Construction of Electrolytic Tinplate镀锡薄钢片(白铁皮/马日铁)制造过程 Production Process of Electrolytic Tinplate锡层质量 Mass of Tin Coating (JIS G3303-1987)两面均等锡层 Both Side Equally Coated Mass两面不均等锡层 Both Side Different Thickness Coated Mass级别、电镀方法、镀层质量及常用称号Grade, Plating type, Designation of Coating Mass & Common Coating Mass镀层质量标记 Markings & Designations of Differential Coatings硬度 Hardness单相轧压镀锡薄铁片(白铁皮/马口铁) Single-Reduced Tinplate双相辗压镀锡薄钢片(马口铁/白铁皮) Dual-Reduction Tinplate钢的种类 Type of Steel常用尺寸 Commonly Used Size电器用硅 [硅] 钢片 Electrical Steel Sheet简介 General软磁材料 Soft Magnetic Material滞后回线 Narrow Hysteresis矫顽磁力 Coercive Force硬磁材料 Hard Magnetic Material最大能量积 Maximum Energy Product硅含量对电器用的低碳钢片的最大好处 The Advantage of Using Silicon low Carbon Steel晶粒取向(Grain-Oriented)及非晶粒取向(Non-Oriented) Grain Oriented & Non-Oriented电器用硅 [硅] 钢片的最终用途及规格 End Usage and Designations of Electrical Steel Strip电器用的硅 [硅] 钢片之分类 Classification of Silicon Steel Sheet for Electrical Use电器用钢片的绝缘涂层 Performance of Surface Insulation of Electrical Steel Sheets晶粒取向电器用硅钢片主要工业标准 International Standard – Grain-Oriented Electrical Steel Silicon Steel Sheet for Electrical Use晶粒取向电器用硅钢片 Grain-Oriented Electrical Steel晶粒取向，定取向芯钢片及高硼定取向芯钢片之磁力性能及夹层系数 (日工标准及美材标准) Magnetic Properties and Lamination Factor of SI-ORIENT-CORE& SI-ORIENT-CORE-HI B Electrical Steel Strip (JIS and AISI Standard)退火 Annealing电器用钢片用家需自行应力退火原因 Annealing of the Electrical Steel Sheet退火时注意事项 Annealing Precautionary碳污染 Prevent Carbon Contamination热力应先从工件边缘透入 Heat from the Laminated Stacks Edges提防过份氧化 No Excessive Oxidation应力退火温度 Stress –relieving Annealing Temperature绝缘表面 Surface Insulation非晶粒取向电力用钢片的电力、磁力、机械性能及夹层系数 Lamination Factors of Electrical, Magnetic & Mechanical Non-Grain Oriented Electrical电器及家电外壳用镀层冷辘 [低碳] 钢片 Coated (Low Carbon) Steel Sheets for Casing,Electricals & Home Appliances镀铝硅钢片 Aluminized Silicon Alloy Steel Sheet镀铝硅合金钢片的特色 Feature of Aluminized Silicon Alloy Steel Sheet用途 End Usages抗化学品能力 Chemical Resistance镀铝(硅)钢片–日工标准 (JIS G3314) Hot-aluminum-coated sheets and coils to JIS G 3314镀铝(硅)钢片–美材试标准 (ASTM A-463-77)35.7 JIS G3314镀热浸铝片的机械性能 Mechanical Properties of JIS G 3314 Hot-Dip Aluminum-coated Sheets andCoils公差 Size Tolerance镀铝(硅)钢片及其它种类钢片的抗腐蚀性能比较 Comparsion of various resistance of aluminized steel & other kinds of steel镀铝(硅)钢片生产流程 Aluminum Steel Sheet, Production Flow Chart焊接能力 Weldability镀铝钢片的焊接状态(比较冷辘钢片) Tips on welding of Aluminized sheet in comparasion with cold rolled steel strip钢板 Steel Plate钢板用途分类及各国钢板的工业标准包括日工标准及美材试标准 Type of steel Plate & Related JIS, ASTM and Other Major Industrial Standards钢板生产流程 Production Flow Chart钢板订货需知 Ordering of Steel Plate不锈钢 Stainless Steel不锈钢的定义 Definition of Stainless Steel不锈钢之分类，耐腐蚀性及耐热性 Classification, Corrosion Resistant & Heat Resistance of Stainless Steel铁铬系不锈钢片Chrome Stainless Steel马氏体不锈钢Martensite Stainless Steel低碳马氏体不锈钢Low Carbon Martensite Stainless Steel含铁体不锈钢Ferrite Stainless Steel镍铬系不锈钢Nickel Chrome Stainless Steel释出硬化不锈钢Precipitation Hardening Stainless Steel铁锰铝不锈钢Fe / Mn / Al / Stainless Steel不锈钢的磁性Magnetic Property & Stainless Steel不锈钢箔、卷片、片及板之厚度分类Classification of Foil, Strip, Sheet & Plate by Thickness表面保护胶纸Surface protection film不锈钢片材常用代号Designation of SUS Steel Special Use Stainless 表面处理 Surface finish 薄卷片及薄片(0.3至 2.9mm 厚之片)机械性能Mechanical Properties of Thin Stainless Steel(Thickness from 0.3mm to 2.9mm) – strip/sheet 不锈钢片机械性能(301, 304, 631, CSP) Mechanical Properties of Spring use Stainless Steel不锈钢–种类，工业标准，化学成份，特点及主要用途Stainless Steel – Type, Industrial Standard, Chemical Composition, Characteristic & end usage of the most commonly used Stainless Steel不锈钢薄片用途例End Usage of Thinner Gauge不锈钢片、板用途例Examples of End Usages of Strip, Sheet & Plate不锈钢应力退火卷片常用规格名词图解General Specification of Tension Annealed Stainless Steel Strips耐热不锈钢Heat-Resistance Stainless Steel镍铬系耐热不锈钢特性、化学成份、及操作温度Heat-Resistance Stainless Steel铬系耐热钢Chrome Heat Resistance Steel镍铬耐热钢Ni - Cr Heat Resistance Steel超耐热钢Special Heat Resistance Steel抗热超级合金Heat Resistance Super Alloy耐热不锈钢比重表Specific Gravity of Heat – resistance steel plates and sheets stainless steel不锈钢材及耐热钢材标准对照表Stainless and Heat-Resisting Steels发条片 Power Spring Strip发条的分类及材料 Power Spring Strip Classification and Materials上链发条 Wind-up Spring倒后擦发条 Pull Back Power Spring圆面("卜竹")发条 Convex Spring Strip拉尺发条 Measure Tape魔术手环 Magic Tape魔术手环尺寸图 Drawing of Magic Tap定型发条 Constant Torque Spring定型发条及上炼发条的驱动力 Spring Force of Constant Torque Spring and Wing-up Spring定型发条的形状及翻动过程 Shape and Spring Back of Constant Torque Spring定型发条驱动力公式及代号The Formula and Symbol of Constant Torque Spring边缘处理 Edge Finish硬度 Hardness高碳钢化学成份及用途 High Carbon Tool Steel, Chemical Composition and Usage每公斤发条的长度简易公式 The Length of 1 Kg of Spring Steel Strip SK-5 & AISI-301每公斤长的重量 /公斤(阔 100-200公厘) Weight per one meter long (kg) (Width 100-200mm) SK-5 & AISI-301每公斤之长度 (阔 100-200公厘) Length per one kg (Width 100-200mm) SK-5 & AISI-301每公尺长的重量 /公斤(阔 2.0-10公厘) Weight per one meter long (kg) (Width 2.0-10mm) SK-5 & AISI-301每公斤之长度 (阔 2.0-10公厘) Length per one kg (Width 2.0-10mm)高碳钢片 High Carbon Steel Strip分类 Classification用组织结构分类 Classification According to Grain Structure用含碳量分类–即低碳钢、中碳钢及高碳钢 Classification According to Carbon Contains弹簧用碳钢片 Carbon Steel Strip For Spring Use冷轧状态 Cold Rolled Strip回火状态 Annealed Strip淬火及回火状态 Hardened & Tempered Strip/ Precision – Quenched Steel Strip贝氏体钢片 Bainite Steel Strip弹簧用碳钢片材之边缘处理 Edge Finished淬火剂 Quenching Media碳钢回火 Tempering回火有低温回火及高温回火 Low & High Temperature Tempering高温回火 High Temperature Tempering退火 Annealing完全退火 Full Annealing扩散退火 Diffusion Annealing低温退火 Low Temperature Annealing中途退火 Process Annealing球化退火 Spheroidizing Annealing光辉退火 Bright Annealing淬火 Quenching时间淬火 Time Quenching奥氏铁孻回火 Austempering马氏铁体淬火 Marquenching高碳钢片用途 End Usage of High Carbon Steel Strip冷轧高碳钢–日本工业标准 Cold-Rolled (Special Steel) Carbon Steel Strip to JIS G3311电镀金属钢片 Plate Metal Strip电镀金属捆片的优点Advantage of Using Plate Metal Strip金属捆片电镀层 Plated Layer of Plated Metal Strip镀镍 Nickel Plated镀铬 Chrome Plated镀黄铜 Brass Plated基层金属 Base Metal of Plated Metal Strip低碳钢或铁基层金属 Iron & Low Carbon as Base Metal不锈钢基层金属 Stainless Steel as Base Metal铜基层金属 Copper as Base Metal黄铜基层金属 Brass as Base Metal轴承合金 Bearing Alloy轴承合金–日工标准 JIS H 5401 Bearing Alloy to JIS H 5401锡基、铅基及锌基轴承合金比较表 Comparison of Tin base, Lead base and Zinc base alloy for Bearing purpose易溶合金 Fusible Alloy焊接合金 Soldering and Brazing Alloy软焊 Soldering Alloy软焊合金–日本标准 JIS H 4341 Soldering Alloy to JIS H 4341硬焊 Brazing Alloy其它焊接材料请参阅日工标准目录 Other Soldering Material细线材、枝材、棒材 Chapter Five Wire, Rod & Bar线材/枝材材质分类及制成品 Classification and End Products of Wire/Rod铁线(低碳钢线)日工标准 JIS G 3532 Low Carbon Steel Wires ( Iron Wire ) to JIS G 3532光线(低碳钢线)，火线 (退火低碳钢线 )，铅水线 (镀锌低碳钢线)及制造钉用低碳钢线之代号、公差及备注 Ordinary Low Carbon Steel Wire, Annealed Low Carbon Steel Wire, Galvanized low Carbon Steel Wire & Low Carbon Steel Wire for nail manufacturing - classification, Symbol of Grade, Tolerance and Remarks.机械性能 Mechanical Properites锌包层之重量，铜硫酸盐试验之酸洗次数及测试用卷筒直径 Weight of Zinc-Coating, Number of Dippings in Cupric Sulphate Test and Diameters of Mandrel Used for Coiling Test冷冲及冷锻用碳钢线枝 Carbon Steel Wire Rods for Cold Heading & Cold Forging (to JIS G3507) 级别，代号及化学成份 Classification, Symbol of Grade and Chemical Composition直径公差，偏圆度及脱碳层的平均深度 Diameter Tolerance, Ovality and Average Decarburized Layer Depth冷拉钢枝材 Cold Drawn Carbon Steel Shafting Bar枝材之美工标准，日工标准，用途及化学成份 AISI, JIS End Usage and Chemical Composition of Cold Drawn Carbon Steel Shafting Bar冷拉钢板重量表 Cold Drawn Steel Bar Weight Table高碳钢线枝 High Carbon Steel Wire Rod (to JIS G3506)冷拉高碳钢线 Hard Drawn High Carbon Steel Wire (to JIS G3521, ISO-84580-1&2)化学成份分析表 Chemical Analysis of Wire Rod线径、公差及机械性能(日本工业标准 G 3521) Mechanical Properties (JIS G 3521)琴线(日本标准 G3522) Piano Wires (to G3522)级别，代号，扭曲特性及可用之线材直径 Classes, symbols, twisting characteristic and applied WireDiameters直径，公差及拉力强度 Diameter, Tolerance and Tensile Strength裂纹之容许深度及脱碳层 Permissible depth of flaw and decarburized layer常用的弹簧不锈钢线-编号，特性，表面处理及化学成份 Stainless Spring Wire – National Standard number, Characteristic, Surface finish & Chemical composition弹簧不锈钢线，线径及拉力列表Stainless Spring Steel, Wire diameter and Tensile strength of Spring Wire处理及表面状况 Finish & Surface各种不锈钢线在不同处理拉力比较表 Tensile Strength of various kinds of Stainless Steel Wire under Different Finish圆径及偏圆度之公差 Tolerance of Wire Diameters & Ovality铬镍不锈钢及抗热钢弹簧线材–美国材验学会 ASTM A313 – 1987 Chromium – Nickel Stainless and Heat-resisting Steel Spring Wire – ASTM A313 – 1987化学成份 Chemical Composition机械性能 Mechanical Properties305, 316, 321及 347之拉力表 Tensile Strength Requirements for Types 305, 316, 321 and 347 A1S1-302贰级线材之拉力表 Tensile Strength of A1S1-302 Wire日本工业标准–不锈钢的化学成份 (先数字后字母排列) JIS –Chemical Composition of Stainless Steel (in order of number & alphabet)美国工业标准–不锈钢及防热钢材的化学成份 (先数字后字母排列) AISI – Chemical Composition of Stainless Steel & Heat-Resistant Steel(in order of number & alphabet)易车碳钢 Free Cutting Carbon Steels (to JIS G4804 )化学成份 Chemical composition圆钢枝，方钢枝及六角钢枝之形状及尺寸之公差 Tolerance on Shape and Dimensions for Round Steel Bar, Square Steel Bar, Hexagonal Steel Bar易车(快削)不锈钢 Free Cutting Stainless Steel易车(快削)不锈钢种类 Type of steel易车(快削)不锈钢拉力表 Tensile Strength of Free Cutting Wires枝/棒无芯磨公差表 (μ) (μ = 1/100 mm) Rod/Bar Centreless Grind Tolerance易车不锈钢及易车钢之不同尺寸及硬度比较 Hardness of Different Types & Size of Free Cutting Steel 扁线、半圆线及异形线 Flat Wire, Half Round Wire, Shaped Wire and Precision Shaped Fine Wire 加工方法 Manufacturing Method应用材料 Material Used特点 Characteristic用途End Usages不锈钢扁线及半圆线常用材料 Commonly used materials for Stainless Flat Wire & Half Round Wire 扁线公差 Flat Wire Tolerance方线公差 Square Wire Tolerance。

晶体模拟退火英语

晶体模拟退火英语Crystal Simulated Annealing in EnglishSimulated annealing is a computational technique for finding an approximation to the global minimum (or maximum) of a given function. The method is based on the process of annealing in metallurgy, where a material is heated and then slowly cooled to decrease the defects in its structure. In the context of optimization, the "temperature" is a parameter that controls the probability of accepting a new solutionthat is worse than the current one.In the field of crystallography, simulated annealing is used to predict the structure of crystals from diffraction data. The process involves the following steps:1. Initialization: Start with a random configuration of atoms in the crystal lattice.2. Heating: Raise the "temperature" to a high value, which allows for large changes in the atomic positions. This step is crucial because it allows the system to escape from local minima.3. Cooling Schedule: Gradually decrease the "temperature" over time. As the temperature drops, the system becomes more likely to settle into lower-energy configurations.4. Energy Evaluation: At each step, calculate the energy of the system based on the current atomic positions. This energy function is a mathematical representation of the stability of the crystal structure.5. Acceptance Criterion: Decide whether to accept a new configuration based on the Metropolis criterion, which compares the energy of the new state with the current state and the current "temperature".6. Convergence: Continue the process until the system reaches a state of minimum energy, indicating that a stable crystal structure has been found.Simulated annealing is particularly useful in crystal structure prediction because it can navigate the complex energy landscapes typical of crystalline materials. It allows for the exploration of a vast number of possible configurations, increasing the chances of finding the true global minimum, which corresponds to the most stable and accurate crystal structure.The technique has been successfully applied to a wide range of materials, from simple inorganic crystals to complex organic compounds and proteins. It remains a powerful tool in the arsenal of computational chemists and materials scientists, contributing to the advancement of our understanding and design of new materials.。

6H_SiC单晶的生长与缺陷

第32卷第3期硅酸盐学报Vol.32,No.3　2004年3月J OURNAL OF THE CHINESE CERAMIC SOCIETY M a r c h,20046H SiC单晶的生长与缺陷胡小波,徐现刚,王继扬,韩荣江,董　捷,李现祥,蒋民华(山东大学,晶体材料国家重点实验室,济南　250100)摘　要:采用升华法,在一定的温度、气体压力和流量的条件下,生长了尺寸<50.8mm的6H SiC单晶。

利用光学显微术观察了原生晶体的表面形貌,发现了微管在晶体表面的露头点具有明显的多个螺位错成核特征。

采用透射模式对抛光晶片进行观察,发现了SiC晶体内的典型缺陷,如:负晶、微管、碳颗粒等,并对它们的形成机理进行了讨论。

关键词:6H碳化硅;微管;负晶;缺陷中图分类号:O771;O782文献标识码:A文章编号:04545648(2004)03024803GR OWTH AN D DEFECTS OF6H SiC MON OCR YSTALSHU Xiaobo,XU Xiangang,W A N G Jiyang,HA N Rongjiang,DON G Jie,L I Xianxiang,J IA N G Minhua (State K ey Laboratory of Crystal Materials,Shandong University,Jinan　250100,China)Abstract:6H SiC monocrystals with the diameter of50.8mm were synthesized by sublimation method.The surface morpho2logy of as2grown SiC crystal was observed by optical microscopy.It is found that outcrops of micropipes on the crystal surface possess remark2 able feature of double or multiple screw dislocations.In addition,polished SiC wafer was also examined by optical microscopy with transmission mode,and it is found that there are some typical defects,such as negative crystals,micropipes,carbon particles in SiC crystals.The formation mechanisms of these defects are discussed.K ey w ords:6H silicon carbide;micropipe;negative crystal;defect SiC是一种重要的半导体材料,它具有优良的热学、力学、化学和电学性质,可以用作蓝光L ED (light emitting diode)衬底材料,同时又是制作高温、高频、大功率电子器件的最佳材料之一[1,2]。

聚乙醇酸退火工艺研究

第３５卷第１期化㊀学㊀研㊀究Ｖｏｌ．３５㊀Ｎｏ．１２０２４年１月ＣＨＥＭＩＣＡＬ㊀ＲＥＳＥＡＲＣＨＪａｎ．２０２４聚乙醇酸退火工艺研究朱爱臣１，李文明１，马丽霞１，朱肖杰１，石㊀锐２∗，李㊀振３，刘㊀阳１∗（１．山东省药学科学院山东省医用高分子材料重点实验室，山东济南２５０１０１；２．北京积水潭医院北京市创伤骨科研究所，北京１０００３５；３．中石油华东设计院有限公司，山东青岛２６６０７１）收稿日期：２０２２⁃１２⁃２１基金项目：山东省科技型中小企业创新能力提升工程项目（２０２２ＴＳＧＣ１０３０），北京市卫生健康委员会项目（ＰＸＭ２０２０－０２６２７５－０００００３）作者简介：朱爱臣（１９８４－），女，高级工程师，研究方向为医用高分子材料的加工工艺研究㊂∗通信作者，Ｅ⁃ｍａｉｌ：138****0830＠１６３．ｃｏｍ摘㊀要：聚乙醇酸（ＰＧＡ）热加工过程中，因冷却速率过快而存在内应力，需要进行退火处理，本文研究了退火温度对ＰＧＡ性能的影响㊂采用差示扫描量热仪（ＤＳＣ）和热台偏光显微镜（ＰＯＭ）研究了ＰＧＡ的冷结晶性及结晶结构，采用Ｘ射线衍射仪（ＸＲＤ）㊁ＤＳＣ和万能材料试验机研究了不同退火温度对ＰＧＡ的结晶性㊁热性能及力学性能的影响㊂实验结果表明，ＰＧＡ在熔融冷却过程中先形成针状晶核，并逐渐成长为球晶，且随着降温速率的增加，冷结晶温度降低，结晶时长变短㊂随退火温度的升高，ＰＧＡ的松弛峰值温度升高，松弛焓增大，应变减小，弹性模量增大，结晶度和应力先增大后减小㊂关键词：聚乙醇酸；内应力；退火；结晶；热性能中图分类号：ＴＱ３２０．６６文献标志码：Ａ文章编号：１００８－１０１１（２０２４）０１－００６８－０６ＥｆｆｅｃｔｏｆａｎｎｅａｌｉｎｇｐｒｏｃｅｓｓｏｎｔｈｅｐｒｏｐｅｒｔｉｅｓｏｆＰＧＡＺＨＵＡｉｃｈｅｎ１ＬＩＷｅｎｍｉｎｇ１ＭＡＬｉｘｉａ１ＺＨＵＸｉａｏｊｉｅ１ＳＨＩＲｕｉ２∗ ＬＩＺｈｅｎ３ＬＩＵＹａｎｇ１∗１．ＳｈａｎｄｏｎｇＫｅｙＬａｂｏｒａｔｏｒｙｏｆＭｅｄｉｃａｌＰｏｌｙｍｅｒＭａｔｅｒｉａｌｓＳｈａｎｄｏｎｇＡｃａｄｅｍｙｏｆＰｈａｒｍａｃｅｕｔｉｃａｌＳｃｉｅｎｃｅｓＪｉｎａｎ２５０１０１ＳｈａｎｄｏｎｇＣｈｉｎａ２．ＢｅｉｊｉｎｇＲｅｓｅａｒｃｈＩｎｓｔｉｔｕｔｅｏｆＴｒａｕｍａｔｏｌｏｇｙａｎｄＯｒｔｈｏｐａｅｄｉｃｓＢｅｉｊｉｎｇＪｉｓｈｕｉｔａｎＨｏｓｐｉｔａｌＢｅｉｊｉｎｇ１０００３５Ｃｈｉｎａ３．ＣＮＰＣＥａｓｔＣｈｉｎａＤｅｓｉｇｎＩｎｓｔｉｔｕｔｅＣＯ．ＬＴＤＱｉｎｇｄａｏ２６６０７１ＳｈａｎｄｏｎｇＣｈｉｎａＡｂｓｔｒａｃｔＰＧＡｈａｓｉｎｔｅｒｎａｌｓｔｒｅｓｓｉｎｔｈｅｍｅｌｔ－ｃｏｏｌｉｎｇｐｒｏｃｅｓｓ，ａｎｄｎｅｅｄｓｔｏｂｅａｎｎｅａｌｅｄ，ｂｅｃａｕｓｅｔｈｅｃｏｏｌｉｎｇｒａｔｅｉｓｔｏｏｆａｓｔ．ＴｈｅｅｆｆｅｃｔｏｆａｎｎｅａｌｉｎｇｔｅｍｐｅｒａｔｕｒｅｏｎｔｈｅｐｒｏｐｅｒｔｉｅｓｏｆＰＧＡｗａｓｓｔｕｄｉｅｄ．ＴｈｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｐｒｏｐｅｒｔｉｅｓａｎｄｓｔｒｕｃｔｕｒｅｏｆＰＧＡｗｅｒｅｉｎｖｅｓｔｉｇａｔｅｄｂｙＤＳＣａｎｄＰＯＭ．Ｔｈｅｅｆｆｅｃｔｓｏｆｄｉｆｆｅｒｅｎｔａｎｎｅａｌｉｎｇｔｅｍｐｅｒａｔｕｒｅｓｏｎｔｈｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎ，ｔｈｅｒｍａｌｐｒｏｐｅｒｔｉｅｓａｎｄｍｅｃｈａｎｉｃａｌｐｒｏｐｅｒｔｉｅｓｏｆＰＧＡｗｅｒｅｉｎｖｅｓｔｉｇａｔｅｄｂｙＸＲＤ，ＤＳＣａｎｄｕｎｉｖｅｒｓａｌｍａｔｅｒｉａｌｔｅｓｔｉｎｇｍａｃｈｉｎｅ．ＴｈｅｅｘｐｅｒｉｍｅｎｔａｌｒｅｓｕｌｔｓｓｈｏｗｔｈａｔＰＧＡｆｉｒｓｔｆｏｒｍｓａｃｉｃｕｌａｒｃｒｙｓｔａｌｎｕｃｌｅｉ，ａｎｄｔｈｅｎｇｒｏｗｓｉｎｔｏｓｐｈｅｒｕｌｉｔｅｇｒａｄｕａｌｌｙｄｕｒｉｎｇｔｈｅｍｅｌｔ－ｃｏｏｌｉｎｇｐｒｏｃｅｓｓ．Ｗｉｔｈｔｈｅｉｎｃｒｅａｓｅｏｆｃｏｏｌｉｎｇｒａｔｅ，ｔｈｅｃｏｌｄｃｒｙｓｔａｌｌｉｚａｔｉｏｎｔｅｍｐｅｒａｔｕｒｅａｎｄｔｈｅｃｒｙｓｔａｌｌｉｚａｔｉｏｎｔｉｍｅｄｅｃｒｅａｓｅ．Ｗｉｔｈｔｈｅｉｎｃｒｅａｓｅｏｆａｎｎｅａｌｉｎｇｔｅｍｐｅｒａｔｕｒｅ，ｔｈｅｒｅｌａｘａｔｉｏｎｔｅｍｐｅｒａｔｕｒｅａｎｄｔｈｅｒｅｌａｘａｔｉｏｎｅｎｔｈａｌｐｙｏｆＰＧＡｉｎｃｒｅａｓｅ，ｔｈｅｓｔｒａｉｎｄｅｃｒｅａｓｅｓ，ａｎｄｔｈｅｅｌａｓｔｉｃｍｏｄｕｌｕｓｉｎｃｒｅａｓｅｓ，ｗｈｉｌｅｔｈｅｃｒｙｓｔａｌｌｉｎｉｔｙａｎｄｔｈｅｍａｘｉｍｕｍｓｔｒｅｓｓｉｎｃｒｅａｓｅｆｉｒｓｔａｎｄｔｈｅｎｄｅｃｒｅａｓｅ．Ｋｅｙｗｏｒｄｓ：ＰＧＡ；ｉｎｎｅｒ－ｓｔｒｅｓｓ；ａｎｎｅａｌｉｎｇ；ｃｒｙｓｔａｌｌｉｚａｔｉｏｎ；ｔｈｅｒｍａｌｐｒｏｐｅｒｔｙ㊀㊀聚乙醇酸（又名聚乙交酯，ＰＧＡ）是一种分子结构类似于聚乳酸（又名聚丙交酯，ＰＬＡ）的可降解脂肪族聚酯，主要应用在缝合线［１］㊁药物缓释［２－３］㊁组织工程支架［４－５］㊁补片［６－７］㊁可吸收结扎夹［８］等医疗器械领域㊂ＰＧＡ的分子结构简单规整，没有ＰＬＡ分子结构中的侧甲基，其晶体结构为平面之字形，类似于聚乙烯（ＰＥ）［９］，为致密型结晶性聚合物，故ＰＧＡ很难溶于普通有机溶剂㊂ＰＧＡ流体为非牛顿假塑性流体，热加工性能较好，属于典型的热塑性聚第１期朱爱臣等：聚乙醇酸退火工艺研究６９㊀合物，适用于挤出㊁模压㊁纺丝㊁注塑㊁吹塑等加工方式㊂由于ＰＧＡ的熔点较高，热分解温度与熔融温度相差不大，热加工区间较窄，需要严格控制加工条件，ＰＧＡ精密注塑工艺已报道［１０］㊂ＰＧＡ熔融冷却加工过程中，其结晶性影响着产品的物理性能，而关于ＰＧＡ冷却结晶的研究较少．喻祖圣等［１１］研究了ＰＧＡ非等温结晶行为，但是在ＰＧＡ热加工过程中，其降温速率远远高于研究中的降温速率２０ħ／ｍｉｎ．李连明等［１２］研究了不同分子量ＰＧＡ的结晶性能，但是其研究的ＰＧＡ的黏度最大为０．７，而通过热加工成型的ＰＧＡ其黏度需要大于１．０，本文中的ＰＧＡ的黏度为１．６１．７㊂本文通过研究高分子量的ＰＧＡ在不同降温速率下的结晶性，研究了ＰＧＡ热加工产品的退火工艺㊂对于高分子纤维制品及工程高分子制品［１３－１４］，退火处理是一种常用的消除内应力的方法，本文研究了不同退火温度对ＰＧＡ结晶性㊁热性能及力学性能的影响㊂１㊀材料与方法１．１㊀材料与仪器㊀㊀聚乙醇酸，参考文献［１５］采用已交酯开环聚合的方法自制，特性黏数为１．６７㊂真空干燥箱：ＤＺＦ型，北京市永光明医疗仪器厂；电子天平，ＢＳＡ２２４Ｓ型，赛多利斯科学仪器（北京）有限公司㊂差示扫描量热仪，ＤＳＣ－６０型，日本岛津㊂热台偏光显微镜，ＡｘｉｏＳｃｏｐｅＡ１ｐｏｌ型，德国ＺＥＩＳＳ公司；万能材料试验机，ＡＧＳ－Ｈ型，日本岛津；注塑机，ＳＳＦ２５０型，宁波圣特龙塑料机械有限公司㊂１．２㊀样品制备ＰＧＡ样品在真空烘箱中进行退火处理，退火温度分别设为：１２０㊁１５０㊁１８０和１９０ħ，退火时间均为１ｈ，真空度高于０．０９ＭＰａ㊂１．３㊀测试方法非等温结晶性：采用日本岛津ＤＳＣ－６０型差示扫描量热仪进行测试㊂取５１０ｍｇＰＧＡ置于铝坩埚中压片制样，以１０ħ／ｍｉｎ的升温速率从室温升至２４０ħ恒温３ｍｉｎ，然后分别以１０㊁２０㊁５０和９０ħ／ｍｉｎ的降温速率降至室温，并记录降温曲线㊂热台偏光显微镜：采用德国ＺＥＩＳＳ公司ＡｘｉｏＳｃｏｐｅＡ１ｐｏｌ型正立偏光显微镜观察㊂取１０１５ｍｇＰＧＡ置于载玻片上，并在２４０ħ的热台上加热，在氮气气氛下熔融３ｍｉｎ后，再覆盖盖玻片，然后加液氮快速降温至１５０ħ，观察ＰＧＡ在１５０ħ时的结晶状况㊂热性能：采用日本岛津ＤＳＣ－６０型差示扫描量热仪进行测试㊂取样品５１０ｍｇ，以１０ħ／ｍｉｎ的升温速率从室温升温至２４０ħ，并记录热流变化曲线㊂Ｘ射线衍射（ＸＲＤ）：采用Ｃｕ（Ｋα）靶辐射，Ｎｉ片滤波，λ＝１．５４０６ˑ１０－１０ｍ，扫描范围２θ＝５ʎ ５０ʎ，扫描步距２θ＝０．０２ʎ，扫描速度５ʎ／ｍｉｎ㊂力学性能：按标准ＧＢ／Ｔ１０４０．２－２００６，采用日本岛津ＡＧＳ－Ｈ型万能材料试验机测试哑铃型１Ａ型样条的拉伸曲线，测试５个平行样，取抗张强度中间值样条绘制拉伸曲线，试验载荷为５００Ｎ，拉伸速率为１００ｍｍ／ｍｉｎ㊂２㊀结果与讨论２．１㊀ＰＧＡ非等温结晶性㊀㊀ＰＧＡ熔融冷却定型过程中，伴随着复杂的结晶行为，研究ＰＧＡ从熔融态在不同降温速率下的ＤＳＣ曲线（见图１），对控制ＰＧＡ的结晶度和调节其力学性能均有十分重要的参考价值㊂图１㊀ＰＧＡ非等温结晶ＤＳＣ曲线Ｆｉｇ．１㊀ＮｏｎｉｓｏｔｈｅｒｍａｌｃｒｙｓｔａｌｌｉｚａｔｉｏｎｃｕｒｖｅｓｏｆＰＧＡ由图１可知，ＰＧＡ在熔融冷却过程中，形成单一放热结晶峰，且随着降温速率的增加，结晶峰向低温移动㊂不同降温速率的结晶峰值温度（ｔｐ）㊁结晶时长（Δｔ）㊁结晶焓变（ΔＨ）见表１㊂由表１可知，随降温速率的增大，结晶峰值温度降低，结晶时长减短，结晶焓变增大㊂这是因为聚合物分子链在熔融态为自由伸展的无规线团，随着温度的降低，分子链由杂乱无序的排列调整为有序的规整晶格结构，而分子链的运动具有温度依赖性，温度越低，分子链运动能力越弱㊂降温速率越大，分子链在高温区的时间越短，构象调整的速度越慢，分子链被冻结的越快，达到结晶峰值的温度越低，结晶时长越短，结晶７０㊀化㊀学㊀研㊀究２０２４年结构的完整性越差㊂表１㊀ＰＧＡ非等温结晶参数Ｔａｂｌｅ１㊀ＮｏｎｉｓｏｔｈｅｒｍａｌｃｒｙｓｔａｌｌｉｚａｔｉｏｎｐａｒａｍｅｔｅｒｓｏｆＰＧＡ降温速率／（ħ／ｍｉｎ）ｔｐ／ħΔｔ／ｍｉｎΔＨ／（Ｊ／ｇ）１０１８３．９２．８－６０．１２０１８０．１１．６－６２．５５０１７２．１０．６－６３．７９０１６１．８０．４－６５．３㊀㊀注：ｔｐ：冷结晶峰值温度，Δｔ：结晶时长，ΔＨ：冷结晶焓２．２㊀ＰＧＡ的结晶形态ＰＧＡ熔融后，用液氮降温，观察降温过程中ＰＧＡ的结晶形态（见图２）㊂由ＰＧＡ非等温结晶ＤＳＣ曲线可知，即使降温速率达到９０ħ／ｍｉｎ，ＰＧＡ在１６１．８ħ也会形成结晶峰，这与通过偏光显微镜观察ＰＧＡ在降温过程中的结晶形态的变化相一致㊂由图２可知，ＰＧＡ熔融冷却降温至１８０ħ时，形成大量的长径比比较大的针状晶核，晶核长约１０μｍ左右；降温至１７０ħ时，针状晶核诱导结晶，形成少量的晶体，晶体尺寸大约在１５μｍ左右，说明ＰＧＡ晶体是沿径向生长的层状晶体；降温至１６０ħ时，视野较亮，形成大量的椭圆形或圆形晶体，晶体尺寸大约在２０μｍ左右；降温至１５０ħ时，晶体尺寸大约在４０μｍ左右㊂因此，ＰＧＡ熔融冷却结晶过程中，随着温度的降低和时间的延长，晶体尺寸增大，视野亮度增大，形成了具有明显双折射性和对称性的球晶，在偏光显微镜下呈现的是典型的Ｍａｌｔｅｓｅ十字消光现象㊂图２㊀ＰＧＡ的偏光显微镜照片Ｆｉｇ．２㊀ＰｏｌａｒｉｚｉｎｇｍｉｃｒｏｓｃｏｐｅｐｈｏｔｏｇｒａｐｈｏｆＰＧＡ２．３㊀ＰＧＡ退火处理后的热性能ＰＧＡ热加工过程中，由熔融态瞬间冷却定型后，分子链处于热不平衡态，存在一定内应力，需要对产品进行退火处理，本文研究了ＰＧＡ退火前后的热性能变化曲线（见图３）㊂由图３可知，未退火样品在４０６０ħ有一个放热结晶峰，而退火后的样品均没有放热结晶峰，说明ＰＧＡ在熔融冷却过程中不能形成完整的结晶结构，因为降温速率过快，分子链被迅速冻结，晶体结构存在一定缺陷㊂ＰＧＡ退火处理后完善了其结晶结构㊂由图３还可看出，在不同温度下退火的样品，在熔融峰之前均存在一个吸热峰，我们定义为松弛峰［１６］㊂ＰＧＡ完全结晶的熔融焓为１３９Ｊ／ｇ［１７］，由ＤＳＣ曲线得出松弛峰和熔融峰的温度和焓变，由此计算不同退火温度样品的结晶度见表２㊂由表２可知，随退火温度的升高，松弛峰值温度（ｔ）升高，松弛焓（ΔＨ）增大，而熔融温度（ｔｍ）没有明显变化，熔融焓（ΔＨｍ）和结晶度先增大后减小㊂由此可以推测，松弛峰对应的是样品中无规部分分子局部有序排列的微晶熔融吸热峰，而熔融峰对应的是样品中结晶部分的熔融吸热峰㊂未退火的样品内非晶部分分子链被冻结后无法形成热平衡状态，样品退火时，被冻结的分子链可以局部运动，形成局部有序的微晶态，随退火温度的升高，分子链运动幅度增大，形成的微晶尺寸增大，松弛峰的熔融焓增大㊂未退火的样品内结晶区结晶结构存在一定缺陷，分子链来不及调节即被冻结，而样品退火时，被冻结的这部分分子链可以局部自由运动，形成完整的结晶第１期朱爱臣等：聚乙醇酸退火工艺研究７１㊀结构，故随退火温度的升高，结晶区对应的熔融温度无明显变化，熔融焓增大，而过高的退火温度，分子链运动幅度过大，不利于形成结晶结构，故熔融焓减小，由此计算的结晶度先增大后减小㊂２．４㊀ＰＧＡ的Ｘ射线衍射图从ＰＧＡ退火前后样品的ＸＲＤ图谱（见图４）可以看出，ＰＧＡ为半结晶结构，２θ在２２．３ʎ和２８．９ʎ附近有尖锐的衍射峰，他们对应的晶面指数分别为１１０和０２０，由图谱结晶峰面积与总面积的比值计算样品的结晶度（见表３）㊂由表３可知，退火后样品的结晶度比退火前明显提高，而随退火温度的升高，结晶度先增大后减少，这与由ＤＳＣ计算的结晶度结果相符㊂由表３还可以看出，２θ在２２．３ʎ和２８．９ʎ附近的衍射峰的半峰宽（ＦＷＨＭ）无明显区别，说明样品在不同温度下退火与退火前的样品相比，晶片厚度无明显区别，即ＰＧＡ为极易结晶型聚合物，退火前样品已形成晶格结构，由于分子链被瞬间冻结，存在一定晶格缺陷，而退火处理可以改善这一现象，消除分子之间的内应力㊂图３㊀ＰＧＡ的ＤＳＣ曲线图Ｆｉｇ．３㊀ＤＳＣｃｕｒｖｅｏｆＰＧＡ图４㊀ＰＧＡ的Ｘ射线衍射图Ｆｉｇ．４㊀ＸＲＤｃｕｒｖｅｏｆＰＧＡ表２㊀ＰＧＡ热性能参数Ｔａｂｌｅ２㊀ＴｈｅｒｍａｌｐａｒａｍｅｔｅｒｓｏｆＰＧＡ退火温度／ħｔ／ħΔＨ／（Ｊ／ｇ）ｔｍ／ħΔＨｍ／（Ｊ／ｇ）结晶度／％未退火－－２２３．７－４１．３２９．７１２０１２３．１－１．６２２４－６０．７４３．７１５０１５９．３－２．４２２３．６－６２．４４４．９１８０１７８．７－３．６２２３．７－６１．９４４．１１９０１９３．１－４．１２２４．１－６１．３４３．９㊀㊀注：ｔ：松弛峰值温度；ΔＨ：松弛焓；ｔｍ：熔融温度；ΔＨｍ：熔融焓㊂表３㊀ＰＧＡ的结晶参数及力学性能Ｔａｂｌｅ３㊀ＣｒｙｓｔａｌｌｉｚａｔｉｏｎｐａｒａｍｅｔｅｒｓａｎｄｍｅｃｈａｎｉｃａｌｐｒｏｐｅｒｔｉｅｓｏｆＰＧＡ退火温度／ħ２θ／（ʎ）ＦＷＨＭ／（ʎ）结晶度／％σｍａｘ／ＭＰａεｍａｘ／％Ｅ／ＭＰａ未退火２２．３０．４９２８．９０．７１８．２４９．６９．８３０１２０２２．３０．５０２８．９０．７１４０．３５８．０７．９３７１５０２２．３０．４９２８．９０．７２４１．５５８．５６．１５５１８０２２．３０．４９２８．９０．７０４０．６５７．６５．３５７１９０２２．３０．４９２８．９０．７１３９．８５６．９４．４６１７２㊀化㊀学㊀研㊀究２０２４年２．５㊀ＰＧＡ的力学性能ＰＧＡ退火前后的力学性能测试，以应变（ε）为横坐标，应力（σ）为纵坐标绘制应力应变曲线（见图５）㊂由应力应变曲线计算最大应力（σｍａｘ），最大应变（εｍａｘ），由应变在０．０５％０．２５％范围之间的曲线斜率计算拉伸模量（Ｅ），其结果见表３㊂图５㊀ＰＧＡ的应力应变曲线Ｆｉｇ．５㊀ＭｅｃｈａｎｉｃａｌｐｒｏｐｅｒｔｉｅｓｏｆＰＧＡ由表３可知，退火后样品的σｍａｘ比退火前明显提高，而随着退火温度的升高，σｍａｘ先增大后减少，这与ＸＲＤ测得的样品的结晶度结果是相符的，因为结晶度越大，其σｍａｘ越大㊂由表３还可以看出，随着退火温度的升高，εｍａｘ减小，Ｅ增大㊂这与ＰＧＡ的松弛焓计算结果是相符的，这是因为随着退火温度的升高，松弛焓增大，样品中非晶部分排列紧密有序，可变位移减小，储能模量增大㊂３㊀结论研究了ＰＧＡ的熔融冷却结晶性，实验结果表明，ＰＧＡ在熔融冷却过程中极易结晶，先形成针状晶核，径向增长，逐步生成球晶，且随着降温速率的增大，冷结晶温度降低，结晶时长变短，结晶焓变增大㊂还研究了不同退火温度对ＰＧＡ结晶性㊁热性能及力学性能的影响，实验结果表明，退火处理能明显改善样品的结晶性㊁热性能及力学性能，随着退火温度的升高，松弛温度升高，松弛焓变增大，应变减少，弹性模量增大，而结晶度和应力表现为先增大后减少㊂参考文献：［１］喻颖．聚乙交酯（ＰＧＡ）纤维的制备及其结构与性能研究［Ｄ］．上海：东华大学，２０２１．ＹＵＹ．Ｓｔｕｄｙｏｎｔｈｅｐｒｅｐａｒａｔｉｏｎ，ｓｔｒｕｃｔｕｒｅａｎｄｐｒｏｐｅｒｔｉｅｓｏｆｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄｆｉｂｅｒ［Ｄ］．Ｓｈａｎｇｈａｉ：ＤｏｎｇｈｕａＵｎｉｖｅｒｓｉｔｙ，２０２１．［２］ＢＵＤＡＫＫ，ＳＯＧＵＴＯ，ＳＥＺＥＲＵＡ．Ａｒｅｖｉｅｗｏｎｓｙｎｔｈｅｓｉｓａｎｄｂｉｏｍｅｄｉｃａｌａｐｐｌｉｃａｔｉｏｎｓｏｆｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄ［Ｊ］．ＪｏｕｒｎａｌｏｆＰｏｌｙｍｅｒＲｅｓｅａｒｃｈ，２０２０，２７（８）：１２５４⁃１２７６．［３］刘淑强，张瑶，杨雅茹，等．可用于医用纺织品的控释型载药ＰＬＡ／ＰＧＡ微球的制备及其释药性能［Ｊ］．上海纺织科技，２０２０，４８（１０）：１０⁃１２，１８．ＬＩＵＳＱ，ＺＨＡＮＧＹ，ＹＡＮＧＹＲ，ｅｔａｌ．Ｐｒｅｐａｒａｔｉｏｎｏｆｃｏｎｔｒｏｌｌｅｄ⁃ｒｅｌｅａｓｅＰＬＡ／ＰＧＡｍｉｃｒｏｓｐｈｅｒｅｓｆｏｒｍｅｄｉｃａｌｔｅｘｔｉｌｅｓａｎｄｔｈｅｉｒｄｒｕｇｒｅｌｅａｓｅｐｒｏｐｅｒｔｉｅｓ［Ｊ］．ＳｈａｎｇｈａｉＴｅｘｔｉｌｅＳｃｉｅｎｃｅ＆Ｔｅｃｈｎｏｌｏｇｙ，２０２０，４８（１０）：１０⁃１２，１８．［４］ＣＯＪＯＣＡＲＵＤＧ，ＨＯＮＤＫＥＳ，ＫＲÜＧＥＲＪＰ，ｅｔａｌ．Ｍｅｎｉｓｃｕｓ⁃ｓｈａｐｅｄｃｅｌｌ⁃ｆｒｅｅｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄｓｃａｆｆｏｌｄｆｏｒｍｅｎｉｓｃａｌｒｅｐａｉｒｉｎａｓｈｅｅｐｍｏｄｅｌ［Ｊ］．ＪｏｕｒｎａｌｏｆＢｉｏｍｅｄｉｃａｌＭａｔｅｒｉａｌｓＲｅｓｅａｒｃｈＰａｒｔＢ：ＡｐｐｌｉｅｄＢｉｏｍａｔｅｒｉａｌｓ，２０２０，１０８（３）：８０９⁃８１８．［５］ＢＲＥＵＥＲＣ．Ｔｉｓｓｕｅ⁃ｅｎｇｉｎｅｅｒｅｄｖａｓｃｕｌａｒｇｒａｆｔｓｆｏｒｃｏｎｇｅｎｉｔａｌｈｅａｒｔｓｕｒｇｅｒｙ［Ｊ］．ＴｈｅＦＡＳＥＢＪｏｕｒｎａｌ，２０２０，３４（Ｓ１）：１．［６］ＺＨＡＮＧＷ，ＷＥＩＺＣ，ＣＨＥＸ．Ｅｆｆｅｃｔｏｆｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄｍｅｓｈｆｏｒｐｒｅｖｅｎｔｉｏｎｏｆｐａｎｃｒｅａｔｉｃｆｉｓｔｕｌａａｆｔｅｒｐａｎｃｒｅａｔｅｃｔｏｍｙ：ａｓｙｓｔｅｍａｔｉｃｒｅｖｉｅｗａｎｄｍｅｔａ⁃ａｎａｌｙｓｉｓ［Ｊ］．Ｍｅｄｉｃｉｎｅ，２０２０，９９（３４）：ｅ２１４５６．［７］王功锦，黄杰．可降解材料聚乙醇酸垫片在全腔镜食管癌根治术中的应用［Ｊ］．临床外科杂志，２０２０，２８（６）：５４８⁃５５０．ＷＡＮＧＧＪ，ＨＵＡＮＧＪ．Ｔｈｅａｐｐｌｉｃａｔｉｏｎｏｆｂｉｏｄｅｇｒａｄａｂｌｅｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄｇａｓｋｅｔｉｎｔｈｏｒａｃｏｓｃｏｐｉｃｒａｄｉｃａｌｒｅｓｅｃｔｉｏｎｏｆｅｓｏｐｈａｇｅａｌｃａｎｃｅｒ［Ｊ］．ＪｏｕｒｎａｌｏｆＣｌｉｎｉｃａｌＳｕｒｇｅｒｙ，２０２０，２８（６）：５４８⁃５５０．［８］葛斌．用于外科手术止血夹的设计及ＰＰＤＯ／ＰＧＡ材料的化学环境模拟体外降解研究［Ｄ］．福州：福建师范大学，２０１８．ＧＥＢ．ＤｅｓｉｇｎｏｆａｈｅｍｏｓｔａｔｉｃｃｌｉｐｆｏｒｓｕｒｇｉｃａｌｏｐｅｒａｔｉｏｎａｎｄｓｔｕｄｙｏｎｉｎｖｉｔｒｏｄｅｇｒａｄａｔｉｏｎｉｎｔｈｅｃｈｅｍｉｓｔｒｙｓｉｍｕｌａｔｅｄｅｎｖｉｒｏｎｍｅｎｔｆｏｒＰＰＤＯ／ＰＧＡｍａｔｅｒｉａｌｓ［Ｄ］．Ｆｕｚｈｏｕ：ＦｕｊｉａｎＮｏｒｍａｌＵｎｉｖｅｒｓｉｔｙ，２０１８．［９］谭博雯，孙朝阳，计扬．聚乙醇酸的合成㊁改性与性能研究综述［Ｊ］．中国塑料，２０２１，３５（１０）：１３７⁃１４６．ＴＡＮＢＷ，ＳＵＮＣＹ，ＪＩＹ．Ａｒｅｖｉｅｗｉｎｓｙｎｔｈｅｓｉｓａｎｄｍｏｄｉｆｉｃａｔｉｏｎｏｆｐｏｌｙ（ｇｌｙｃｏｌｉｃａｃｉｄ）［Ｊ］．ＣｈｉｎａＰｌａｓｔｉｃｓ，２０２１，３５（１０）：１３７⁃１４６．［１０］朱爱臣，史建华，张宪科，等．基于正交试验法的聚乙醇酸精密注塑工艺研究［Ｊ］．化学研究，２０２１，３２（５）：４５２⁃４５７．ＺＨＵＡＣ，ＳＨＩＪＨ，ＺＨＡＮＧＸＫ，ｅｔａｌ．Ｓｔｕｄｙｏｎｔｈｅｐｒｅｃｉｓｉｏｎｉｎｊｅｃｔｉｏｎｍｏｌｄｉｎｇｐａｒａｍｅｔｅｒｓｏｆｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄｂｙｏｒｔｈｏｇｏｎａｌｅｘｐｅｒｉｍｅｎｔ［Ｊ］．ＣｈｅｍｉｃａｌＲｅｓｅａｒｃｈ，２０２１，第１期朱爱臣等：聚乙醇酸退火工艺研究７３㊀３２（５）：４５２⁃４５７．［１１］喻祖圣，郑翔睿，奚桢浩，等．聚乙交酯非等温结晶行为研究［Ｊ］．功能高分子学报，２０１９，３２（３）：３１６⁃３２１．ＹＵＺＳ，ＺＨＥＮＧＸＲ，ＸＩＺＨ，ｅｔａｌ．Ｎｏｎｉｓｏｔｈｅｒｍａｌｃｒｙｓｔａｌｌｉｚａｔｉｏｎｂｅｈａｖｉｏｒｏｆｐｏｌｙｇｌｙｃｏｌｉｄｅ［Ｊ］．ＪｏｕｒｎａｌｏｆＦｕｎｃｔｉｏｎａｌＰｏｌｙｍｅｒｓ，２０１９，３２（３）：３１６⁃３２１．［１２］李连明．不同分子量ＰＧＡ结晶性能及其对聚乳酸结晶性能影响的研究［Ｄ］．镇江：江苏科技大学，２０１９．ＬＩＬＭ．Ｓｔｕｄｙｏｎｃｒｙｓｔａｌｌｉｚａｔｉｏｎｏｆｐｏｌｙｇｌｙｃｏｌｉｄｅｗｉｔｈｄｉｆｆｅｒｅｎｔｍｏｌｅｃｕｌａｒｗｅｉｇｈｔａｎｄｉｔｓｅｆｆｅｃｔｏｎｃｒｙｓｔａｌｌｉｚａｔｉｏｎｏｆｐｏｌｙｌａｃｔｉｃａｃｉｄ［Ｄ］．Ｚｈｅｎｊｉａｎｇ：ＪｉａｎｇｓｕＵｎｉｖｅｒｓｉｔｙｏｆＳｃｉｅｎｃｅａｎｄＴｅｃｈｎｏｌｏｇｙ，２０１９．［１３］刘雨微，崔晓娜，韩铭，等．嵌段共聚物自组装薄膜中纳米结构的形成及应用［Ｊ］．化学研究，２０２１，３２（３）：１８９⁃２０４．ＬＩＵＹＷ，ＣＵＩＸＮ，ＨＡＮＭ，ｅｔａｌ．Ｆｏｒｍａｔｉｏｎａｎｄａｐｐｌｉｃａｔｉｏｎｓｏｆｎａｎｏｓｔｒｕｃｔｕｒｅｓｉｎｓｅｌｆ⁃ａｓｓｅｍｂｌｅｄｔｈｉｎｆｉｌｍｓｏｆｂｌｏｃｋｃｏｐｏｌｙｍｅｒｓ［Ｊ］．ＣｈｅｍｉｃａｌＲｅｓｅａｒｃｈ，２０２１，３２（３）：１８９⁃２０４．［１４］梁静谊，王大千，王槟鑫，等．玻纤增强聚甲醛复合材料性能研究［Ｊ］．化学研究，２０２１，３２（５）：４５８⁃４６５．ＬＩＡＮＧＪＹ，ＷＡＮＧＤＱ，ＷＡＮＧＢＸ，ｅｔａｌ．Ｐｒｏｐｅｒｔｉｅｓｏｆｇｌａｓｓｆｉｂｅｒｒｅｉｎｆｏｒｃｅｄｐｏｌｙｏｘｙｍｅｔｈｙｌｅｎｅｃｏｍｐｏｓｉｔｅｓ［Ｊ］．ＣｈｅｍｉｃａｌＲｅｓｅａｒｃｈ，２０２１，３２（５）：４５８⁃４６５．［１５］房鑫卿，肖敏，王拴紧，等．高分子量聚乙交酯的合成及表征［Ｊ］．高分子材料科学与工程，２０１２，２８（１）：１⁃４．ＦＡＮＧＸＱ，ＸＩＡＯＭ，ＷＡＮＧＳＪ，ｅｔａｌ．Ｓｙｎｔｈｅｓｉｓａｎｄｃｈａｒａｃｔｅｒｉｚａｔｉｏｎｏｆｈｉｇｈｍｏｌｅｃｕｌａｒｗｅｉｇｈｔｐｏｌｙｇｌｙｃｏｌｉｄｅ［Ｊ］．ＰｏｌｙｍｅｒＭａｔｅｒｉａｌｓＳｃｉｅｎｃｅ＆Ｅｎｇｉｎｅｅｒｉｎｇ，２０１２，２８（１）：１⁃４．［１６］薛锋，丁恩勇，程镕时．聚丙烯退火过程中高分子链的缠结［Ｃ］／／２００７年全国高分子学术论文报告会论文摘要集（上册）．成都：中国化学会，２００７：３２７．ＸＵＥＦ，ＤＩＮＧＥＹ，ＣＨＥＮＧＲＳ．Ｔｈｅｅｎｔａｎｇｌｅｍｅｎｔｏｆｐｏｌｙｐｒｏｐｙｌｅｎｅｃｈａｉｎｄｕｒｉｎｇａｎｎｅａｌｉｎｇ［Ｃ］／／ＳｕｍｍａｒｙＣｏｌｌｅｃｔｉｏｎｏｆ２００７ＮａｔｉｏｎａｌＰｏｌｙｍｅｒＡｃａｄｅｍｉｃＰａｐｅｒＲｅｐｏｒｔＣｏｎｆｅｒｅｎｃｅ（Ｖｏｌｕｍｅ１）．Ｃｈｅｎｇｄｕ：ＣｈｉｎｅｓｅＣｈｅｍｉｃａｌＳｏｃｉｅｔｙ，２００７：３２７［１７］ＳＨＡＷＥＳ，ＢＵＣＨＡＮＡＮＦ，ＨＡＲＫＩＮ⁃ＪＯＮＥＳＥ，ｅｔａｌ．Ａｓｔｕｄｙｏｎｔｈｅｒａｔｅｏｆｄｅｇｒａｄａｔｉｏｎｏｆｔｈｅｂｉｏａｂｓｏｒｂａｂｌｅｐｏｌｙｍｅｒｐｏｌｙｇｌｙｃｏｌｉｃａｃｉｄ（ＰＧＡ）［Ｊ］．ＪｏｕｒｎａｌｏｆＭａｔｅｒｉａｌｓＳｃｉｅｎｃｅ，２００６，４１（１５）：４８３２⁃４８３８．［责任编辑：任艳蓉］。

拉晶工艺英文介绍

拉晶工艺英文介绍The Introduction to the Process of Crystal Pulling.The process of crystal pulling, also known as the crystallization process, is a crucial step in the manufacturing of semiconductors and other electronic materials. This process involves the growth of single crystals from a molten state, typically in a controlled environment. The quality of the final crystal directly affects the performance and reliability of the resulting product.1. Raw Materials Preparation.The first step in the crystal pulling process is the preparation of the raw materials. Depending on the type of crystal desired, the raw materials can vary but typically include high-purity elements or compounds. These materials are carefully selected and mixed in precise ratios to ensure the desired chemical composition of the finalcrystal.2. Melting the Raw Materials.The next step is to melt the raw materials in a furnace or other heating device. The temperature and atmosphere within the furnace are carefully controlled to prevent impurities from entering the molten mixture. The molten material is then allowed to settle and homogenize, ensuring a uniform composition throughout.3. Seed Crystal Insertion.Once the molten material has reached the desired temperature and homogeneity, a small seed crystal is inserted into the molten mass. This seed crystal serves as the starting point for the growth of the larger crystal. The seed crystal must be carefully chosen to have the desired crystal structure and orientation.4. Crystal Growth.With the seed crystal inserted, the growth process begins. The molten material slowly solidifies around the seed, forming a single crystal. This process can take hours, days, or even weeks, depending on the size and complexityof the desired crystal. During this time, the temperature, pressure, and other parameters are carefully controlled to ensure the crystal grows with the desired properties.5. Cooling and Annealing.Once the crystal has reached its desired size, the growth process is stopped, and the crystal is allowed to cool slowly. This cooling process must be carefully controlled to prevent the formation of cracks or other defects. Additionally, the crystal may undergo an annealing process, which involves heating and cooling cycles torelieve internal stress and improve its mechanical properties.6. Post-Processing and Quality Control.After cooling, the crystal is removed from the furnaceand undergoes further processing, such as cutting, polishing, and shaping, to meet the final specifications. Quality control measures are employed throughout this process to ensure the crystal meets the required standards for purity, structure, and mechanical properties.7. Applications of Crystal Pulling.The crystal pulling process finds applications in various industries, including electronics, semiconductors, optics, and more. The resulting single crystals are used in devices such as transistors, lasers, and optical components due to their excellent optical, electrical, and thermal properties.In conclusion, the crystal pulling process is a crucial step in the production of high-quality single crystals. By carefully controlling the raw materials, melting process, seed crystal selection, growth conditions, cooling, and post-processing steps, manufacturers can produce crystals with the desired properties for a wide range of applications.。

超声退火法制备抗性淀粉及其理化性质分析

曾徐睿，刘良忠，齐婷，等. 超声退火法制备抗性淀粉及其理化性质分析[J]. 食品工业科技，2023，44（18）：292−299. doi:10.13386/j.issn1002-0306.2022120135ZENG Xurui, LIU Liangzhong, QI Ting, et al. Preparation of Resistant Starch by Ultrasonic Annealing Method and Its Physicochemical Properties[J]. Science and Technology of Food Industry, 2023, 44(18): 292−299. (in Chinese with English abstract).doi: 10.13386/j.issn1002-0306.2022120135· 工艺技术 ·超声退火法制备抗性淀粉及其理化性质分析曾徐睿1，刘良忠1, *，齐　婷1，朱　哲2（1.武汉轻工大学食品科学与工程学院，湖北武汉 430023；2.武汉隆丰园生物科技有限公司，湖北武汉 430040）摘　要：为优化超声退火法制备糯米抗性淀粉的工艺条件，本研究以抗性淀粉含量为指标，在单因素实验的基础上运用响应面分析法探究水分含量、超声时间、退火温度、退火时间对抗性淀粉含量的影响，并对所制备的抗性淀粉进行理化性质分析。

结果表明：超声退火法制备糯米抗性淀粉的最佳工艺条件为水分含量68%，超声时间23 min ，退火温度55 ℃，退火时间25 h 。

在该条件下制得的抗性淀粉含量为45.61%±0.95%。

与原淀粉相比，糯米抗性淀粉的透明度提高了38.19%，在90 ℃下的溶解度从6.03%增加到52.45%，而膨胀能力却从41.04 g/g 降至6.38 g/g 。

在扫描电镜的观察下，抗性淀粉的表面形貌发生明显变化，表面呈粗糙多孔结构。

不同退火条件对PEALD制备的Ga_(2)O_(3)薄膜特性的影响

annealing in different annealing atmospheres, and the increase in the proportion of N2 in annealing atmosphere is beneficial to
the recrystallization of Ga2 O3 . Furthermore, the effect of annealing time was further studied under N2 atmosphere, and the
中图分类号:O47;O484(2021)05-0838-07
Effects of Different Annealing Conditions on the Characteristics of
Ga2 O3 Thin Films Prepared by PEALD
stabilize after annealing for 30 min, and the surface grain density no longer increases. In addition, the average transmittance
of the sample in the range of 400 nm to 800 nm is almost 100% , and the light absorption edge is steep. Annealing in N2
Ga2 O3 作为氮化镓半导体的外延衬底制作垂直结构光电器件也非常具有吸引力 [14-15] 。尽管同质外延 Ga2 O3
薄膜获得了高导通电压和低漏电流的金属-半导体场效应晶体管 [16] , 但由于体材料衬底昂贵, 目前制备
Ga2 O3 薄膜多采用异质外延生长方法,制备方法主要有:金属有机物化学气相沉积( MOCVD) [17-18] 、分子束外

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Annealing effect on the crystal structure and exchange bias in Heusler Ni 45.5Mn 43.0In 11.5alloyribbonsL.González-Legarreta a ,W.O.Rosa b ,J.García a ,M.Ipatov c ,M.Nazmunnahar c ,L.Escoda d ,J.J.Suñol d ,V.M.Prida a ,R.L.Sommer b ,J.González c ,M.Leoni e ,B.Hernando a ,⇑aDepartment of Physics,University of Oviedo,Calvo Sotelo s/n,33007Oviedo,SpainbCentro Brasileiro de Pesquisas Físicas,Rua Dr.Xavier Sigaud,150Urca.,22290-180Rio de Janeiro,RJ,Brazil cDepartment of Materials Physics,Faculty of Chemistry,University of the Basque Country,20018San Sebastian,Spain dDepartment of Physics,Campus Montilivi s/n,University of Girona,17071Girona,Spain eDepartment of Material Engineering and Industrial Technologies,University of Trento,Via Mesiano 77,I-38123Trento,Italya r t i c l e i n f o Article history:Received 3July 2013Received in revised form 7August 2013Accepted 8August 2013Available online 21August 2013Keywords:Heusler alloysMartensitic transformation Magnetic properties Exchange biasa b s t r a c tA Heusler Ni 45.5Mn 43.0In 11.5alloy has been prepared by arc melting and produced in a ribbon shape by rapid solidiﬁcation using melt spinning technique.Structural properties have been investigated,at differ-ent temperatures,by using X-ray diffraction.Austenite is the stable phase at room temperature with a L21cubic crystal structure.Exchange bias effect was observed after ﬁeld cooling by means of hysteresis loop measurements.At 5K,hysteresis loop shifts along the axis of the applied magnetic ﬁeld and that shift magnitude decreases signiﬁcantly with increasing temperature.A piece of ribbon was annealed at 973K during 10min in order to investigate the inﬂuence of annealing on crystal structure and magnetic properties.After annealing,a martensitic phase with a monoclinic 10M structure at room temperature is observed.The onset of the martensitic phase transformation shifts to 365K,temperatures associated with both martensitic and reverse transitions do not change noticeably under an applied magnetic ﬁeld up to 30kOe,and a drastic decrease on magnetization is observed in comparison with the as-quenched ribbon meanwhile the exchange bias effect is enhanced.Ó2013Elsevier B.V.All rights reserved.1.IntroductionHeusler alloys are deﬁned as ternary intermetallic compounds with X 2YZ composition [1](where X and Y are transition metals and Z is a III,IV,or V group element)[2].Off-stoichiometric Ni 50Mn 50Àx In x ferromagnetic Heusler alloys,with x within a certain range of concentration values,exhibit several functional properties as the shape-memory effect [3],large inverse magnetocaloric and magnetic superelasticity effects [4],barocaloric effect [5]and mag-netoresistance effect and exchange bias (EB)behavior [6–10].Most of these effects could be ascribed to the existence of martensitic phase transition with a strong magneto-structural coupling,which are related to the total valence electron density (e/a ratio)as a key parameter [11].In this work we investigate the effects of annealing on structural and magnetic properties of the off-stoichiometric Ni 45.5Mn 43.5In 11.5Heusler alloy ribbons.Some preliminary results about EB behavior of these alloys are also reported.2.Sample preparation and experimental procedureA polycrystalline alloy with nominal composition Ni 45.5Mn 43.0In 11.5was pre-pared using conventional arc melting with high purity elements (>99.98at.%)under pure argon atmosphere.The ingot was melted several times to ensure chemical composition homogeneity.Then,the ingot was induction melted in a quartz tube in a melt spinning system and ejected,in argon environment,onto the polished sur-face of copper wheel rotating at an elevated linear speed of 48m/s.Ribbon ﬂakes were obtained with 7–12l m in thickness, 2.0mm in width and several mm (4–12)in length.In order to investigate the annealing effects on structural and mag-netic properties of Heusler alloy,a piece of sample was kept as reference (as-quenched ribbon)and some pieces of the ribbon were annealed for 10min at 973K in Ar atmosphere quartz tubes.For this,tantalum foil was used for wrapping each ribbon before introducing it in the quartz container for avoiding Si contamina-tion.Then,after removing from the furnace they were quenched in ice water.Microstructure,as well as composition determination,was performed through Scanning Electron Microscopy (SEM-JEOL6100)together with Energy Dispersive X-ray microanalysis (EDX-Inca Energy 200).The crystal structures of the annealed ribbon was analyzed by X-ray diffraction (XRD)at 150K,300K and 350K and as-quenched ribbon at 100K and 300K.XRD patterns were measured at Diamond Light Source,UK,in the range between 10°62h 6110°,by employing radiation of k =1.127Å.Thermomagnetic measurements,M (T ),with zero-ﬁeld cooling (ZFC),ﬁeld cool-ing (FC)and ﬁeld heating (FH)routines were performed in the temperature range of 50–400K applying in the ribbon plane different magnetic ﬁelds up to 30kOe by Vibrating Sample Magnetometry (VSM,Versalab,QD).The heating and cooling rate0925-8388/$-see front matter Ó2013Elsevier B.V.All rights reserved./10.1016/j.jallcom.2013.08.078Corresponding author.Tel.:+34985103307;fax:+34985103324.E-mail address:grande@uniovi.es (B.Hernando).was of 5K/min.Isothermal magnetization hysteresis loops,M (H ),were measured from À30to +30kOe,at a rate of 50Oe/s,in the temperature range of 5–300K by a physical property measurement system (PPMS-9,QD).diffraction patterns of the ribbon annealed at 973K.Monoclinic at 150K;(b)at 300K,and (c)at 350K.Insets show the patterns in the range 52°<2h <54°.(T)monoclinic 150K10M 0.429(6)0.584(5) 2.094(1)88.87300K 10M 0.437(3)0.592(4) 2.092(2)88.44350K10M0.436(3)0.589(2)2.094(1)88.39L.González-Legarreta et al./Journal of Alloys and Compounds 582(2014)588–593589becomes more important whereas the relative intensity of other peaks is reduced.The room temperature (RT)microstructure of the as-quenched and the annealed ribbon is shown in Fig.3a and b,respectively.As the martensitic transformation temperature is above RT for the annealed ribbon,the observed microstructure in this sample corresponds to the martensitic phase meanwhile austenite deter-mines the microstructure observed in the as-quenched sample.Thus,the difference observed in the microstructure of both sam-ples at RT is most likely related to the different crystal structure than to the annealing process.In these ﬁgures,it can be observed columnar grains located at the ribbon free surface which grown perpendicularly from small grains formed at the ribbon surface that was in contact with the wheel during the quenching process.For the as-quenched ribbon,see Fig.3a,grains with non-homoge-neous width size ranging between 0.5and 2l m size are observed in the columnar structure.However,for the annealed ribbon,see Fig.3b,mainly because the different crystalline structure,recrys-tallization,and a lower ordering degree than in the as-quenched sample,the grain size has increased in width between 1and 4l m,and the columnar structure is broaden.After an exhaustive study by EDX microanalysis,see Fig.3c and d,on different zones at the ribbons surface and transversal section,an averaged composition of Ni 45.8Mn 42.6In 11.6,rather near to the nominal one,for either the as-quenched and/or annealed samples was determined.The estimated error in determining the concen-tration of each element was of ±0.1%,and no Si contamination was appreciated either in the as-quenched ribbon,and/or in the annealed one probably due to the tantalum wrapping of this sam-ple inside the quartz tube during the thermal process.The micro-analysis conﬁrmed the chemical homogeneity of the alloy without inﬂuence of the grain shape on composition.The valence electron concentration per atom for this alloy results e/a =7.91.The temperature dependence of magnetization M (T )of both as-quenched and annealed ribbon measured under H =50Oe after ZFC and FC from 400K are shown in Fig.4a.M (T )curves after FC and FH at H =50Oe are collected separately in Fig.4b.As can be observed,the martensitic transformation and reverse transforma-tion temperatures of the Ni 45.5Mn 43.5In 11.5ribbon signiﬁcantly in-crease after annealing at 973K in comparison with the as-quenched one [12].The annealed ribbon presents the structural transition at M s (martensite start)of 365K ﬁnishing at M f (martensite ﬁnish)of 320K,being A s =330K and A f =380K the corresponding temperatures of the austenite start and ﬁnish,respectively.These characteristic ﬁrst order transition tempera-tures should be different when no external magnetic ﬁeld is ap-plied.Then,austenite is not the main phase observed in the XRD pattern at 350K and the detected texture probably could be only associated to the beginning of the transformation.Besides,a drastic decrease of the magnetization in the temper-ature range (50K <T <400K)results at an applied ﬁeld H =50Oe for this sample,see the magnetization axis on the right in both Fig.4a and b.This behavior could be ascribed to the annealing inﬂuence characterized by an increase in the grain size and crystal defects,along with the variation of internal stresses induced dur-ing the ribbon quenching process.The sharp increase of M at the Curie temperature,T C ,of austenite and the sharp drop of M after the martensitic transformation in the as-quenched ribbon are ab-sent in the treated sample.For both ribbons,the M (T )behavior can be explained by considering the variation of the magnetic ex-change parameters across the martensitic transformation from the paramagnetic cubic high-temperature austenite to the low-tem-perature ferromagnetic 10M monoclinic martensite [13].In the martensitic phase,the Mn-excess atoms occupying the In sublat-tice sites interact with the Mn atoms on the Mn sublattice sites antiferromagnetically.Then,an enhancement of the antiferromag-netic (AFM)interactions occurs at low temperatures,being also responsible for the drop in the magnetization during the martens-itic transformation [14,15].A bifurcation of FC and ZFC curve at low ﬁeld is revealed for both samples.An irreversibility between ZFC and FC curves also occurs at a spin freezing temperature,T f ,around 220K for the as-quenched ribbon [12]and at 310K for the annealed one.Besides,two peaks in the ZFC at 125K and 280K are observed for the annealed ribbon.This behavior is con-sistent with a magnetic state with superparamagnetic (SPM)or fer-romagnetic (FM)domains embedded in an antiferromagnetic matrix in the low temperature region that are collective frozen forming a super-spin glass sate at lower temperatures [16].It has been found for non-stoichiometric Heusler similar alloys such as in Ni 50Mn 25+y In 25Ày [14,16]and other relates as Ni 50Mn 25+y Sn 25Ày [17]and Ni 50Àx Co x Mn 40Sn 10[18,19].As the applied ﬁeld H is increased,see Fig.5,the martensitic phase transformation temperature keeps almost unchanged while the magnetization difference between austenite andmartensitic43.6In 11.5alloy:(a)as-quenched ribbon proﬁle,and (b)proﬁle of the ribbon annealed at 973K.EDX microanalysis:increases.Also,the splitting between ZFC and FC curves shift temperatures with increasing H.Similar tendency has observed in bulk polycrystalline Ni49.5Mn34.5In16alloyto a few ferromagnetic components embedded antiferromagnetic matrix[15].These features have been reported for Ni50.0Mn36.5In13.5Heusler alloywhere the structural transformation between martensitic and paramagnetic austenitic phasesfor In-based Heusler alloys.Therefore,the samplewould be in a paramagnetic state above and belowtemperature interval,see Fig.4a.off-stoichiometric Heusler Ni–Mn–In alloys it is wellexchange-bias of the hysteresis loops is often observed compositions where FM and AFM exchange interactionsThis effect can be interpreted in terms of magneticinto distinct FM and AFM regions in theFM clusters in an AFM matrix[16].The FM andchange interactions are modiﬁed after annealing by the decrease inthe interatomic distances,since the cell parameters of the as-quenched ribbon at100K are larger than the respective for the an-nealed sample at150K.The EB effect is attributed to the existence of a FM unidirectional anisotropy at the interface between differ-ent magnetic phases[20],and the general process to obtain this anisotropy isﬁeld-cooling the system from high temperature for reconﬁguring the FM spins at the mentioned interface.In order to detect the exchange bias effect in as-quenched and annealed Ni45.5Mn43.5In11.5ribbons,we measured M(H)fromÀ30to +30kOe at5K afterﬁeld cooling in+10kOe from375K,but for more clear visualization,hysteresis loops are only shown from À10to+10kOe,in Figs.6and7.We also measured the hysteresis loop at40K following the same protocol.This corresponds to cool-ing below T f,for both samples,starting from a temperature that is above the corresponding one to the respective martensitic phase transformation.As can be seen from Figs.6a and7a,at T=5K clear exchange bias effect is observed.Theﬁeld shift of exchange bias ﬁeld(H E),in M(H)loops are in the opposite sense to the cooling ﬁeld as expected.H E,is calculated using H E=À(H1+H2)/2,and coercivity(H C)is deﬁned as H C=À(H1ÀH2)/2,where H1and H2 are left and rightﬁelds at which the magnetization equals to zero.H E is about0.27kOe for the as-quenched ribbon while increases up to1.21kOe in the case of the annealed one.Nevertheless,a small increase is just observed at T=5K in the coerciveﬁeld value from 0.31kOe(as-quenched ribbon)to0.44kOe for the ribbon annealed at973K.At T=40K,see Figs.6b and7b,the EBﬁeld decreases to 0.13kOe for the annealed sample and to H E=0.006kOe for the as-quenched ribbon.This decreasing of exchange bias with increasing T is quite typical,and thus H E eventually vanishes fordependence(50–400K)of the magnetization for as-quenched and annealed Ni45Mn43.6In11.5ribbons at an applied magneticﬁeld of conditions from400K,and(b)FC from400K and FH conditions.Temperature dependence(50–400K)of the ZFC,FC and FH magnetization of11.5ribbon annealed at973K at several applied magneticﬁelds.Isothermal magnetization hysteresis loops of as-quenched Ni45Mnmeasured at(a)5K and(b)40K after FC of+10kOe magnetic(below the martensitic phase transformation)and(d)270K(abovemartensitic phase transformation)in ZFC regime.the as-quenched sample at around 50K,usually referred to as the blocking temperature,T B .However,the EB effect in the annealed ribbon exits in a wider temperature range up to 80K,probably due to the modiﬁcation of FM and AFM exchange interactions.An-other interesting feature,is the fact that H C also decreases with increasing T for the as-quenched sample to a value of 0.19kOe meanwhile an increase up to H C =1.04kOe occurs after annealing.This rather different behavior could be interpreted as follows.The decrease of H C with increasing T should be realized if the martens-itic phase was dominated by AFM order,and the decreasing of both H E and H C would be due to thermal instability of the nanoscopic FM clusters but not to the AFM order parameter [19].This situation re-sults similar to AFM/FM bilayers with Curie and Néel temperatures such that T C <T N [21].Moreover,if the thermal stability of the FM component is low relative to the AFM one,lower FM volume and/or magnetocrystalline anisotropy,the magnetization reversal in the FM is easy and does not imply signiﬁcant rearrangement of the AFM spin structure,and thus H C is decreasing at around T N [22].This behavior is consistent with a martensitic phase where nanoscopic FM clusters near the thermal stability would be embedded in a long-range order AFM matrix [19].In the case of the annealed sample,the fact that H C exhibits an increase with T has been previously observed in several non-stoichiometric Heus-ler Ni 50Mn 25+y X 25Ày (X =Sn,Sb,In)alloys [10,16,23,24].This coer-civity increasing is common in exchange-bias systems and occurs at FM/AFM interfaces existent between regions with Curie and Néel temperatures such that T N <T C ,and T B next to T N [20].Thus,as T N is approached from lower temperature,the decreasing anisot-ropy in the AFM region originates an increasingly signiﬁcant mod-iﬁcation of the AFM interfacial spin structure induced by themagnetization reversal in the FM region,leading to a signiﬁcant energy loss in the AFM,and as a consequence a H C enhancement [20].Additional information is provided by the isothermal M (H )loops from À30to +30kOe for two temperature values in each sample after zero-ﬁeld cooling from 375K.In the case of the as-quenched ribbon,the 270K M (H )curve (shown in Fig.6d)is above the martensitic transformation temperature and below the T C for austenite indicated by M (T ),see Fig.4,and a FM-like hysteresis loop is observed,with a coercive ﬁeld value of H C =49.6Oe.M (H )curve at 150K (Fig.6c)exhibits similar behavior for the martensite phase,but with a higher value of H C =67.0Oe indicating higher magnetocrystalline anisotropy than the respective to austenite phase and a decrease in magnetization due to the AFM exchange arising from the decrease in the Mn–Mn distance after annealing under martensitic transformation,resulting in a coexistence of FM and AFM phases in the martensitic state [16].More interesting and a quite different behavior is observed in Fig.7c and d,focusing on ZFC hysteresis loops for the annealed sample at 300K and 150K,both lower than the spin freezing T f ,into the martensitic phase.At 300K,a low nonlinearity is observed in M (H )superim-posed on a linear background.It should be remarked that the non-linear FM-like contribution (H C =7.2Oe)is restricted to low ﬁelds H <1kOe.As T is further decreased to 150K,a more signiﬁcantly nonlinearity (H C =91.6Oe)appears extending to higher ﬁeld H .Similar magnetic phenomena in the martensitic phase,has been observed in several related alloy system and interpreted as some form of superperamagnetic freezing of FM clusters [19,25]or nano-scopic spin cluster of unknown origin [16],that besides the exis-tence of FM and AFM exchange interactions in close competition,would explain magnetometry results including the observed ex-change bias effect and help to a more extensive understanding of the magnetic behavior in martensitic phase of Ni 45.5Mn 43.0In 11.5al-loy ribbons.4.ConclusionsOur experimental results reveal an important inﬂuence of short annealing on the crystalline structure and magnetic properties of Ni 45.5Mn 43.0In 11.5alloy ribbons.Martensitic transition is shifted to higher temperature but this does not increase with increasing ap-plied magnetic ﬁeld suggesting that the sample is in a paramag-netic state above and below the martensitic starting temperature.Magnetization is sensitively diminishing displaying a complex behavior in the low-temperature martensite,including a signiﬁ-cant exchange bias effect consistent with superparamagnetic freezing of ferromagnetic,FM,clusters embedded in an antiferro-magnetic,AFM,matrix.As-quenched and annealed samples exhibit the same temperature dependence of exchange-bias ﬁeld but opposite in the case of coercivity.Thus,coercivity decreasing with increasing temperature observed in the as-quenched ribbon should be realized if the martensitic phase was dominated by AFM order-ing and/or thermal stability of the FM component was low relative to the AFM one.However,the coercivity increasing with increasing temperature in the annealed ribbon can be attributed to the mag-netization reversal in the FM region as responsible to that enhance-ment.These features should be further investigated to realize the exchange bias effect in all its temperature range of existence.AcknowledgmentsThis work has been supported in part by Spanish MICINN under the Projects MAT2009-13108-C02-02and MAT2010-18914.Authors from UPV/EHU acknowledge the ﬁnancial support from the Department of Industry of the BasqueGovernmentIsothermal magnetization hysteresis loops of 973K annealed Ni ribbon measured at (a)5K and (b)40K after FC of +10kOe magnetic 150K and (d)300K (both temperatures below the martensitic transformation)in ZFC regime.and Compounds 582(2014)588–593(Programme SAIOTEK2011,Projects:S-PE11UN013,S-PE11UN087).Authors from UOviedo acknowledge theﬁnancial support from the Principado de Asturias(SV-PA-13-ECOEMP-47). 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