Effect of Size-Dependent Thermal Instability on Synthesis of Zn2SiO4-SiOx Core–Shell Nanotube Array

合集下载

《公共建筑节能(绿色建筑)工程施工质量验收规范》DBJ50-234-2016

本规范共分 22 章和 13 个附录，主要内容包括：总则，术语，基本规定，墙体节能工程，幕墙节能工程，门窗节能工程，屋面节能工程，地面节能工程，供暖节能工程，通风与空调设备节能工程，空调与供暖系统冷热源节能工程，空调与供暖系统管网节能工程，配电节能工程，照明节能工程，地源热泵系统节能工程，太阳能光热系统节能工程，太阳能光伏节能工程，监测与控制节能工程，建筑环境工程，资源综合利用工程，建筑节能（绿色建筑）工程现场实体检验，建筑节能（绿色建筑）工程质量验收。
（ 7 ）本规范第 16.2.10 条依据国家标准《太阳能供热采暖工程技术规范》 GB50495-2009 第 5.3.5 条的规定。
（8）本规范第 3.4.4 条为绿色建筑工程涉及的建筑环境与资源综合利用子分部工程验收方式的规定。
本规范由重庆市城乡建设委员会负责管理，由重庆市建设技术发展中心（重庆市建筑节能中心）、重庆市绿色建筑技术促进中心负责具体技术内容解释。在本规范的实施过程中，希望各单位注意收集资料，总结经验，并将需要修改、补充的意见和有关资料交重庆市建设技术发展中心（重庆市渝中区牛角沱上清寺路 69 号 7 楼，邮编：400015，电话：023-63601374，传真：023-63861277），以便今后修订时参考。
建设部备案号： J13144-2015
DB
重庆市工程建设标准 DBJ50-234-2016Leabharlann 公共建筑节能（绿色建筑）工程
施工质量验收规范
Code for acceptance of energy efficient public building(green building) construction
（3）本规范第 1.0.4、3.1.2、11.2.4、22.0.6、22.0.7 条内容分别依据国家标准《建筑节能工程施工质量验收规范》GB50411-2007 第 1.0.5、3.1.2 条、11.2.3、15.0.5、15.0.5 条等强制性条文要求。

元器件考题分析

第一章何谓电容器? 描述电容器的主要参数有哪些?由介质隔开的两块金属极板构成的电子元件，广泛用于储能和传递信息。

１. 标称容量和允许误差２. 额定电压３. 绝缘电阻４. 电容器的损耗５. 其他参数（频率特性，温度系数，稳定性，可靠性）有哪些类型的电容器? 各有何特点?无机介质电容器：介电常数高,介质损耗角正切小，电容温度系数范围宽，可靠性高，寿命长．有机介质电容器：电容量范围大，绝缘电阻(时间常数)大，工作电压高(范围宽)，温度系数互为偿，损耗角正切值小，适合自动化生产．热稳定性不如无机介质电容器，化学稳定性差, 易老化，具有不同程度的吸湿性电解电容器：比电容大，体积小、质量轻，有自愈特性双电层电容器(超级电容器)何谓电阻器? 可采用哪些主要的参数描述电阻器的性能优劣?具有吸收电能作用的电子元件，可使电路中各元件按需要分配电能,稳定和调节电路中的电流和电压。

１. 标称阻值和允许误差２. 额定功率3. 额定电压 4. 噪声有哪些主要类型的电阻器? 分别采用什么符号表示这些类型的电阻器?１．薄膜电阻器2．合金型电阻器3无帽结构电阻器4高频电阻器5小片式电阻器何谓电感器? 描述电感器的主要参数有哪些?凡能产生电感作用的元件统称为电感器１. 电感量及其误差２. 线圈的品质因数３．固有电容４．额定电流５．稳定性有哪些主要类型的电感器?固定电感器，平面电感器，天线线圈，震荡线圈，阻流圈简述矩形片式厚膜电阻器的电极结构特点及其设计依据．简述MLC的结构特点．试述片式铝电解电容器的结构特点和制作工艺过程，并简要说明各工序的作用．结构模型：电容器芯子/密封料/电极端子/封装树脂/矩形端子板工艺过程：腐蚀，形成，切断，柳接，卷绕，浸渍，装配，卷边，组合装配，特性测试，包装试述片式钽电解电容器的结构特点和制作工艺过程，并简要说明各工序的作用．结构模型：钽/氧化钽/二氧化锰/胶状石墨层/导电涂料层/导电性粘接剂层工艺过程：掺合，引线埋人加压成型，干燥烧结，阳极氧化，粘附二氧化碳，中间形成，石墨层，涂敷固化，连接，模塑封装，引线成型，老化，检验，包装比较片式铝电解电容器和钽电解电容器的结构差别．简述片式薄膜电容器的结构特点．带基薄膜—蒸镀电极一喷镀薄膜一外部电极有哪三种片式电感器？与目前使用的片式电感器相比，编织型片式电感器有何特点？简述框式电感器的制作过程．I．框式：结构最简单，但电感量较小II．螺旋式：能够获得较大的电感量，但要从中间抽头出来比较困难III．叉指式：电感量介于上述两者之间，但高频损耗大编织型片式电感器特点：单位体积电感量比目前已使用的片式电感器有所提高，但使用频率不高，Ｑ值低，尚未进入实用阶段，但这种工艺技术新颖．第二章:1.何谓微电子学?一种将电子产品微小型化的技术，他的核心产品是集成电路。

SIZE-DEPENDENT EFFECTS IN PROPERTIES OF

© 2009 Advanced Study Center Co. Ltd.Corresponding author: R.A. Andrievski, e-mail: ara@icp.ac.ruSIZE-DEPENDENT EFFECTS IN PROPERTIES OFNANOSTRUCTURED MATERIALSR.A. AndrievskiInstitute of Problems of Chemical Physics, Russian Academy of Sciences, Chernogolovka, Moscow Region,142432, Russia Received: February 17, 2009Abstract. Size-dependent effects in nanostructured (nanocrystalline, nanophase or nanocomposite)materials are of great importance both for fundamental considerations and modern technology.The effect of the nanoparticle/nanocrystallite size on surface energy, melting point, phase transfor-mations, and phase equilibriums is considered as applied to nanostructured materials. The role of size-dependent effects in phonon, electronic, superconducting, magnetic, and partly mechanical properties is also analyzed in detail. Special attention is paid to the contribution of other factors such as the grain boundary segregations, interface structure, residual stresses and pores, non-uniform distribution of grain sizes, and so on. The little explored and unresolved problems are pointed and discussed.1. INTRODUCTIONSize-dependent effects (SDE, i.e. the characteris-tic size influence of grains, particles, phase inclu-sions, pores, etc ., on the properties of materials and substances) have been studied in physics,chemistry, and materials science for a long time. It is quite enough to list the following well-known equations of Laplace, Thomson (Kelvin), Gibbs-Ostwald, Tolman, J. Thomson, Hall-Petch,Nabarro-Herring, and Coble, which connect the capillary pressure (P ), saturated vapor pressure (p ), saturated solubility (C ), surface energy of flat surface (σ0), conductivity (λ), hardness (H ) and creep rate (′ε) correspondingly with the pore/in-clusion radius (r ), thickness film (h ) and grains/crys-tallites size (L ):P r =20σ/, (1)p r p rRT =B =B =002exp /,σΩ (2)C r C rRT =B =B=002exp /,σΩ (3)σσξ0012r r =B =B =+//,(4)λλh f =B =B =+0510./. (5)H L H AL =B =+−005., (6)~/,ε1L m (7)where p o is the equilibrium pressure of saturated vapor on flat surface, C o is the equilibrium solubil-ity of subject with flat surface, Ω is an atomic (mo-lar) volume, R is the gas constant, T is tempera-ture, ξ is the Tolman constant, λo is conductivity of grain-coarse material (thick film), f = h /l o , l o is the mean free pass of carriers (f < 1), H o is the friction stress in the absence of grain boundaries (GB) and A is a constant. The factors m =2 and 3 correspond to the Nabarro-Herring diffusion creep and to theCoble diffusion creep correspondingly.The development of modern advancednanostructured materials (NM) manifests some new problems such as the identification adaptabil-ity of these equations (1-7) and other those in na-108R.A. Andrievskinometer interval; the role of quantum effects as well as an influence of other factors in the NM prop-erties or their stability and reproducibility. All these things are very important and actual to consider the SDE features taking in mind the update high rate in this field and not only fundamental consid-erations, but the strategy optimal development for nanotechnology. These questions have been ana-lyzed early by the present author [1-5] and other investigators (see, e.g., [6-19]). It seems to be nec-essary to tell apart the SDE in the isolated nanosubjects (clusters, nanoparticles, nanowires,and nanotubes) and consolidated ones (nanobulks,nanocomposites, and nanofilms/nanocoatings).Analyses of SDE in consolidated NM can be found in reviews [1-6], but the most publications were devoted to nanoparticles and clusters. Naturally,the SDE topic as a whole seems not to be cov-ered, because there are many new results scat-tered throughout the literature.Taking into account the comprehensive books [9,10,13,15,16], this review is mainly devoted to the SDE recent analysis in consolidated NM (papers mainly from 2001-2002), although some results for isolated nanosubjects will be also considered with the exception of catalytic and biological properties.Moreover, the comprehensive problem on the SDE in mechanical properties will be also discussed in limited scale as applied only to brittle high-melting point compounds (see Section 4 below).It is important to point that there are at least five principal features of size effects in consolidated NM [1-5]:- the crystalline size reduction to the nanometer scale results in a significant increase of the role of the interface defects such as GB, triple junc-tions (TJ), and elastically distorted layers;- interfacial properties on a nanometer scale can be different from those of conventional grain-coarse materials;- the reduced crystallite size can be overlapped by characteristic physical lengths such as the mean free path of carriers or the Frank-Read loop size for dislocations and so on;- size effects in NM can have quantum nature if the crystallite size is commensurable with the De Broyle wavelength l B =D /(2m *E )1/2 where D is the Planck constant, m * is the electron effective mass, E is the electron energy. Based on known values of m * and E, it is possible to predict that the quantum size effects can be exhibited in met-als only when the crystallite size is lower than ~ 1nm. For semiconductors (especially with narrow gap-zone, e.g., InSb) and semimetals (e.g., Bi),the l B value is significantly higher and can be of about 100 nm;- unlike conventional grain-coarse materials, in NM there are many factors for masking SDE such as residual stresses, pores, TJ and the presence of other defects, progressive accumulation of inter-face segregations, non-equilibrium phase ap-pearance, and so on.All these peculiarities may reveal in the pres-ence of certain specific points in the size depen-dencies and in the non-monotonous change of the properties determined by decreasing of the grain size. As a rule, such non-monotonous change is especially characteristic for structure-sensitive properties. For example, such specific features, i.e.twist points, and non-monotonous change of prop-erties have been revealed in the known hardness and magnetic coercivity studies (see, e.g., [1,2,4]).Such changes can be connected with different fac-tors such as other mechanisms of SDE, progres-sive accumulation of segregations on interfaces,and so on.2. THERMODYNAMIC PROPERTIES 2.1. Surface energyThe problem concerning surface energy is impor-tant both for fundamental sense and for many materials science and engineering applications such as the prognosis of phase diagrams, the es-timations of the fracture, milling, dilution, wetting,nucleation, coagulation, recrystallization charac-teristics, and so on. In this connection, the reveal-ing of surface energy change in nanometer inter-val is of special interest.Tolman’s equation (4), i.e. the effect of the drop radius on surface energy, has been widely dis-cussed (see, e.g., [7,8,11,14,20-23]). Omitting the details of these calculations, which are based on different physical and chemical approaches, let us to consider these results as a whole. Being ap-plied to isolated nanocrystals under calculations [11,14], the amendment g (r ) to the conventional σ0value fixed on the capillary pressure (1) is describedby the following expressiong r r r rr rB=B =B =−+−−211422182ln /ln /For many nanocrystals this amendment is very small and comes into particular prominence (about 10%) only at r < 5 nm.The radius effect on surface energy can be also described by equations:109Size-dependent effects in properties of nanostructured materialsσσr Kr at r r at r r =B =<>456,,, (9)where K is a constant, r o is critical radius (of 2-10nm) [21]. There are other theoretical calculations that have revealed about the similar values of criti-cal radius lower of which surface energy is de-creased (see, e.g., [20,22]). The experimental re-sults in this field are very limited, ambiguous, and often mutually contradictory. From one side, elec-trochemical measurements have revealed the σr values about 0.2 J/m 2 and 0.7 J/m 2 for Ag nanoparticles of r ~ 50 nm and r ~ 150 nm accord-ingly to [24]. These values seem to be not so pre-cise but are reasonable (it is appropriate to com-pare them with the value of σo = 1.14 J/m 2 and with the GB energy of ~0.27 J/m 2). From the other side,the experimental study of size-dependent evapo-ration of free-spherical Ag and PbS nanoparticles using relationship (2) results in values of 7.2 J/m 2and 2.45 J/m 2 respectively, these values are sig-nificantly higher as compared to that of the flat bulks [25]. May be, this is an effect of the evaporation conditions (see Section 2.2) or there is a system-atic error in the experimental results [25].As applied to nanocrystalline solids the analy-sis of size- or curvature-dependent interface, free energy results to the following relationshipσσr t r =B =B≈−01/, (10)where t is an interfacial layer (the semi-width of GB in consolidated nanometals) [7,8]. It is easy to see that at the reasonable values of 2t ~ 1nm and r = 5 nm the σ decreasing is of 10% that is in close agreement with estimations under Eq. (8).In Debye approximation, the Tolman constant ξin Eq. (4) can be expressed asξα=′−151.,h =B(11)where h ’ is the height of atomic monolayer, α is the ratio of root-mean-square amplitude of thermal vi-brations of atoms on surface and in volume [23]. If α > 1 and ξ > 0, then σr < σo and in the opposite case α < 1 and ξ < 0 then σr > σo . As it will be seen from here on, these two cases are important as applied to the behavior analysis of nanoparticles in matrices.It is known that distinct feature of nanostructure in NM is that, with the grain size confinement, the part of TJ is increased and, accordingly, the part of GB is decreased (see, e.g., [4]). The computer modeling with account of GB and TJ contributions reveals that Ni grain size reducing is accompaniedby decrease of the total excess enthalpy (i.e. the total GB energy) [26]; this result quantitatively agrees with experimental data for nano Se [27].The calculated interface GB energy in Cu-Ni bilayered nanofilm changes from 0.9 J/m 2 to about 0.7 J/m 2 when bilayer thickness decreases from ~1 nm to ~0.4 nm [28].From the foregoing data it has been evident that in spite of different theoretical approaches in cal-culations, they are evidence of the σ decrease in the confinement of isolated nanosubjects and nanostructures in NM. It seems to be reasonable to introduce the amendments into grain/crystalline boundary energy values at L lower than near 10nm. However, it is also evident that experimental studies in this field are necessary for further de-tailed continuation.The comprehensive description of the SDE role in the other surface phenomena such as wetting,nucleation, adsorption, and capillarity can be found in book [16]. It is worth to note that the observation by high-resolution transmission electron micro-scope (HRTEM) methods is in common practice (see, e.g., [19,29,30]).2.2. Melting pointIt is well known for a long time that small particles and thin films are characterized by the lower melt-ing points (T m ) as compared with their counterparts in bulk form as a result of the atom thermal vibra-tion amplitude increase in the surface layers. There are many relationships contacting T m and r for free-standing (supporting) metal (Au, Ag, Cu, Sn, etc .)clusters/nanoparticles as well as h for thin films (see, e.g., [4,10,13,15,16,19]). Generally, all pro-posed equations have the form of T m ~ 1/r (h ) type.Recent theoretical and experimental works tend to obtain more precise specifications: the size,shape, and stress effects on the T m nanograins on a substrate [31]; the role of fractal structure [32];the comparison of SDE in the case of spherical nanoparticles, nanofilms, and nanowires [9,33]; the effect of transition to amorphous state [34], kinetic study of the Cu film melting-dispergation [35], etc .Thus the consideration of size-dependent cohesive energy results in the T m variation for spherical nanoparticle, nanowire, and nanofilm of a material in the same characteristic size as 3:2:1 that is con-firmed by experimental values for In [33]. The mo-lecular dynamics simulation of the melting behav-ior for “model ” nanocrystalline Ag has exhibited two characteristic regions on the grain size decrease [34]. The first is above about 4 nm where the T m110R.A. Andrievskivalues decrease with grain size decreasing. The second one (r < 4 nm) is a size-independent re-gion where the T m values almost keep a constant.The dominant factor in this situation is supposed to be the nanocrystal T m shifts from grain to GB.The amorphous phase (r < 1.3 nm), as indicated from the radial distribution function and common neighbor analysis, is different from the GB by the sharp enhancement of local fivefold symmetry and is characterized by a much lower solid-to-liquid transformation temperature than that of Ag nanocrystal and GB. Melting and non-melting of solid surfaces including partly behavior of nanosystems are discussed in review [12].When the nanosolids are entirely coated by a high-melting point substance or embedded in such matrix, the nanosolid T m values can not only de-crease but increase too; this interesting phenom-ena is named by superheating or overheating. TheSME detailed analysis at melting and superheat-Fig. 1. Melting point of In nanoparticles embedded in Al matrix for two methods of preparation as a function of particle diameter: ball milling (n ) and spinning (l ). HRTEM images illustrated the different In/Al interface structures, replotted from [19].ing of bulks and nanosolids can be found in review [19]. The crucial role of interface in superheating is clearly demonstrated by the measurement re-sults of the ball-milled and the melt-spun Al-In samples T m values as a function of particle size of In (Fig. 1) [19].Fig. 1 clearly demonstrates that samples pre-pared by the two methods undergo completely dif-ferent melting behaviors which were investigated using differential scanning calorimetry (DSC), in-situ HRTEM, and in-situ X-ray diffraction (XRD). In the case of ball-milling procedure, there are irregu-lar In particles and the incoherent random inter-faces were formed between particles and the ma-trix, then the nanoparticles exhibit the SDE T m de-pression, as described previously. In the case of melt-spun samples, the In particles were found to be distributed both in the Al GB and within the Al grains. The particles within Al grains are truncated by octahedral shapes bounded by {111}/(100) fac-111Size-dependent effects in properties of nanostructured materials Compound Usual bulk phase Observed(metal, alloy)at RT and P =106 Pa nanocrystalline phaseZrO 2 Monoclinic Tetragonal BaTiO 3 Cubic Tetragonal PbTiO 3 Cubic Tetragonal Al 2O 3 α-Al 2O 3 γ-Al 2O 3 TiO 2 RutileAnataseY 2O 3Cubic (α-Y 2O 3) Monoclinic (γ-Y 2O 3)CdSe, CdSWurtzite Rock salt Co hcp * (α-Co) fcc * (β-Co) T hcp (α-Ti) bcc * (β-Ti) Fe-NiMartensiteAustenite (fcc)*hcp is hexagonal-closed-packed structure; fcc is face-centered-cubic structure; bcc is body-centered-cubic structure.Table 1. Structure of unusual phases observed in some NM [37].ets and can be considered as epitaxial coherent interfaces revealing the T m increase with decreas-ing of r . It is interesting to know whether other prop-erties, not only melting point, are also changing under the ball-milled and melt-spun nanoparicles in matrices. The superheating has been also ob-served in the systems of Pb-Al and Ag-Ni both in the Al and Ni matrixes as well as partly in layered Pb/Al films [19].On phenomenological level, superheating is explained by change of root-mean-square ampli-tude of thermal vibrations of atoms on surface and in volume (Eq. (11) [23]) and by modification of in-terfacial energy [19] or the coherent interface pa-rameter between inclusions and matrix [36].2.3. Phase transformationsApart of melting point SDE make an impact on other phase transformations in NM. Table 1 summarizes some data in this field [37].Observed solid-state transformations are con-sidered to be connected with an increased effec-tive internal pressure due to high surface/interfa-cial curvature or with the whole and surface en-ergy difference between allotropic phases [37]. The tetragonal-to-monoclinic (T → M) phase transfor-mation in the zirconia-yttria system is the most ex-plored subject. Fig. 2 shows experimental data on the T → M transformation temperature (T tr ) for dif-ferent YSZ powders and sintered pellets with par-ticles/grains in the nanometer interval [38].In Fig. 2 each line demarcated the tetragonal phase stability region (defined as “T ”) and that toleft of the line (defined as “T+M ”) where the T-phase starts to transform into the M-phase which was fixed in both the DSC and the dilatometer studies. It is evident from these results that the T tr value is de-creased with a reduction of crystallite/grain size.The curve extrapolation provide the opportunity to predict the particle critical size values at room tem-perature of the T-phase stability: 15, 30, 51, and 71 nm for 0YSZ, 0.5YSZ, 1.0YSZ, and 1.5YSZ powders, respectively. In the case of the 0.5YSZ,1.0YSZ, and 1.5YSZ sintered pellets, the grain criti-cal size values are 70, 100, and 155 nm, respec-tively. The observed difference in crystalline/grain size (about 2 times) can be attributed to the strain energy term and difference in the surface and in-terfacial energies.In the thermodynamic approach frame, the fol-lowing relationships between the T tr value and criti-cal crystalline/grain size (D c , L c ) have been pro-posed116///for powders ,D H T T c b tr b =−∆∆σ=B =B =B (12)1166////for pellets ,L H T T U cbtrbd=−+∆∆Σ∆∆Σ=B =B =B (13)where ∆H b and T b are the enthalpy and transfor-mation temperature for bulk (coarse-grained) sol-ids, ∆σ and ∆Σ are the differences in the surface/interfacial energies for T/M phases in powders and pellets, ∆U d is the strain energy involved in the transformation (this additional strain-energy term estimates only for consolidated pellets) [37,38].112R.A. AndrievskiFig. 2. Inverse critical grain/crystalline size versus transformation temperature (T → M) for different yttria-doped zirconia (a) powders and (b) pellets: 0YSZ – pure ZrO2, 0.5YSZ – 0.5 mol.%Y2O3, 1.0YSZ – 1.0mol.%Y2O3, 1.5YSZ – 1.5 mol.%Y2O3, replotted from [38].113Size-dependent effects in properties of nanostructured materials Fig. 3. The phase stability diagram for Nb/Zr multilayers as a function of the Nb volume fraction (f NB ) and the inverse of bilayer thickness (λ-1): # 1 is the stability region of hcp Zr/bcc Nb; # 2 is the stability region of bcc Zr/bcc Nb; # 3 is the stability region of hcp Zr/hcp Nb. The circles indicate experimental results and the line boundary calculated from the thermodynamic model, replotted from [41].The possibility to observe also the high-tem-perature phase at low temperature is the situation when the elementary crystal (germ) size of this phase is lower than crystallite size. Such situation tends to the transformation blocking and the high-temperature phase is fixing. This example is typi-cal for some martensite transformations in the Fe-Ni and Ti-Ni-Co micro- and nanocrystalline alloys [39]. The martensite volume part (M ) depends on the high-temperature initial phase crystallite size under the equationM M K Lm =−−005., (14)where M o and K m are some constants. The rela-tionship (14) coincides with that of Hall-Petch one (6). This agreement could be connected with the analogous nature of elastic strain fields at the crack and germ propagation.The availability of non-equilibrium phases in thin film is observed for a long time. Under film thick-ness lower than some critical value (h c ), non-com-mon phases can be fixed; that is connected with the Gibbs free energy excess effect due to the sur-face effect. One of the first relationships describ-ing this effect was proposed more than 50 years agoh F F c =−−σσ1221=B =B/, (15)where indexes 1 and 2 correspond to equilibrium and non-equilibrium phases, respectively, σi and F i are the total surface energy and the Gibbs free energy of these phases [40]. The h c rough estima-tion by Eq. (15) resulted in the quite reasonable values (h c ~ 10 nm).Basically, the idea [40] based on a classical ther-modynamic approach with some modifications is used in many modern studies especially with the development of nanostructured thin film multilayers such as Ti/Al(Nb, Zr) and TiN/AlN(NbN, ZrN), Fe/Cr, etc ., which can exhibit metastable structures in one or both layers. Fig. 3 shows the phase stability diagram for Nb/Zr multilayers with varying volume fractions and bilayer (λNb + λZr ) thickness [41].The phase stability region boundaries were cal-culated using a classical thermodynamic model taking into account the structural energy assess-114R.A. AndrievskiFig. 4. The Bi-Pb phase diagram. The dashed lines are the equilibrium bulk phase diagram and the solid lines are experimental results for nanoparticles of (a) 10 nm radii and (b) 5 nm radii, replotted from [42].115Size-dependent effects in properties of nanostructured materials Fig. 5. Pseudobinary TiN-TiB 2 phase diagram for equilibrium state (the solid lines) and for nanostructured films (L ~ 10 nm), replotted from [44].ment in the interfacial energy value. The equilib-rium phases (hcp Zr and bcc Nb) were observed only in the region #1; at that, the effect of the Nb volume fraction is not monotonous. The presence of the non-equilibrium bcc Zr and hcp Nb phases in sputtered films (regions #2 and #3, respectively)was confirmed experimentally by XRD and elec-tron diffraction.2.4. Phase equilibriumsBy now there are several examples of theoretical calculations and experimental studies of the phase diagram SDE in nanointerval: ZrO 2-Y 2O 3 [37], Pb-Bi [42], In-Sn [30], Au-Sn [30], carbon phase dia-gram [43], Al-Pb [18], Bi-Cd [18], TiB 2-TiN [44], TiB 2-B 4C [45], hydrogen systems [46,47], etc . Fig. 2 data have been used for the T/T+M line design in the ZrO 2-Y 2O 3 diagram as a function of particle size [37].The most detailed study was undertaken in the Pb-Bi diagram observing in-situ the melting behav-ior of isolated alloyed nanoparticles by hot stage TEM. Fig. 4 compares this diagram for particles of 10 nm radii and 5 nm radii as well as for bulks [42].As shown in Fig. 4, T m is a size-depend decreas-ing from the bulk value. The progressive narrow-ing of two-phase fields and an increase of solubil-ity are also observed. The theoretical calculations developed from thermodynamic first principles were in agreement with experimental results.The significant increase of solubility for nanosized subjects was observed in many cases including the hydrogen-metal (intermetallic) sys-tems [46,47], Pb-Fe and Pb-Sn systems [48,49],and systems-based high-melting point compounds [44,45]. Fig. 5 shows the TiN- TiB 2 phase diagram both in equilibrium state and film one (the grain size about of 10 nm) [44].The significant increase in solubility was fixed by the experimental XRD test [50] and the eutectic temperature decrease that was calculated in the regular solution approximation using the phase equilibrium at the eutectic temperature. The con-116R.A. Andrievskitribution made by the surface energy excess for solid state was written in per mole terms as ∆F S = 6V σgb /L where V is the molar volume and σgb is the interfacial (grain boundary) energy (see also [51]). The same approach was used for the SDE analysis in the case of the TiB 2-B 4C phase diagram [45]. Unfortunately,the σgb value information is very limited. It was supposed in the TiN-TiB 2 phase diagram calculations (Fig.5) that σgb ~ 3 J/m 2; in the case of the TiB 2-B 4C phase diagram the more reasonable σgb interval from 1J/m 2 to 3 J/m 2 was used[45].The solubility features in nanosystems are theoretically discussed in papers [52,53]. In many cases,the great role of the grain boundary segregations is pointed (see, e.g., [48]). However, as a whole, the question on the oversaturated solution formation nature remains insufficiently explored as applied to mechanical alloying, ion plating, severe plastic deformation, electrodeposition, and other preparation methods of NM.Another interesting feature in the SDE on phase equilibriums is that the spinodal decomposition and the triple point coordinates. It was theoretically shown that the grain segregations (thickness of about 1nm) increase the solid solution decomposition degree when the grain size does not exceed 10 nm [54].The spinodal decomposition regularities of supersaturated TiN-AlN system with the thermal stable nanostructure formation are very interesting and important for the development of new advanced NM. Ab initio calculations revealed the contribution role of strain and surface energy on the energetic balance for decomposition process [55].The calculation of SDE on the carbon diagram results in the following data for the triple point coordi-nates defined regions for solid nanodiamond, liquid that and amorphous nanocarbon [43] Particle diameter, (nm) 1 1.2 1.5 2 3 4 6 Singlecrystal Pressure (GPa) 15.2 16.5 16.115.615.214.814.513.5 T, K2210316035503820409041904300 4470These results are in qualitative agreement with some experimental/theoretical information [56]. The movement of triple point, defined equilibrium between solid state, liquid one, and vapor one, to the tem-perature decrease and pressure increase with the particle decrease was also described in book [9].In the most cited works dedicated to the SDE in phase equilibriums, the conditional thermodynamic approaches were used and the interface structure influence was not analyzed. In mean time, as it was shown at the Fig. 1, the impact of external interfaces on T m can be very great and opposite, so that interface contributions to the thermodynamic functions become significant [18,19]. The analysis of two-phase equilibrium in alloy nanoparticles has revealed the possibility of the eutectic point degeneration into intervals of composition where the alloy melts discontinuously and of the phase diagram topology dra-matic changes [57]. These observations tended to conclusion that the concept of equilibrium in nanostructures yet remains to be understood and needs in further detailed study [18,51].3. PHYSICAL PROPERTIES3.1. Phonon and electronic properties. SuperconductivityMeasurements of electronic heat capacity (γe ) of Cu-Nb (Pb) nanocomposites and studies of their phonon spectra revealed the decrease of the electronic state density at the Fermi level N (E F ) and the increase of low-frequency oscillations as compared with their coarse-grained (cg ) counterparts [58,59]. The γe value for Nb inclusions (L ~20 nm) in Cu 90Nb 10 nanocomposite was of 3.0 mJ/mol K 2 that is about three times lower than that in conventional Nb. The Debye temperature (ΘD ) value is also reduced with the inclusion decrease. These changes were connected with the weakening of interatomic interactions at the GB and assisted by the superconducting transition temperature (T C ) decrease.The size dependence of T C , the residual resistance ratio (RRR), the upper critical magnetic field (H C2)and the Ginzburg-Landau coherence length (ξGL ) for the nanocrystalline Nb films are shown in Table 2[60,61].It is clear from Table 2 that while the T C and RRR values decrease with size monotonically, the H C2 and ξGL values increase with size but there is non-monotonous change. The first change is supposed to be connected with the N (E F ) decrease rather than by phonon softening, i.e. the effect of the electron-phonon。

HPGe探测器对圆柱型体γ源效率刻度的经验公式

HPGe探测器对圆柱型体γ源效率刻度的经验公式聂鹏;倪邦发;田伟之【摘要】介绍了HPGe γ探测器的圆柱型体源效率刻度的经验公式.利用单能量γ射线放射性源~(141)Ce(145.4keV)、~(51)Cr(320.8 keV)、~(198)Au(411 keV)、~(65)Zn(1115.5 keV)的水介质体源对经验公式进行了检验.结果表明,用经验公式计算的相对效率与实验结果的相对偏差在5%以内,证明了该公式的实用性.【期刊名称】《核技术》【年(卷),期】2010(033)001【总页数】2页(P79-80)【关键词】HPGeγ探测器;体源;效率刻度;经验公式【作者】聂鹏;倪邦发;田伟之【作者单位】中国原子能科学研究院核物理研究所,北京,102413;中国原子能科学研究院核物理研究所,北京,102413;中国原子能科学研究院核物理研究所,北京,102413【正文语种】中文【中图分类】O571.5在体源γ射线发射率的HPGe γ谱学测定中，通常采用点源近似法或相同几何和近似基体的标准体源比较测定的方法。

点源近似法是一种粗略的处理方法，不适合于精确的计算。

用相同体积和近似基体的标准体源进行探测效率的相对测定方法，虽然精确，但要制作满足上述要求的标准体源，既耗费劳力又浪费物力。

本文通过对体源放射性探测的物理过程、体源的几何特征以及体源与探测器的相对位置的分析，推导出以给定几何点源效率归一的、圆柱型体源中级联符合效应可忽略的、γ射线全能峰效率计算的经验公式以放射性水溶液体源141Ce (145.4 keV)、51Cr (320.8 keV)、198Au (411 keV)、65Zn (1115.5 keV)γ射线全能峰相对效率实测值与计算值的比较，证明了该公式的实用性和简便性。

设圆柱型体源的半径为R，高度为H，体源底部到HPGe探测器表面的距离为S，探测器的半径为R’，探测器的有效作用深度为S0[1]。

Nano-glasses--A_new_type_of_glasses

Nano-glasses: A New Type of Glasses?Herbert GleiterInstitut fuer NanotechnologieForschungszentrum KarlsruhePostfach 3640 , 76021 Karlsruhe/Germanye-mail: herbert.gleiter@int.fzk.deFax: +49-7247-82-6368AbstractToday, glasses are produced mostly by rapidly cooling the melt. In the area of metallic glasses, research attention has been focused in the past on finding chemical compositions that result in metallic glasses even if the cooling rates are as low as 10 degrees per second or less. These studies led to the development of the so called bulk metallic glasses.However, so far no systematic attempt has apparently been made to modify the atomic structure and,thus, the properties of glasses by changing their free volume in a controlled manner, although it is well known that the structure and the properties of glasses not only depend on their chemical composition but also on their free volume. In this paper it is proposed to generate a new type of glasses, called nano-glasses. Nano-glasses consist of nanometer-sized glassy regions connected by “glass/glass interfaces” These "interfaces" between adjacent glassy regions may have an enhanced free volume up to 15% (as compared to typically 1-2% of a melt grown glass). Due to this enhanced free volume, nano-glasses differ structurally and propertywise from glasses(with the same chemical composition) produced by cooling the melt. Nano-glasses may be synthesized either by assembling nanometer-sized clusters of atoms with a glassy structure or by introducing (by plastic deformation) such a high density of shear bands into a metallic glass that the spacing between adjacent bands is a few nanometers.IntroductionSoon after their discovery (1), the attractive properties of metallic glasses (produced by rapidly cooling the melt) became apparent. For example, in the tightly packed glassy structure, the displacement of atoms (e. g. to accommodate plastic flow) is obstructed. The metallic glass absorbs less energy than a crystalline body upon stress-induced deformation through damping and returns more by rebounding elastically to its initial shape. Due to the absence of crystal defects, glassy materials the following remarkable properties:- Strength (twice that of stainless steel but lighter);- Hardness (desirable e.g. for surface coatings);- Toughness (more fracture resistant than ceramics); and- Elasticity (high yield strength).The absence of grain boundaries renders many glassy materials resistant to corrosion and wear, as well as possessing soft magnetic properties, specifically in alloys of glass formers (B, Si, P) and ferrous magnetic transition metals (Fe, Co, Ni). High electrical resistivity leads to low eddy current losses. Easy magnetization and demagnetization allows lower losses in applications, operation at high temperatures with minimal flux density reduction, and annealing. On the other hand, in the years after the discovery of metallic glasses, high cooling rates were required (typically 105 to 106 °Cs-1) to prevent the melt from crystallizing. These high cooling rates limited initially (1) the maximum thickness of metallic glasses to about 0.1 mm or less. In order to remove this drawback, one of the main research thrusts in the area of metallic glasses focussed in the last four decades on finding metallic glasses with chemical compositions that require cooling rates as low as possible. In 1969 Chen and Turnbull (2) found ternary Pd-M-Si glasses (with M=Ag, Au or Cu) requiring cooling rates less than 103 °Cs-1. These cooling rates allowed casting thickness of about 0.5 mm. In 1974, Chen obtained a critical casting thickness of 1 mm in Pd-T-P (T=Ni, Co, Fe) (2) and in the early 1980s, Turnbull's group produced glassy ingots of Pd40Ni40P20 with a diameter of 5 mm using surface etching followed by heating and cooling cycles – the first bulk metallic glass (3). In 1988, Inoue discovered that La-Al-TM (TM=Ni, Cu) is highly glass-forming. It allowed him to cast glassy La55Al25Ni20 up to 5 mm thick and, in 1991, glassy La55Al25Ni10Cu10 specimens that were up to 9 mm thick (4), (5). At about the same time Johnson and Peker at Caltech developed a pentary alloy based on Zr-Ti-Cu-Ni-Be (Zr41.2Ti13.8Cu12.5Ni10.0Be22.5). With critical casting thickness of up to 10 cm (6) An overview of the critical casting thickness and the date of the discovery of these alloys is shown in Fig. 1Fig. 1 The critical casting thickness versus the year in which alloys were discovered. Over the last 40 years, the critical casting thickness has increased by more than three order of magnitude (6).Recently, Loeffler et al (7) described a method based on the centrifugal processing of the melt in order to find systematically the chemical compositions of multi component alloys with the lowest melting point. Alloys with these compositions are likely to require low cooling rates for glass formation.In all of these studies, relatively little attention was paid to the free volume1 of glasses, although it is well known that the properties of glasses depend strongly on their free volume. For example, if metallic glasses are plastically deformed at temperatures far below their glass transition temperature, shear bands may develop and slip is localized to these shear bands (9) – (11). Spaepen (8) first put forward the idea that shear band formation is to be attributed to a multiplication of free volume at an incipient band, which reduces flow stress locally so that shear is localized to this mechanically weakened region. Several authors (9), (10) concluded from the overall density reduction of plastically deformed metallic glasses that the enhanced free volume in a shear band may be as high as 50%. Moreover, it was found that this excess free volume in the shear bands was dissipated during a subsequent relaxational annealing treatment (10). According to more recent studies (11), the largest enhancement of the free volume in the shear bands exists during the deformation process and is called a ‘mechanically dilated flow state’. Once the deformation process is over, irreversible relaxation processes occur which reduce the free volume in the shear bands. The idea of an enhanced free volume and thus an enhanced atomic mobility in the shear bands is consistent with the observation that nanometer-sized crystals grow in the shear bands at ambient temperature, whereas outside of these bands no crystallization occurs (12), (13). The reduction of the enhanced free volume in the shear bands during subsequent annealing agrees with the following observation. Annealing of glasses is found to affect their properties e.g. the diffusivity, their optical properties etc. These property variations are believed to result from atomic rearrangements in the glasses which are often associated with the reduction of the free volume.It is the purpose of this paper to point out that the way to a new kind of glasses - called nano-glasses – may be opened by varying the free volume of glasses in a controlled way during the preparation procedure. In fact, this approach may result in glasses with new atomic structures and properties.Basic IdeasThe basic idea to generate glasses with enhanced free volume is explained schematically in Figs. 2a to 2f (14).Fig. 2 a-d. Generation of nano-glasses (schematic) by assembling nanometer-sized glassy solids.1 The free volume of a glass is the excess volume of a glassy solid relative to a crystalline solid withthe same chemical composition.Fig. 2a: Two-dimensional glass made up by hard spheres. This piece of glass is cut into four segments (A, B, C, D) along the lines indicated.Fig. 2b: Separation of the four segments.Fig. 2c: The positions of the four segments have been changed as indicated by the arrows in Fig. 2b, so that the segment A is now in the initial position (Fig. 2a) of segment B, B in the initial position of C etc. After this cyclic positional exchange, the four segments are re-assembled so that no atomic overlap occurs. This procedure results in the structure shown in Fig. 2c. Fig. 2d: All sites larger than one atomic volume in the interfaces between the four segments (Fig. 2c) are filled by inserting new atomsLet us start by considering a piece of a two-dimensional glass formed by hard spheres (Fig. 2a). This piece of glass is cut into four segments labelled A, B, C and D as is indicated in Fig. 2a. Subsequently, these segments are separated and positionally moved so that A gets into the position were the segment B was formerly, B to C etc. (Fig. 2b). After this positional exchange, the four segments are reassembled so that some of the atoms at the "surfaces" of the four segments (e. g. atoms 1 and 2 in Fig. 2c) get into contact without any atomic overlap. Finally, the empty space between the four segments is filled by inserting new atoms into all sites that are larger than one atomic volume (Fig. 2d). As may be seen from Fig. 2 d, the resulting glass consists of the four segments (A, B. C and D) connected by the horizontal and vertical contact regions between them. These contact regions are called glass/glass interfaces. In these interfaces the atomic density is lower than in the interior of the four segments. A simple estimate shows that the free volume in such glass/glass interfaces may be as high as 15%. In other words, it may be about one order of magnitude larger that the free volume trapped in glasses produced by rapidly cooling a melt with the same chemical composition. If the hard spheres (black circles in Fig. 2d) are assumed to represent atoms, the glass shown in Fig. 2d is a glass that exhibits a microstructure on a nanometer scale consisting of nanometer-sized glassy regions (the segments A, B, C, D) connected by interfaces with an enhanced free volume (relative to the free volume in the interior of the four segment). In this paper, a glass with this kind of microstructure will be called a nano-glass. Clearly, the microstructure and, thus, the enhanced free volume of a nano-glass maybe varied by varying the size of the assembled glassy regions (Figs. 2c and 2d). Infact, if the size of these regions is increased, a nano-glass becomes structurally more and more comparable to a bulk glass produced by cooling the melt.The procedure to generate nano-glasses by assembling nanometer-sized glassy regions (Figs. 2a-d) is not the only conceivable approach. The incorporation of a high density of (nanometer-spaced) shear bands (by plastically deforming a bulk metallic glass far below its glass transition temperature) seems to be an additional alternative. The generation of nano-glasses by this approach is indicated schematically in Figs. 2e and 2f.Fig. 2e Fig. 2fFig. 2e: Two-dimensional model of a glass (schematic). The atomic arrangement shown is the same as the one displayed in Fig. 2 a.Fig. 2f: Two-dimensional model of a plastically deformed glass (cf. Fig. 2a). The plastic deformation has introduced two shear bands (one in the horizontal and one in the vertical direction). The resulting atomic structure is the one of an nano-glass consisting of nanometer-sized glassy regions (A, B,C, and D) connected by glass/glass interfaces (shear bands).Fig. 2e displays the same two-dimensional models of a glass that was already shown in Fig. 2a. This glass is plastically deformed (Fig. 2f) by introducing two shear bands. These shear bands section the glass into four nanometer-sized regions (labelled A, B, C and D in Fig. 2f). The four regions are displaced relative to one another in directions parallel to the corresponding shear bands. Due to the aperiodic atomic structure of the glass, these displacements result in similar atomic structures in the contact regions between A, B, C and D as were described in Figs. 2c and 2d when four glassy segments (labelled A, B, C, and D in Figs. 2c and 2d) were assembled. In order to avoid atomic overlap in the shear bands, the atomic density in the bands has to be reduced relative to the atomic density in the interior of the glassy regions (A, B, C and D in Fig 2f). In other words, the introduction of nanometer-spaced shear bands results in a microstructure consisting of nanometer-sized glassy regionsconnected by glass/glass interfaces. The ideas just explained agree with the experimentally observed large excess free volume (of up to 50 %) in the shear bands of plastically deformed glasses (8) – (13) (cf. Introduction of this paper).Let us add a few general comments concerning the physics behind the idea of enhancing the free volume of a glass so that a nano-glass results. The enhanced free volume of an nano-glass is defined as the excess free volume of a nano-glass relative to the free volume of a fully relaxed bulk glass with the same chemical composition. In other words, the bulk glass is used as the reference state for the nano-glass. In the most simple case of a one component glassy model system, this fully relaxed bulk glass is the densest disordered packing of spheres of equal size. The use of a fully relaxed glassy state as a reference has been successfully applied in the past e.g. for defining the glass temperature (15) (16).The significance of the enhanced free volume of a nano-glass may become physically more transparent by considering it in the light of our present understanding of the physical significance of the free volume of bulk glasses. Molecular dynamics simulations of the atomic rearrangement processes in glasses and in undercooled melts indicate that glasses are heterogeneous as far as the dynamic properties of their structure are concerned. Due to the heterogeneous spacial distribution of the free volume, glasses exhibit localized regions of high and low atomic ability. In the regions of high mbility, atomic rearrangements processes occur (often in a co-operative mode) at a higher rate than outside of these regions. Such atomic rearrangements are important for transport processes in glasses such as diffusion, relaxations etc. (17). This implies that the free volume of a glass is a simplification of the complex reality of the atomic structure and, thus, in general, does not allow to understand many properties in detail. In principle, the same seems to apply to nano-glasses. However, the basic idea of nano-glasses is not affected by this complexity. Production of Nano-glassesOne way to produce nano-glasses (18) (19), is by means of inert gas condensation (Fig. 3). Nanometer-sized glassy spheres are generated by evaporating (or sputtering) the material in an inert gas atmosphere. These spheres accumulate at the surface of a cold finger. After scraping them off, they are compacted into a pellet shaped nano-glass specimen. Compaction is performed in a high pressure consolidation device. This procedure is basically identical to the one used to generate nanocrystalline materials. So far, nano-glasses have been synthesized by inert gas condensation from a variety of alloys: Au-Si, Au-La, Fe-Si, La-Si, Pd-Si, Ni-Ti, Ni-Zr, Ti-P (20) – (22).Fig. 3: Production of nano-glasses by consolidation of nm-sized glassy spheres produced by inert gas condensation (18) (19).The production of nano-glasses by heavily deforming bulk metallic glasses (e.g. by rolling, ball milling or extrusion) has not yet been reported.Structural Studies--Small Angle X-ray ScatteringThe nanometer sized microstructure of nano-glasses is expected to show up in SAXS experiments, as was observed (20), (22). Fig. 4 displays the SAXS intensity (20),(21) for a model system with glassy regions of a size of 3.4 nm and a standard size deviation of 1.53 which corresponds to the value found in inert-gas condensed powders. In these experiments, the radial density profile of the glassy regions was described by a power law function. The deviation between the model and the experiment at small scattering angles (Fig. 4) is likely to result from the fact that the log-normal size distribution assumed, underestimates the number of large glassy regions. Ti-P alloys have a negative heat of mixing (~65 kJ mol-1). Hence, concentration fluctuations increase the free energy of the system significantly. In other words, the (Ti70P30) nano-glass (Fig. 4) is likely to be uniform in chemical composition and the SAXS curve shown in Fig. 4 reflects the modulation of theatomic density in this nano-glassFig. 4: Experimental small-angle X-ray scattering intensity of a Ti70P30(•••)nano-glass and the calculated scattering intensity for a model nano-glass with delocalized interfaces (----) as described in this paragraph (22).--Mössbauer-spectroscopyFig. 5 shows the Moessbauer spectra and the corresponding quadrupole splitting (QS) distributions of a melt spun Fe72 Si18 Fe10 glass and a nano-glass with the same chemical composition produced by consolidating 3.6 nm-sized glassy spheres.Fig. 5 Mössbauer spectra and corresponding QS distribution (p(QS)) of a melt-spun Pd72Fe10Si18 metallic glass (upper part of the figure) and a nano-glass with the same chemical composition generated by consolidating 3.6 nm-sized glassy spheres (18).Fig. 6: Relative spectral fraction of the interfacial component versus the inverse droplet size (18).Compared to the melt-spun glass, the nano-glass exhibits two peaks in the QS distribution. The first one coincides with the one of the melt-spun glass. The second one (at about 0.9 mm s-1) was observed only in the nano-glass, and was, thus, assumed to originate from the "interfaces" between the glassy regions. On the other hand, an oxide and/or a crystalline phase may result in a similar peak. Chemical analysis revealed, however, less than 2% oxygen in the nano-glass and crystallization of the nano-glass removed the second peak in the QS distribution (Fig.5).If the interfacial thickness depends little on the size of the glassy regions, the spectral fraction of the second component of the QS distribution in the Mössbauer spectrum (Fig. 5) should be inversely proportional to the size of the droplets consolidated. This was indeed observed experimentally (Fig. 6). From the slope of Fig. 6, the thickness of the interfaces was deduced to be 0.4 nm, corresponding to about two atomic spacings (Fig. 2d). Hence, the structural model (Fig. 2d) of a nano-glass (consisting on nm-sized glassy regions connected by glass/glass interfaces with a reduced density) seems to agree with the Mössbauer results. In fact, this observation also agrees with another result of the Mössbauer measurements (18). In addition to the higher QS value, the interfacial component also exhibited an enhanced isomer shift. This result is similar to the observations reported for other nanometer-sized systems (24) and indicates a reduced s-electron density. The reduced s-electron density may be interpreted in terms of a lower average atomic density and/or electron localization effects at the interfaces.The density profile of the glass/glass interfaces of nano-glasses seem to depend on the chemical composition and/or the time temperature history of specimen preparation. In Fig.7, the SAXS diffraction curves for Si75Au25 and Ti70P30 nano-Fig. 7: Comparison of the small angle scattering curves of Si-Au and Ti-P nano-glasses with the scattering curve expected for atomically sharp glass/glass interfaces (20)-(22).glasses are compared. If the glass/glass interfaces would be atomically sharp, a intensity profile with a slope of 4 (Porod scattering) should be seen. This was not the case. According to Fig. 7, the slope of the Ti/Pd-nano-glass is about 3. A slope of 3 suggests that the glass/glass interfaces are not atomically sharp but delocalized on a scale of several atomic layers. The slope of the SAXS curve for the Si-Au glass was about two (Fig. 7). A SAXS-slope of about two is seen for density fluctuations at the critical point (Ornstein-Zernicke Scattering), (20)-(22)2. In summary, the SAXS results seem to disagree with the localized glass/glass interface structure deduced from the Mössbauer data (Fig. 6).Multi-component Nano-glassesSo far the discussion was limited to nano-glasses consisting of one kind of atoms only. If a nano-glass contains atoms of different chemical elements, multi-component nano-glass may be formed. For the sake of brevity, we shall limit ourselves to two types of multi-component nano-glasses.Fig. 8 displays schematically the atomic structure of a nano-glass consisting of glassy regions made up by chemically different atoms (symbolized by open and closed circles). In this case, two kinds of glass/glass interfaces result. Interfaces between chemically identical glassy regions and interfaces between regions of different chemical compositions.______________________________________________2 In both interpretations void scattering (sample porosity) is assumed to be insignificant.Fig.8: Model (schematic) of a nano-glass that consists of regions with different chemical compositions. The chemically different atoms are represented by open and close circles.. Fig. 9: Schematic cross section through a two-dimensional nano-glass. The atoms are represented by circles. The material consists of nanometer-sized, glassy regions. In the central part of these regions (full circles), the interatomic spacings are similar to the ones of a bulk glass. In the interfacial areas (broken lines, open circles), a broad spectrum of interatomic spacings exists.Fig. 9 (19) shows schematically the structure of a nano-glass made up of glassy regions the surfaces of which differ chemically (atoms represented by open circles) from the centers of these regions (full circles).Research PerspectivesThe studies performed so far on nano-glasses seem to reveal their existence and evidence an atomic structure in which the density varies on a namnoeter scale (Figs.2 , 4 and 7).However, so far little is known on the atomic structure of nano-glasses as a function of chemical composition, the preparation procedure and their time-temperature history. MD-simulations of the atomic structure and the dynamics of nano-glasses appear to be useful to deduce a structural model. In fact, the results of such simulations could be related directly to observations by electron microscopy as well as measurements by SAXS, WAXS, C p and of the elastic constants of nano-glasses. Moreover, studies of the diffusivity, plasticity, the ferromagnetic as well as the optical properties and the thermal stability of nano-glasses seem to be of interest in order to improve our basic understanding of these materials and to evaluate their technological potential as well.AcknowledgementsThe author gratefully acknowledges the support and numerous helpful discussions with numerous colleagues, in particular with R. Birringer, H. Hahn, U. Herr, H. J. Hoefler, U. Gonser, J. Jing, H.Roesner,J. Weissmueller and G.Wilde.Post-deadline addition:Nano-glasses and crystalline materials seem to represent two different classes of materials concerning the coupling between the atomic structure and the defect structure. In crystalline materials, the cores of lattice defects are localized to a few interatomic spacings. For example, the core diameter of a dislocation or the width of a grain boundary are in the order of a few lattice constants. Outside of the defect cores, the atomic arrangement in a crystal is the arrangement given by the crystal structure (slightly distorted by the elastic strain fields of the defects). The physical reason for the localization is the fact that the crystalline state represents the state of lowest free energy at low temperatures. Hence, all regions in which the atomic arrangements deviate significantly from the crystal structure (e.g. in the cores of defects) are minimized in size. In nano-glasses the situation is differentin the following sense. Just like in a liquid, the local arrangement of the atoms in a glass can vary as a function of time at constant free energy and at constant average density of the glass. For example, a glass may deform by viscous flow (i.e. by local variations of the atomic arrangements) near the glass transition temperature. Hence, if a localized region of reduced atomic density - such as a glass/glass interface - is introduced into a glass, the glass can reduce its free energy by delocalizing the locally reduced density of the interface. In other words, the density reduction initially localized in the interface is gradually “smeared” out into the regions of the glass adjacent to the interface, so that a glass of slightly reduced average density (and modified atomic structure) results in the vicinity of the glass/glass interface. (The same process is known to happen, if a local density reduction is introduced into a liquid, e.g. by intense ultrasonic radiation). This implies, by introducing defects into glasses, one can modify the glass structure (and hence the properties of a glass) if the defects are allowed (e.g. by annealing) to delocalize. A comparable process does not exist in crystals. The low energy of the crystalline state prevents this process. In other words, there seems to be a fundamental difference between glassy and crystalline materials. In glasses (e.g. in nano-glasses) defects (e.g. glass/glass interfaces) may be used modify the atomic structure (and, hence, the properties) of the glass by controlling the density, the type and the degree of delocalization of the defects introduced.References(1) Klement, W., Willens R. H., Duwez, P., Nature 187 (1960) 869(2) Chen, H. S., Turnbull, D., Acta Metall. 22 (1974) 1505(3) Kui, H. W., Greer, Al. L., Turnbull D., Appl. Phys. Lett. 45 (1984) 615(4) Inoue, A., Zhang, I., Masumoto, T., Mater. Trans. JIM 30 (1989) 965; 31 (1990)177(5) Kato, A., Zhang T., Kim, S. G., Masumoto, T., Mater. Trans. JIM 31 (1991) 177and 425(6) Peker, A., and Johnson, W. L., Appl. Phys. Lett. 63 (1993) 2342(7) Loeffler, J. F., Bossuyt, S., Peker, A., W. L. Johnson, Phil. Mag. 83 (1003)2797(8) Spaepen, F., Acta Metall. 25 (1977) 407(9) Cahn, R. W. Pratten, N. A., Scott, M. G., Sinning, H. R., Leonhardson, E., Mat.Res. Soc. Symp. 28 (1984) 241(10) Krishnand. K. D., Cahn, R. W., Scripta Metall. 9 (1975) 1259(11) Kim, J. J., Choi, Y., Suresh, S. Argon, A. S., Science 295 (2002) 654(12) Hebert, R. J., Roesner, H., Perepezko, J., Wilde, G., Scripta Mat. inpress(13) Wilde, G., Boucharat, N., Hebert, R. J., Roesner, H., Valiev, R. Z., Mat. Sci.and Eng., submitted(14) Gleiter, H., Europhysics News 20 (1989) 130(15) Tool, A. Q., Eichlin, C. G., J. Amer. Cer. Soc. 14 (1981) 27616) Mitsch, C., Goerler, G. P., Wilde, G., Willnecker, R., J. Non-Cryst. Solids 270‚ (2000) 172(17) Raetzke, K., Hueppe, P. W., Faupel, F., Phys. Rev. Lett. 68 (1992) 2347(18) Jing, J., Kramer, A., Birringer, R., Gleiter, H., Gonser, U., J. Non-Cryst. Solids ´ 113 (1989) 167(19) Gleiter, H., J. Appl. Cryst. 24 (1991) 79(20) Weissmueller, J., Schubert, P., Franz, H., Birringer, R., Gleiter, H., Proc. VIINatl. Conf. on the Physics of Non-Crystall. Solids Cambridge, England August4-9, 1991(21) Weissmüller, J., Birringer, R., Gleiter, H., Key Engineering Materials 77-78,(1993) 161(22) Weissmueller, J., Birringer, R., Gleiter, H., Microcomposites and NanophaseMaterials publ. in: Proc. of the TMS Annual Meeting New Orleans 1991, p. 291 (23) Averback, R. S., Hahn, H., Hoefler, H. J., Logas, J. C., Appl. Phys. Lett. 57(1990) 1745(24) Herr, U., Jing, J., Birringer, R., Gonser, U., Gleiter, H., Appl. Phys. Lett. 50(1987) 472。

微通道翅片换热参数模拟

微通道翅片换热参数模拟英文回答：Microtubular Finned Heat Exchanger Parameters Simulation.Microtubular finned heat exchangers (MTFHEs) are compact and high-performance heat exchangers that have gained increasing attention in various industrial applications. The thermal performance of MTFHEs is influenced by several parameters, including the geometry of the microtubes and fins, the working fluids, and the operating conditions. Computational fluid dynamics (CFD) simulations can provide valuable insights into the complex flow and heat transfer phenomena occurring within MTFHEs and help optimize their design and operation.CFD Modeling of MTFHEs.CFD simulations of MTFHEs typically involve solving thegoverning equations of fluid flow and heat transfer, such as the Navier-Stokes equations and the energy equation. The geometry of the MTFHE is discretized into a computational mesh, and the governing equations are solved at each mesh node. The resulting solutions provide detailed information about the flow field, temperature distribution, and heat transfer rates within the MTFHE.Parameters Considered in CFD Simulations.Microtubes:Diameter.Length.Pitch.Material.Fins:Fin height.Fin thickness.Fin spacing.Fin material.Working Fluids:Properties (density, viscosity, thermal conductivity, specific heat capacity)。

PCB_Thermal_Considerations (1)

Gear2-Materials is an encyclopedia in your pocket for ThermoPhysical properties of over 1,300 materials. These properties include: Density, Specific Heat, and Thermal Conductivity.
Test ambient
Measurement of θJA Devices soldered to special thermal resistance test boards 8-9 mil (200-225µm) standoff from board Placed in box of known volume (1cu ft if you’re American!) Temperature rise measured
PLCC θJC test setup
θJC data
Power dissipation has an effect on thermal resistance Must be considered when calculating cooling requirements
Other factors affecting θJC
Test setups
Test device on board
Air flow test setups
θJC Tests
Test device held against an infinite heatsink This comprises a massive, water-cooled copper block, kept at 20°C In this way, θCA (case-ambient) is very close to zero, so any measurement is purely θJC (junction-case)

原子层沉积硫化钼

Atomic Layer Deposited MoS 2as a Carbon and Binder FreeLi-ion BatteryDipK Nandi,UttamK Sen1,DevikaChoudhury,Sagar Mitra,Shaibal K Sarkar *Department of Energy Science and Engineering,IIT Bombay,Mumbai 400076,IndiaA R T I C L E I N F O Article history:Received 3June 2014Received in revised form 22August 2014Accepted 3September 2014Available online 22September 2014Keywords:Atomic layer deposition Molybdenum sul ﬁdeQuartz crystal microbalanceFourier transform infrared spectroscopy Lithium-ion batteriesA B S T R A C TMolybdenum sul ﬁde is deposited by atomic layer deposition (ALD)using molybdenum hexacarbonyl and sul ﬁde.Film growth is studied using in-situ quartz crystal microbalance,ex-situ X-ray ﬂectivity and ellipsometry.Deposition chemistry is further investigated with in-situ Fourier transform spectroscopy.Self-limiting nature of the reaction is observed,typical of ALD.Saturated growth of 2.5Åper cycle at 170 C is obtained.As-deposited ﬁlms are found amorphous in nature.As-grown lms are tested as lithium-ion battery anode under half cell con ﬁguration.Electrochemical charge-discharge measurements demonstrate a stable cyclic performance with good capacity retention.Discharge capacity of 851mAh g À1is obtained after 50cycles which corresponds to 77%of capacity retention of the initial capacity.ã2014Elsevier Ltd.All rights reserved.1.IntroductionSimilar to other transition metals,Mo has multiple oxidation states leading to different sul ﬁde like MoS 2,Mo 2S 3and MoS 3[1–4].Among the above,MoS 2was studied most due to its layered crystallographic assembly [5,6].A schematic of layered MoS 2structure is shown in Fig.1.Thin ﬁlm MoS 2was demonstrated in a variety of electronic applications [7,8].Lately,MoS 2also drew considerable attention for its application as anode material in Li-ion battery [9–16].Theoretically MoS 2can accommodate 4Li ions through conversion reaction leading to a gravimetric capacity of 670mAh g À1.In literature,a reversible capacity in the range of 850to 900mAh g À1has been reported using bare MoS 2[9–16].As lithiation induced microstructural and chemical changes take place in MoS 2,the governing chemistry of MoS 2opens up the possibility of employing amorphous materials directly as an electrode for lithium ion battery application.It is expected that chemical and microstructural changes in amorphous materials during lithiation can be energetically preferred in comparison to its crystalline phase.Among several existing thin ﬁlm deposition technologies,atomic layer deposition (ALD)has an advantage of conformal deposition on complex substrates [17].Surface bound self-limiting growth with precise thickness control ensures the applicability of ALD in variety of applications [18–21].Though the application of ALD in electrochemical storage is overwhelming,but till now the use is mostly limited to the protection layer to enhance the electrochemical performance of the electrode [22–27].In recent literature,ALD was used to produce thin ﬁlms of active material for lithium ion battery application [28–32].TiO 2was mostly studied system in this category [28,31,32].Recently SnO 2[29]and V 2O 5[32]were also prepared by ALD technology and successfully used as active material for lithium ion batteries.In this paper,we report a new ALD chemistry to deposit amorphous molybdenum sulphide.Film growth was studied by in-situ quartz crystal microbalance (QCM),ex-situ X-ray re ﬂectivity (XRR)and Ellipsometry measurements.A narrow ALD temperature window between 155–170 C with characteristic saturated growth of ca.2.5Åper cycle was observed under self-limiting condition.In-situ Fourier transform infrared spectroscopy (FTIR)was employed to understand the surface limited half-reactions.The as-deposited molybdenum sul ﬁde ﬁlms were found to exhibit good electrochemical activity as anode material in lithium ion battery.A stable electrochemical performance was achieved with 375nm of ﬁlm thickness.Here we emphasize on a carbon and binder free electrode that showed a well comparable electrochemical performance at par with the existing literature.*Corresponding author.Tel.:+912225767846;fax:+912225764890.E-mail address:shaibal.sarkar@iitb.ac.in (S.K.Sarkar).1Contributed equally./10.1016/j.electacta.2014.09.0770013-4686/ã2014Elsevier Ltd.All rights reserved.Electrochimica Acta 146(2014)706–713Contents lists available at ScienceDirectElectrochimica Actaj o u rn a l h o m e p a g e :w w w.e l s e v i e r.c o m /l o c a t e /e l e c t a c t a2.Experimental Section2.1.Material SynthesisMoS2depositions were carried out in a custom-built hot-wall viscousﬂow ALD reactor,equipped with capacitance manometer and rotary vent pump.A constant pressure of1Torr was maintained byﬂowing ca.350sccm of high purity(99.999%) nitrogen gas during the reaction.The N2stream ensures laminar ﬂow of the reactants over the substrate surface.Reactants were introduced into the reactor with the help of a combination of metered valves and computer controlled pneumatic valves connected in series.This arrangement ensured controlled dosage of the reactants into the reactor.The precursors were kept at room temperature but fed to the reactor along differentially heated channels.Mo(CO)6was dosed with an overhead N2ﬂow arrangement.During the reaction,a minimum of15s purging time was maintained between two consecutive doses.2.2.Material characterizationsFilm deposition was monitored in-situ with quartz crystal microbalance(QCM).QCM measures change in the resonance frequency when mass is deposited on the crystal.The deposited mass was derived using Sauerbrey equation.We used6MHz AT cut QCM crystal with polished gold coating that was glued to the drawer and retainer assembly purchased from Inﬁcon.Under steady state condition,a mass resolution of<1nm was obtained. Before any deposition,100cycles of ALD grown Al2O3was deposited to keep the starting surface similar during different reactions.In-situ Fourier transformed infrared spectroscopy was per-formed to determine the surface species during each ALD half reactions.KBr pellets were used as substrates for this in-situ FTIR measurement.The measurement was carried out in a different reactor with similar conﬁguration apart from the ZnSe windows through which the IR beam was passed.The reactor was isolated from the window and the spectrometer by a combination of pneumatic gate valves.The transmitted beams were recorded with a liquid N2cooled MCT detector.The measurements were performed with Bruker Vertex-70FTIR spectrophotometer.X-ray reﬂectivity measurements were executed with Bruker D8Advanced diffractometer equipped with Cu-K a source.For reﬂectivity measurements,ﬁlms were deposited on polished Si surface to minimize the interfacial roughness while for the diffraction measurements standard microscope glass substrates were used to eliminate any substrate effect.Experimentally obtained XRR data wasﬁtted with the help of commercially available software,Parratt32to get the thickness and the density of theﬁlm.The thickness measurements thus obtained from the XRR were further veriﬁed with spectroscopic ellipsometry(Sentech) measurements at an incident angle of70degree.All elemental analysis was performed with X-ray photoelectron spectroscopy measurements in Thermo VG Scientiﬁc photoelectron spectrome-ter(MultiLab)equipped with Al-K a source(1486.6eV).Polished Si was used as a substrate for preparing theﬁlms for XPS measure-ments.Peaksﬁtting and analysis was done using XPS peak 4.1software.Raman spectroscopy was performed in a Raman spectrometer(Jobin Yvon HR8000)having514.5nm laser toﬁnd the different vibrational modes of the as depositedﬁlms.Films are grownon microscopic glass substrates for Raman spectroscopic measurements.2.3.Cell fabrication and electrochemical characterizationConventional charge-discharge measurements were performed under galvanostatic condition in a coin cell assembly with a cell conﬁguration of Li/Electrolyte/MoS2.Theﬁlm was directly grown on a stainless steel substrate under saturated condition that was used as an anode without any further modiﬁcations.All cells were assembled under Argon atmosphere inside glove box(Mbraun, Germany).Thin Li foil was used for both reference and counter electrodes.A well-documented standard electrolyte consisting of 1M LiPF6in EC/DMC(1:1wt/wt)(LP-30,Merck,Germany)was used for all electrochemical measurements.Watman borosilicate glass microﬁberﬁlters were used as a separator between the electrodes.Cyclic voltammetry was obtained with a scan rate of 0.2mV sÀ1using Biologic VMP-3within potential window3V to 0.01V vs.Li/Li+.Electrochemical impedance spectroscopy(EIS) measurements were carried out at open circuit voltage for the frequency range of1MHz to0.1Hz using Biologic VMP-3. Electrochemical charge-discharge measurements were performed with Arbin Instruments at constant current density of20m A cmÀ2. During all electrochemical measurements,a constant temperature of20 C was maintained.3.Results and DiscussionChemical vapor deposition(CVD)of molybdenum sulﬁde using Mo(CO)6and H2S was reported earlier[33].The thermal CVD reaction can be read as follows,Mo(CO)6+H2S!MoS2+n H2+m COFor atomic layer deposition,the above reaction can be split in to two separate half-reactionsMo(CO)x*+H2S!Mo-SH y*+n H2+m CO(A)Mo-SH y*+Mo(CO)6!MoS-Mo(CO)x*+n H2+m CO(B)Where*denotes the surface species.The above pair of surface reactions predicted at this point, however does not obsolete the possibilities of other reaction pathways.From the as predicted half reactions,it can be wellFig.1.Schematic of a layered MoS2structure.D.K.Nandi et al./Electrochimica Acta146(2014)706–713707understood that the ALD reactions are surface limited.Gaseous precursors are dosed into the reaction chamber that reacts with the surface species.Reactions continue till all the surface species gets converted to a newer functionality.The unconsumed reactants and the reaction products are simultaneously pumped out from the reaction chamber.The rate limiting surface reaction thus ensures the layer-by-layer growth of the material.As mentioned earlier,repeating the AB cycles sequentially felicitates MoS 2deposition.During the course of these reactions,the Mo(CO)6and H 2S exposures and purging time was maintained by time sequence n*1s-t 1-m*1s-t 2where n and m are the number of doses for Mo(CO)6and H 2S respectively while t 1and t 2are the purge times,which is invariably set for 15s,unless otherwise speci ﬁed.The dose time for Mo(CO)6and H 2S were set for 1s,where n and m,the respective number of exposures were varied from 1to 5.To validate that the nature of growth follows the self saturating mechanism,mass change per cycle was measured using QCM with increase in the number of reactant doses.Mo(CO)6being the solid reactant,we preferred to vary the number of doses rather than the dose time.Fig.2a shows the deposition growth rate versus the number of Mo(CO)6precursor doses,measured in-situ with QCM and veri ﬁed with ellipsometry and XRR measurements.For ellipsometry and XRR measurements,ﬁlms were grown on Si wafer with native SiO 2.In case of the QCM measurements,the mass gain was calculated from the change in oscillation frequency of the quartz crystal using Sauerbrey's equation.The QCM was also calibrated with the known deposition of Al 2O 3using trimethyl aluminum (TMA)and H 2O prior to start all of our measurementsrelated to MoS 2deposition.The growth per cycle was observed to level off at 2.5Åper cycle corresponding to 3–4metal precursor doses that yielded a total dose of 18x106L where L stands for Langmur,the unit of precursor exposure into the reactor (1L is equivalent to the exposure of 10À6Torr during 1s).Further increase in the reactants above the limiting value did not provide any extra mass gain indicating the self-limiting behaviour in the growth,as found typically for ALD.Similar characteristic was observed with varied H 2S pulses that saturated from the ﬁrst pulse itself (not shown here).The saturation dose for H 2S was found to be ca.2.5Â106L.The self-limiting nature of the ALD reaction from QCM was also veri ﬁed independently by spectroscopic ellipsom-etry and XRR measurements.Experimentally obtained Kiessig fringes with ﬁtting model that include 20ÅSiO 2and the metal sul ﬁde layer is shown in the inset of ﬁgure 2a.The program ﬁts individual densities and thicknesses of individual layer separately.The material density as calculated from the scattering length density was found to be 3.85g cm À3,which was substantially lower than the bulk value of MoS 2.This discrepancy can be attributed to the amorphous nature of the deposit,discussed below.Furthermore,the study was extended to understand the temperature dependence of the ﬁlm growth with in-situ QCM.The ALD temperature window was found out by QCM measure-ments with the saturating pulsing condition (4*1s-15s-1s-15s)as observed from self-limiting growth (Fig.2a).A relatively narrow ALD temperature window was observed within the temperature range of 155–170 C,as shown in Fig.2b with an average growth rate of 2.4Åper ALD cycle.Similar temperature window was alsoFig.2.(a)Growth rate of ﬁlm with the increasing number of Mo(CO)6precursor exposure measured by QCM,XRR and Ellipsometry (in inset XRR scan of MoS 2ALD ﬁlm on a silicon substrate after 200cycles at 170 C that yields a ﬁlm thickness of 10nm)and (b)Growth per ALD cycle with increasing substrate temperature.Fig.3.(a)Two complete ALD cycles showing the mass gain in two individual half reactions along with their corresponding precursor pressures;(b)Mass gain vs.ALD cycles monitored by QCM at a substrate temperature 170 C pulsing condition and corresponding ﬁlm thickness calculated from mass gain (in inset).708 D.K.Nandi et al./Electrochimica Acta 146(2014)706–713obtained for an unsaturated pulsing condition (1s-15s-1s-15s)with a lower growth rate that ensured the independency of the temperature window with pulsing condition.Obtained narrow range of ALD window for molybdenum sul ﬁde deposition was found to be similar to one reported earlier for oxides using the same Mo-hexacarbonyl precursor [34].Couple of representative ALD cycle is depicted in the Fig.3a to show how the mass gain and hence the growth rate was calculated using QCM measurements.The Mo(CO)6exposure resulted in a pronounced mass gain during the ﬁrst half cycle.This mass gain was measured ca.18ng cm À2.The total mass change after H 2S exposure during the second half cycle of the reaction,was found to be ca.27ng cm À2.Considering the density of MoS 2found from the XRR measurements,explained later,as 3.85g cm À3,a growth rate of ca.0.5Åper ALD cycle was obtained that is consistent with the overall growth rate as shown in Fig.2a.Fig.3b shows the mass gain measured on Al 2O 3coated QCM crystal during ﬁrst 50cycles of deposition at 170 C for 4*1s-15s-1s-15s pulse sequences.The ﬁlm growth was found to be highly linear with the number of ALD cycles with an average growth rate of 2.4Åper cycle under saturated condition of growth for ﬁrst 50ALD cycles as shown in Fig.3b.Highly linear growth was also observed under 1s-15s-1s-15s pulsing condition proving the linear growth of the ﬁlm under both saturated and unsaturated condition.Absence of nonlinearity during the initial stages of growth indicated zero incubation period.The corresponding thickness with increase number of ALD cycles under this two different pulsing conditions are shown in the inset of Fig.3b.The linear growth characteristic was found similar to classic ALD growth chemistry.In-situ FTIR spectroscopy was carried out to validate the proposed deposition mechanism.It monitors the vibration spectra of the surface species during the molybdenum sul ﬁde ﬁlm growth.Differential FTIR absorption spectra were obtained by subtracting each spectrum from the previously recorded one.Fig.4represents the difference spectra obtained during each half reaction.Here the positive absorption represents the accumulated vibration from the surface bound chemical species,while the negative feature indicates the removal of the same.We performed all the experiments at 170 C in semi-static ﬂow mode under surface saturated condition.Before recording the spectra,we deposited 50cycles of MoS 2on KBr palette.This replicates the reaction under the linear growth regime avoiding any unwanted substrate related effects.After saturated Mo(CO)6doses,the absorption peak corresponding to C=O stretch was found at around 2002cm À1[35].The C=O stretch is due to the dissosiatively chemisorbed metal-carbonyl species.An opposite but symmetrical signature peak due to the removal of carbonyl species was found after saturated H 2S doses.This clearly indicated a complete removal of the same.Thus the appearance-disappearance of the carbonyl peak during the two-half reactions cycles of the ALD,as shown in Fig.4,clearly con ﬁrms the proposed reaction mechanism.Chemical composition of the as deposited ﬁlms was examined using X-ray photoelectron spectroscopy (XPS)equipped with Al-K a source.Fig.5shows the character-istic elemental peaks for Mo and S con ﬁrming the presence of both components in the as-deposited molybdenum sul ﬁde ﬁlm.As shown in Fig.5a,doublet peaks of Mo at 229.2and 232.3eV correspond to the binding-energy of 3d 5/2and 3d 3/2electrons respectively.For sulfur the binding energy of 2p 3/2and 2p 1/2electrons were obtained at 162.5and 163.4eV (Fig.5b).In addition to this 2p doublet,S2s peak was also observed at 226.9eV alongside the Mo3d peaks.The complete survey of the XPS in the binding range of 0-1000eV and the individual scan for C1s are given in the supporting information of this article (Fig.S1).The absence of C1s in the complete survey and a very small intense peak in the individual scan re ﬂects the almost non-existence of C in the prepared ﬁlm.Elemental analysis was obtained by calculating the relative atomic ratio from the normalized area under the peak dividedFig.4.FTIR difference spectra during ﬁrst four alternate dosing of Mo(CO)6and H 2S showing the presence and removal of C=O surface species.Fig.5.XPS of the ALD grown MoS 2ﬁlm on Si substrate at 170 C (a)Peaks correspond to S2p doublet and (b)Peaks correspond to Mo3d doublet and S2s.D.K.Nandi et al./Electrochimica Acta 146(2014)706–713709by the sensitivity factors.Elemental ratio of Mo:S was found to be1:2.1.This indicated that the resultant material was MoS2.Furthermore,as reported earlier,the peak positions of the constituent elements also conﬁrmed that the chemical nature of the deposit was MoS2[36,37]. This was further veriﬁed with Raman spectroscopy.Fig.6 depicts the sharp Raman peaks at383cmÀ1and406cmÀ1 corresponding to E12g and A1g respectively.The said signature peaks occurred due to1st order Raman vibrational modes within the S-Mo-S layers[38,39]. Fig.7a and7b shows the cross-sectional SEM image and surface AFM image of the as-depositedﬁlm on Si wafer.The samples used for both these characterizations consisted a ﬁlm of150cycles under saturated condition that should yield a thickness of ca.37.5nm according to the growth rate.The SEM showed a uniform thickness of40nm over the whole substrate and the surface AFM image RMS roughness of ca.0.65nm over1m m x1m m area was found.Similar to other transition metal oxides and sulﬁdes,electro-chemistry of MoS2against lithium is also dominated by conversion reactions[40–42].During the reaction,four moles lithium were taken by one mole of MoS2as shown below[9,11,12], MoS2+4Li=Mo+2Li2SCyclic voltammogram,as shown in Fig.8a,demonstrated a stable electrochemical performance of MoS2electrode.During the cathodic process prominent peaks were observed at0.43V that characteristically corresponds to the conversion reaction with Li [9,10].During the anodic sweep,the formation of MoS x occurred via multistep reactions,which were observed at1.16V,1.4V and 1.8V vs.Li/Li+[10,43–45].A broad shoulder was also found at2.3V that was due to the formation of Li2S to S[10–12].From the2nd cycle onwards an extra peak at1.35V was observed during the cathodic sweep.The origin of this peak is the formation of Li2S/ Li2S2from the reaction between Li and sulfur[10–12].According to literature the formation of Mo-S bond(MoS x)during the delithiation process was not the prominent process[11–16]. However,small peaks observed at$1.5and$1.8V was also evident in earlier reports which are assigned for Mo-S bond formation [9,10,43–45].In this case the peak intensities of1.16,1.4,1.8V (associated for MoS x formation)were more prominent than literature work.The actual reason for such different behavior needs to be studied further.Solid electrolyte interface(SEI)formation is prominent for this electrode which was usually seen for nano-sized particle as reported earlier[46,47].Broad peaks at0.25V during cathodic sweep and at0.35V during anodic sweep was attributed due to formation and breakage of SEI layer.Formation of SEI layer was also supported by electrochemical impedance spectroscopy shown in ﬁgure8b.EIS was taken at OCV which was recorded at2.7V.Two semicircles were observed followed by a linear regime.The semicircle at high frequency regime can be ascribed to the surface ﬁlm impedance,while the other semicircle at relatively lower frequency regime can be ascribed to the charge transfer resistance [48,49].The corresponding equivalent circuit is depicted in inset of ﬁgure8b.R1indicates the electrolyte resistance.R2and Q1 represent the resistance of the Li-ion migration process and capacitance of SEI layer respectively.Similarly R3and Q2are designated to the charge transfer resistance and double layer capacitance respectively.W4represents the diffusion controlled Warburg impedance.The excellent correlation between the simulated and experimentally obtained data indicated the accuracy of the model within the experimental error limit.Charge discharge proﬁle was obtained at a constant current rate of20m A cmÀ2for the potential window of3.0to0.01V vs.Li/Li+.As shown in Figure9a the charge discharge proﬁle are in good agreement with the cyclic voltammetry results.One prominent plateau(around0.5V)was observed during theﬁrst discharge process,which was denoted for conversion reaction of MoS2with Li.The plateau at ca.1.4V in the charge proﬁle was due to the reverse reaction.During the2nd cycle another plateau was observed at ca.1.5V in the discharge process which was assignedFig.6.Raman shifts of the as-depositedﬁlm showing the characteristics peaks ofMoS2.Fig.7.(a)Cross-section SEM image of as deposited MoS2ﬁlm on Si wafer showing an uniformﬁlm throughout the substrate and(b)Surface AFM image of the same to determine the RMS roughness of the depositedﬁlm.710 D.K.Nandi et al./Electrochimica Acta146(2014)706–713for Li/S reaction [50,51].A discharge capacity of 150m Ah cm À2(corresponding to 1110mAh g À1calculated based on mass gain approximated from QCM measurements)was achieved during the ﬁrst discharge cycle that was reversible for the remaining cycles.At the end of 50cycles a discharge capacity of around 115m Ah cm À2(corresponds to 851mAh g À1)was achieved.Thus ca.77%of the initial capacity was retained at the end of 50charge-discharge cycles (Figure 9b).Often severe capacity fading is found in literature due to the volume change that occurs when a carbon and binder free pure material is used as an active electrode [52–54].However under present scenario minimal (only 23%)capacity fading was observed due to uniform deposition,low sample thickness (375nm)and more importantly the amorphous nature of the material.A coulombic ef ﬁciency in the range of 97to 98%was observed for initial 25cycles which came down to 94.4%at the end of 50cycles (Fig.9b inset).It is to be noted that in ﬁrst two cycles there was a sudden change in the coulombic ef ﬁciency as some of the side reactions were prominent in ﬁrst couple of cycles which were stabilized thereafter.The lower coulombic ef ﬁciency (or higher coulombic inef ﬁciency)could be due to the fact that the electrode used in this case is carbon and binder free electrode.Conductive carbon and polymeric binder plays an important role in lithium percolation and mechanical integrity of the electrode.As the present electrode does not contain any kind of conductive additive (such as carbon,CNTs or graphene)it is quite obvious that lithium percolation could be dif ﬁcult upon cycling as more and more resistive interfaces will be included during cycling.Similar observations were also found in literature where free standing electrodes were used [55–57].To improve the coulombic ef ﬁciency of the electrode different kind of coating were reported in the literature [56]whereas the present work demonstrates theperformance of the as deposited pristine material without any further treatment.Though it is dif ﬁcult to compare the obtained results with the literature,the current work highlights the use of ALD grown MoS 2as an active material for lithium ion battery anode.In present case a thin ﬁlm electrode was used instead of conventional tape cast method of electrode preparation.There is no literature available on the use of ALD grown MoS 2electrode in lithium ion battery,as best of our knowledge MoS 2itself is deposited for the ﬁrst time by ALD.In most of the literatures,the ALD technique was used to coat a protective layer on active material for Li-ion batteries [22–27].However,there are few reports on ALD grown materials used as electrodes in Li-ion batteries where template based (such as AAO,CNT,graphene etc.)synthesis approach has been adopted [28,29,31,32].Instead of having a carbon and binder free electrode the obtained results are very much comparable with the reported results that demonstrate a reversible capacity in the range of 850to 900mAh g À1using bare MoS 2[9–16].The higher achieved capacity can be explained by considering the other components of the electrochemical reactions.The theoretical capacity of 670mAh g À1is calculated based on the faradic reaction (diffusion controlled reactions)between Li and MoS 2where 4moles of Li reacts with MoS 2to form Li 2S and Mo.In practice,the obtained capacity from a lithium ion battery system not only comes from the faradic reactions but also from non faradic (capacitive reactions such as double layer capacitance and pseudocapacitance)[58–60]and side reactions (electrolyte decomposition and Li +adsorption on SEI layers)[61].Several researchers observed that the capacity contribution from capacitive reaction could be up to 30to 60%of the total capacity [60].In the current study we use the thin ﬁlm electrode with very less electrode loading (0.144mg cm À2)Fig.9.(a)Charge-discharge pro ﬁle and (b)Cyclic performance of MoS 2electrode at a current density of 20m A cm-2within the potential limit of 3.0V to 0.01V.Fig.8.(a)Cyclic voltammogram of MoS 2at a scan rate of 0.2mV S À1within a potential window of 3.0V to 0.01V vs.Li/Li +and (b)Electrochemical impedance spectroscopy of MoS 2at OCV and corresponding equivalent circuit.D.K.Nandi et al./Electrochimica Acta 146(2014)706–713711whereas in case of conventional electrode higher electrode loading (2–5mg cmÀ2)was used.So it is expected that the capacitive contribution should be higher in present case.Apart from low material loading,a higher order of porosity was developed with charge-discharge cycling(shown in Fig.S3in supporting informa-tion)which can improve the Li+storage via capacitive reactions. Capacity contributions from side reactions are also expected to be signiﬁcant as nanomaterials are known to possess catalytic effect in electrolyte degradation.Around15%of the total capacity is obtained from below0.2V region which occurs mainly due to Li+ adsorption in SEI layer along with irreversible solvent reduction processes.So,the capacity obtained in the range of3.0to0.2V is around723mAh gÀ1(at the end of50cycles)which is accounted for both MoS2/Li redox reaction and pseudocapacitance.Therefore, the capacity contribution only from MoS2/Li redox reaction would be lower than the theoretical capacity of670mAh gÀ1.4.ConclusionsIn this paper we successfully demonstrated amorphous Molyb-denum sulﬁde(MoS2)deposition by ALD using Mo(CO)6and H2S.A combination of complementary in-situ and ex-situ characterization techniques were employed to study the growth and self-limiting nature of the deposit.Saturated growth rate of2.5Åwas observed by QCM and veriﬁed by ellipsometry and XRR measurements.The as-deposited MoS2ﬁlm was successfully tested as lithium ion battery anode.Stable electrochemical performance was demon-strated using binder and carbon free MoS2electrode.Discharge capacity of around851mA gÀ1was achieved after50cycles at a current density rate of20m A cmÀ2.The study indicated the importance of ALD grown ultra-thinﬁlm that could be considered as freestanding electrode for lithium ion batteries.AcknowledgmentThis work was supported by the National Centre for Photovol-taic Research and Education(NCPRE)funded by the Ministry of New and Renewable Energy,Govt.of India.Appendix A.Supplementary dataSupplementary data associated with this article can be found,in the online version,at /10.1016/j.elec-tacta.2014.09.077.References[1]Q.Q.Ji,Y.F.Zhang,T.Gao,Y.Zhang,D.L.Ma,M.X.Liu,Y.B.Chen,X.F.Qiao,P.H.Tan,M.Kan,J.Feng,Q.Sun,Z.Liu,Epitaxial Monolayer MoS2on Mica with Novel Photoluminescence,Nano Lett.13(2013)3870.[2]T.W.Lin, C.J.Liu,J.Y.Lin,Facile Synthesis of MoS3/Carbon NanotubeNanocomposite with High Catalytic Activity Toward Hydrogen Evolution Reaction,Appl.Catal.B-Environ.134(2013)75.[3]R.C.Che,N.Bai,L.M.Peng,Structure and Growth of Monoclinic Mo2S3Nanorods,Appl.Phys.Lett.83(2003)3561.[4]S.Najmaei,Z.Liu,W.Zhou,X.L.Zou,G.Shi,S.D.Lei,B.I.Yakobson,J.C.Idrobo,P.M.Ajayan,J.Lou,Vapour Phase Growth and Grain Boundary Structure of Molybdenum Disulphide Atomic Layers,Nat.Mater.12(2013)754.[5]M.Shanmugam,C.A.Durcan,B.Yu,Layered Semiconductor MolybdenumDisulﬁde Nanomembrane Based Schottky-Barrier Solar Cells,Nanoscale4 (2012)7399.[6]A.N.Enyashin,M.Bar-Sadan,L.Houben,G.Seifert,Line Defects inMolybdenum Disulﬁde Layers,J.Phys.Chem.C117(2013)10842.[7]R.Thamankar,T.L.Yap,K.E.J.Goh,C.Troadec,C.Joachim,Low TemperatureNanoscale Electronic Transport on the MoS2Surface,Appl.Phys.Lett.103 (2013)083106.[8]M.Shanmugam,T.Bansal, C.A.Durcan, B.Yu,Molybdenum Disulphide/Titanium Dioxide Nanocomposite-poly3-hexylthiophene Bulk Heterojunction Solar Cell,Appl.Phys.Lett.100(2012)153901.[9]U.K.Sen,S.Mitra,High-Rate and High-Energy-Density Lithium-Ion BatteryAnode Containing2D MoS2Nanowall Cellulose Binder,ACS Appl.Mater.Interfaces5(2013)1240.[10]Y.M.Shi,Y.Wang,J.I.Wong,A.Y.S.Tan,C.L.Hsu,L.J.Li,Y.C.Lu,H.Y.Yang,Self-Assembly of Hierarchical MoS x/CNT Nanocomposites(2<x<3):Towards High Performance Anode Materials for Lithium ion Batteries,Sci.Rep.3(2013)2169.[11]M.Wang,G.D.Li,H.Y.Xu,Y.T.Qian,J.Yang,Enhanced Lithium StoragePerformances of Hierarchical Hollow MoS2Nanoparticles Assembled from Nanosheets,ACS Appl.Mater.Interfaces5(2013)1003.[12]T.Stephenson,Z.Li,B.Olsen,D.Mitlin,Lithium Ion Battery Applications ofMolybdenumdisulﬁde(MoS2)Nanocomposites,Energy Environ.Sci.7(2014) 209.[13]J.Xiao,X.Wang,X.Q.Yang,S.Xun,G.Liu,P.K.Koech,J.Liu,J.P.Lemmon,Electrochemically Induced High Capacity Displacement Reaction of PEO/ MoS2/Graphene Nanocomposites with Lithium,Adv.Funct.Mater.21(2011) 2840.[14]H.Hwang,H.Kim,J.Cho,MoS2Nanoplates Consisting of Disordered Graphene-like Layers for High Rate Lithium Battery Anode Materials,Nano Lett.11(2011) 4826.[15]S.K.Das,R.Mallavajula,N.Jayaprakash,L.A.Archer,Self-assembled MoS2–Carbon Nanostructures:Inﬂuence of Nanostructuring and Carbon on Lithium battery Performance,J.Mater.Chem.22(2012)12988.[16]G.Du,Z.Guo,S.Wang,R.Zeng,Z.Chen,H.Liu,Superior Stability and HighCapacity of Restacked Molybdenum Disulﬁde as Anode Material for Lithium Ion Batteries,mun.46(2010)1106.[17]S.M.George,Atomic Layer Deposition:An Overview,Chem.Rev.110(2010)111.[18]S.K.Sarkar,J.Y.Kim,D.N.Goldstein,N.R.Neale,K.Zhu,C.M.Elliot,A.J.Frank,S.M.George,In2S3Atomic Layer Deposition and Its Application as a Sensitizer on TiO2Nanotube Arrays for Solar Energy Conversion,J.Phys.Chem.C114(2010) 8032.[19]S.Gieraltowska,L.Wachnicki,B.S.Witkowski,E.Guziewicz,M.Godlewski,Thin Films of High-k Oxides and ZnO for Transparent Electronic Devices,Chem.Vapor Depos.19(2013)213.[20]T.K.Minton,B.H.Wu,J.M.Zhang,N.F.Lindholm,A.I.Abdulagatov,J.O'Patchen,S.M.George,M.D.Groner,Protecting Polymers in Space with Atomic Layer Deposition Coatings,ACS Appl.Mater.Interfaces2(2010)2515.[21]A.A.Dameron,S.D.Davidson,B.B.Burton,P.F.Carcia,R.S.McLean,S.M.George,Gas Diffusion Barriers on Polymers Using Multilayers Fabricated by Al2O3and Rapid SiO2Atomic Layer Deposition,J.Phys.Chem.C112(2008)4573. [22]E.M.Lotfabad,P.Kalisvaart,K.Cui,A.Kohandehghan,M.Kupsta,B.Olsen,D.Mitlin,ALD TiO2Coated Silicon Nanowires for Lithium Ion Battery Anodes with Enhanced Cycling Stability and Coulombic Efﬁciency,Phys.Chem.Chem.Phys.15(2013)13646.[23]Y.S.Jung,P.Lu,A.S.Cavanagh,C.Ban,G.H.Kim,S.H.Lee,S.M.George,S.J.Harris,A.C.Dillon,Unexpected Improved Performance of ALD Coated LiCoO2/GraphiteLi-Ion Batteries,Adv.Energy Mater.3(2013)213.[24]M.Bettge,Y.Li,B.Sankaran,N.D.Rago,T.Spila,R.T.Haasch,I.Petrov,D.P.Abraham,Improving High-capacity Li1.2Ni0.15Mn0.55Co0.1O2-Based Lithium-Ion Cells by Modiﬁying the Positive Electrode with Alumina,J.Power Sources 233(2013)346.[25]J.Zhao,Y.Wang,Ultrathin Surface Coatings for Improved ElectrochemicalPerformance of Lithium Ion Battery Electrodes at Elevated Temperature,J.Phys.Chem.C116(2012)11867.[26]B.Díaza,J.Swiatowska,V.Mauricea,A.Seyeuxa,E.Härkönenb,M.Ritalab,S.Tervakangas,J.Kolehmainenc,P.Marcus,Tantalum Oxide Nanocoatings Prepared by Atomic Layer and Filtered Cathodic Arc Deposition for Corrosion Protection of Steel:Comparative Surface and Electrochemical Analysis, Electrochim.Acta90(2013)232.[27]J.Liua,X.Lia,M.Caib,R.Lia,X.Sun,Ultrathin Atomic Layer Deposited ZrO2Coating to Enhance the Electrochemical Performance of Li4Ti5O12as an anode material,Electrochim.Acta93(2013)195.[28]X.Meng,X.Q.Yang,X.Sun,Emerging Applications of Atomic Layer Depositionfor Lithium-Ion Battery Studies,Adv.Mater.24(2012)3589.[29]X.Li,X.Meng,J.Liu,D.Geng,Y.Zhang,M.N.Banis,Y.Li,J.Yang,R.Li,X.Sun,M.Cai,M.W.Verbrugge,Tin Oxide with Controlled Morphology and Crystallinity by Atomic Layer Deposition onto Graphene Nanosheets for Enhanced Lithium Storage,Adv.Funct.Mater.22(2012)1647.[30]X.Chen,E.Pomerantseva,K.Gregorczyk,R.Ghodssi,G.Rubloff,Cathodic ALDV2O5Thin Films for High-rate Electrochemical Energy Storage,RSC Adv.3 (2013)4294.[31]K.Gerasopoulos,X.Chen,J.Culver,C.Wang,R.Ghodssi,Self-assembled Ni/TiO2Nanocomposite Anodes Synthesized via Electroless Plating and Atomic Layer Deposition on Biological Scaffolds,mun.46(2010)7349.[32]S.K.Panda,Y.Yoon,H.S.Jung,W.S.Yoon,H.Shin,Nanoscale Size Effect of Titania(Anatase)Nanotubes with Uniform Wall Thickness as High Performance Anode for Lithium-ion Secondary Battery,J.Power Sources204(2012)162.[33]J.Mun,D.Kim,J.Yun,Y.Shin,S.Kang,T.Kim,Chemical Vapor Deposition ofMoS2Films,ECS Trans.58(2013)199.[34]M.Diskus,O.Nilsen,H.Fjellvag,Growth of Thin Films of Molybdenum Oxideby Atomic Layer Deposition,J.Mater.Chem.21(2011)705.[35]M.Kurhinen,T.Venäläinen,T.A.Pakkanen,Subcarbonyl Species of Molybde-num Hexacarbonyl Supported on Silica,J.Phys.Chem.98(1994)10237. [36]C.J.Liu,S.Y.Tai,S.W.Chou,Y.C.Yu,K.D.Chang,S.Wang,F.S.S.Chien,J.Y.Lin,T.W.Lin,Facile Synthesis of MoS2/Graphene Nanocomposite with High Catalytic Activity Toward Triiodide Reduction in Dye-Sensitized Solar Cells,J.Mater.Chem.22(2012)21057.712 D.K.Nandi et al./Electrochimica Acta146(2014)706–713。

使用伪氨基酸组成和BP神经网络预测类弹性蛋白多肽的相变温度

使用伪氨基酸组成和BP神经网络预测类弹性蛋白多肽的相变温度黄凯宗;张光亚【摘要】根据获得的16条ELP序列及相变温度的数据,利用伪氨基酸组成方法提取其序列特征值.将伪氨基酸组成中的相关系数部分作为类弹性蛋白的特征向量,从类弹性蛋白序列出发,利用最小中位方差回归,找出与其序列相关系数的最佳阶数.运用均匀设计法,分别对支持向量机与BP神经网络参数进行优化.结果表明:BP神经网络获得的预测模型最佳,相变温度绝对误差为0.39℃,均方根误差为0.89℃.【期刊名称】《华侨大学学报（自然科学版）》【年(卷),期】2011(032)002【总页数】4页(P194-197)【关键词】类弹性蛋白;相变温度;伪氨基酸组成方法;支持向量机;BP神经网络【作者】黄凯宗;张光亚【作者单位】华侨大学,化工学院,福建,泉州,362021;华侨大学,化工学院,福建,泉州,362021【正文语种】中文【中图分类】Q516.02类弹性蛋白多肽(Elastin-Like Polypep tides,ELPs)是一种具有弹性功能且对环境非常敏感的生物高分子,它由五肽重复序列单元构成.如果环境温度低于ELP的相变温度,则该多肽在水溶液中是高度可溶的,聚合物链就保持无序结构,且相当伸展;反之,当环境温度高于相变温度时,这一含水的多肽链结构就会瓦解,并开始聚集,形成一个富含 ELPs的聚集物[1].利用类弹性蛋白的可逆相变特性,使其在蛋白纯化、药物载体、组织工程等方面得到广泛的应用[2].U rry等[3]认为,相变温度是关于 ELP序列、多肽链长度、Xaa种类摩尔分数的函数.Chilkoti等[4]利用重组基因进行克隆表达,得到了在序列和多肽链长均能精确控制的ELP.他们用非线性回归分析描述了ELP序列链长及浓度与相变温度的关系,但所得到的模型仅能预测3种ELP文库的相变温度.本文根据获得的16条ELP序列及相变温度的数据,利用伪氨基酸组成方法提取其序列特征值,采用BP神经网络、支持向量机方法、最小中位方差回归预测ELP 的相变温度值.1.1 试验数据来源文中所用的数据取自于文献[5].1.2 伪氨基酸组成伪氨基酸组成包含20+λ个变量,最早由Chou等[6]提出.由于文中所涉及的ELP氨基酸组成极为相似,而且种类很少,为了减少输入变量数目,对其略作调整,仅取其后的λ个变量,即氨基酸相关系数. ELP相关系数的阶数λ从1取到10,氨基酸相关系数计算参见文献[7].1.3 均匀设计在运行时,支持向量机(SVM)[8]和BP神经网络[9]都需要选择参数,以达到最佳效果.因此,采用均匀设计法(UD)[10]来选择适当的运行参数.定义3个特征指标[11],即平均绝对百分比误差δMPAE、均方根误差δMSE和平均绝对误差δMAE.模型预测的结果采用常用的“留一法”,即对 n组数据,每次取1组作测试,其他n-1组作为训练样本,共进行 n次循环,使得样本中所有数据都能进行预测.2.1 氨基酸相关系数的阶数的选择根据文献[6],氨基酸相关系数的阶数(λ)是伪氨基酸组成一重要参数.文献数据的相变温度呈离散分布,使用最小中位方差回归会更为精确 [11-12],且运行过程中无需调整参数.参数λ经最小中位方差(Least Median of Squares Regression,LM SQ)回归检测,获得的平均绝对百分比误差δMPAE、均方根误差δMSE和平均绝对误差δMAE关系,如表1所示.由表1可知,当λ=8时,δMAE为3.04,δMSE为5.73,δMPAE为40.91%.即拟合所得ELP相变温度准确率最高,因此取λ=8.当λ=8时,执行最小中位方差回归得到ELP的相变温度拟合模型为其中:x1～x8分别为伪氨基酸组中相关系数;x9～x10分别为ELP的相对分子质量、ELP每一单体的Xaa数量;ELP浓度对ELP相变温度没有影响,故为其相关系数零.从模型(1)可见,第1,第4和第6个相关系数对相变温度有较大的负面影响,而第5个相关系数则有较大的正面影响;伪氨基酸组的相关系数对ELP的相变温度影响较大.当ELP浓度较高时,其浓度在一定范围变化对相变温度几乎不影响.这与Chilkoti 等[4]的实验结果较为一致.使用最小中位方差回归获得的拟合值与实测值关系,如图1所示.由图1可知,一些拟合值非常好,而另外一些预测值与实测值差距比较大,从而导致其回归直线的斜率偏离较大.2.2 利用支持向量机预测相变温度如前所述,λ=8为氨基酸相关系数的阶数最佳运行参数.利用均匀设计法对支持向量机的运行参数进行优化,交叉验证后的结果如表2所示.由表2可得出,3个误差特征指标在交叉验证中变化的幅度较小.这说明SVM对运行的参数不是很敏感.当惩罚系数C=100,ε为1.0×10-5,γ为0.3 (即方案7)时,其δMAE,δMSE和δMP AE值均最小,分别为1.85,3.31和23.39%.即所建立的模型对 ELP相变温度预测准确率最高,故为最佳方案.在方案7中,使用用支持向量机方法建立相变温度模型.通过该模型对实际测得的数据进行预测,预测的效果,如图2所示.从图2可知,模型预测的结果与实际测量值的相关系数达0.93,模型预测的结果较好.2.3 利用神经网络预测相变温度对神经网络而言,由于训练样本集的大小有限,网络训练后对训练集外的输入的响应,直接决定网络的性能.为了检验所建立的神经网络的可靠性,对其进行3因素9水平交叉验证,结果如表3所示.从表3可知,3个特征值变化幅度较大,神经网络对运行参数比较敏感.在9组验证中,采用默认参数获得的特征值最好.即隐含层节点数(n)为6,学习速率(v)为0.3,动态参数(σ)为0.2时,准确率最高,其δMAE,δMSE和δMPAE值均最小,分别为0.39,0.89和4.86%.用BP神经网络建立的相变温度模型.通过该模型对实际测得的数据进行预测,结果如图3所示.从图3可知,模型预测的结果与实际测量值的相关系数达0.99.由图1～3可知,BP神经网络所建立的预测相变温度的精度,比使用支持向量机和最小中位方差回归建立的相变温度要好,可作为后续使用的模型.当实测的ELP相变温度为60℃(此时ELP的序列最短浓度最高),与3种算法所预测(回归的结果是拟合的)出来相变温度值均差距较大.这可能是因为当序列较短时,ELP 浓度与长度的变化对相变温度影响更大[4],而ELP的序列组成对相变温度影响较小. 与传统的拟合方法预测ELP的相变温度相比,基于支持向量机和神经网络对相变温度进行预测,不用通过预测相变温度具体形式,就可以直接从数据中得到相变温度与ELP序列、分子量、Xaa组成、浓度之间的关系.同时,只要能加以一定的先验知识,还能够更大范围地反映它们之间的关系,其应用的范围也将更为广阔.文中基于Chou等提出的伪氨基酸概念[6],考虑到ELP的氨基酸组成极为相似,构造了一种λ维的伪氨基酸组成来表示蛋白质序列.采用BP神经网络、支持向量机方法、最小中位方差回归预测ELP的相变温度值.结果表明,当λ=8为氨基酸相关系数的阶数最佳运行参数时,使用BP神经网络所建立的相变温度预测模型为最佳.【相关文献】[1]URRYDW.Physical chemistry of biological free energy transduction as demonstrated by elastic protein-based polymers[J].Phys Chem(B),1997,101(51):11007-11028.[2]CHOW D,NUNALEE M L,CH IL KOTIA,et al.Pep tide-based biopolymers in biomedicine and biotechnology [J].Mater Sci Eng R Rep,2008,62(4):125-155.[3]URRYD W,LUAN C H,PARKER T M,et al.Temperature of polypep tide inverse temperature transition depends on mean residue hydrophobicity[J].J Am ChemSoc,1991,113(11):4346-4348.[4]M EYER D E,CH ILKOTIA.Quantification of the effects of chain length and concentration on the thermal behavior of elastin-like polypep tides[J].Biomacromolecules,2004,5(3):846-851.[5]OlSON SD.Mathematical models for analysisof tissue regeneration in articular cartilage[D].No rth Carolina State: North Carolina State University,2009.[6]CHOU Kuo-chen.Prediction of protein cellular attributes using pseudo amino acid composition[J].Proteins:Structure,Function,and Bioinfo rmatics,2001,43(3):246-255. [7]SHEN Hong-bin,CHOU Kuo-chen.PseAAC:A flexible web-server for generating various kinds of protein pseudo amino acid composition[J].AnalyticalBiochemistry,2008,373(2):386-388.[8]VANPN IK V N.The nature of statistical learning theory[M].New York:Sp ringer-Verlag,1995.[9]黄永恒,曹平,汪亦显.基于BP神经网络的岩土工程预测模型研究[J].科技导报,2009,27(6):61-64.[10]方开泰.均匀设计:数论方法在试验设计的应用[J].应用数学学报,1980(3):363-372.[11]张光亚,葛慧华,方柏山.一种预测木聚糖酶最适温度的PCANN模型[J].华侨大学学报:自然科学版,2007,28 (1):55-58.[12]ROUSSEEUW PJ.Leastmedian of squares regression[J].Journal of the American Statistical Association,1984,79 (388):871-880.[13]STEELE JM,STEIGERW L.Algorithms and complexity for least median of squares regression[J].Discrete Applied Mathematics,1986,14(1) :93-100.。

以大豆胰蛋白酶抑制因子为研究对象

实验流程
交联壳聚糖床的制备环氧氯丙烷对交联壳聚糖床的活化胰蛋白酶的固定化 SBTI粗品的制备天然胰蛋白酶酶活鉴定固定化胰蛋白酶酶活鉴定 SBTI的亲和分离纯化产品效果的SDS-PAGE鉴定
Scheme 1
Crosslinking of chitosan by glutaraldehyde

SBTI属于丝氨酸蛋白酶抑制因子，包括 Kunitz trypsin inhibitor(KTI)和Bowman-Birk trypsin inhibitor(BBTI)两类，前者 Mr~20kDa,后者Mr~8-10kDa

我们从大豆研磨液中分离出来的BBTI是一种强力的胰凝乳蛋白酶抑制剂，纯化的 BBTI经研究表明，BBTI还具有杀菌活性，抑制HIV-1的反转录活性。此外，在 Sprgue-Dawley鼠中发现，它还能有效抑制乙酰基亚硝基脲引起的神经系统瘤.
Affinity chromatography of SBTI by trypsin-immobilised chitosan beads (0.2mol L−1 borax/borate (pH 7.8) as equilibrium and washing buffer, HCl/H2O solution as eluent, pH gradient elution at flow rate of 30mL h−1, 7.5mL per tube).
Effect of pH on SBTI adsorption by free and immobilised trypsin.
Effect of temperature on SBTI adsorption by free and immobilised trypsin.

Thermo-calc软件-TCCP用户指南 (有目录索引)

Thermo-Calc®User’s GuideVersion PThermo-Calc Software ABStockholm Technology ParkBjörnnäsvägen 21SE-113 47 Stockholm, SwedenCopyright © 1995-2003 Foundation of Computational ThermodynamicsStockholm, Sweden目录第1部分一般介绍 (12)1.1 计算热力学 (12)1.2 Thermo-Calc软件/数据库/界面包 (12)1.3 致谢 (13)1.4 版本历史 (13)1.5 Thermo-Calc软件包的通用结构 (13)1.6 各类硬件上Thermo-Calc软件包的有效性 (14)1.7 使用Thermo-Calc软件包的好处 (14)第2部分如何成为Thermo-Calc专家 (14)2.1 如何容易地使用本用户指南 (14)2.2 如何安装和维护Thermo-Calc软件包 (16)2.2.1 许可要求 (16)2.2.2 安装程序 (16)2.2.3 维护当前和以前版本 (16)2.2.4 使TCC执行更方便 (16)2.3 如何成为Thermo-Calc专家 (16)2.3.1 从TCSAB与其世界各地的代理获得迅速技术支持 (17)2.3.2 日常使用各种Thermo-Calc功能 (17)2.3.3 以专业的和高质量的标准提交结果 (17)2.3.4 通过各种渠道相互交换经验 (17)第3部分Thermo-Calc软件系统 (17)3.1 Thermo-Calc软件系统的目标 (17)3.2 一些热力学术语的介绍 (18)3.2.1 热力学 (18)3.2.2 体系、组元、相、组成、物种（System, component, phases, constituents and species） (18)3.2.3 结构、亚点阵和位置 (19)3.2.4 成分、构成、位置分数、摩尔分数和浓度(composition, constitution, site fractions, molefractions and concentration) (19)3.2.5 平衡态和状态变量 (19)3.2.6 导出变量 (22)3.2.7 Gibbs相规则 (25)3.2.8 状态的热力学函数 (25)3.2.9 具有多相的体系 (25)3.2.10 不可逆热力学 (26)3.2.11 热力学模型 (26)3.2.12 与各种状态变量有关的Gibbs能 (27)3.2.13 参考态与标准态 (27)3.2.14 溶解度范围 (28)3.2.15 驱动力 (28)3.2.16 化学反应 (28)3.2.17 与平衡常数方法相对的Gibbs能最小化技术 (28)3.2.18 平衡计算 (29)3.3 热力学数据 (30)3.3.1 数据结构 (30)3.3.3 数据估价 (32)3.3.6 数据加密 (33)3.4 用户界面 (34)3.4.1 普通结构 (34)3.4.2 缩写 (34)3.4.3 过程机制（history mechanism） (35)3.4.4 工作目录和目标目录(Working directory and target directory) (35)3.4.5 参数转换为命令 (36)3.4.6 缺省值 (36)3.4.7 不理解的问题 (36)3.4.8 帮助与信息 (36)3.4.9 出错消息 (36)3.4.10 控制符 (36)3.4.11 私人文件 (36)3.4.12 宏工具 (37)3.4.13 模块性 (37)3.5 Thermo-Calc中的模块 (37)3.5.1 基本模块 (37)3.7 Thermo-Calc编程界面 (39)3.7.1 Thermo-Calc作为引肇 (39)3.7.2 Thermo-Calc应用编程界面：TQ和TCAPI (40)3.7.3 在其它软件包中开发Thermo-Calc工具箱 (43)3.7.4 材料性质计算核材料工艺模拟的应用 (43)3.8 Thermo-Calc的功能 (44)3.9 Thermo-Calc应用 (44)第4部分Thermo-Calc数据库描述 (45)4.1 引言 (45)4.2 Thermo-Calc数据库描述形式 (45)第5部分数据库模块（TDB）——用户指南 (55)5.1 引言 (55)5.2 TDB模块中用户界面 (56)5.3 开始 (56)5.3.1 SWITCH-DATABASE (56)5.3.2 LIST-DATABASE ELEMENT (56)5.3.3 DEFINE_ELEMENTS (56)5.3.4 LIST_SYSTEM CONSTITUENT (56)5.3.5 REJECT PHASE (56)5.3.6 RESTORE PHASE (56)5.3.7 GET_DATA (56)5.4 所有TDB监视命令的描述 (56)5.4.1 AMEND_SELACTION (56)5.4.6 DEFINE_SPECIES (58)5.4.7 DEFINE_SYSTEM (58)5.4.8 EXCLUDE_UNUSED_SPECIES (58)5.4.9 EXIT (58)5.4.10 GET_DATA (58)5.4.11 GOTO_MODULE (59)5.4.12 HELP (59)5.4.13 INFORMA TION (59)5.4.14 LIST_DATABASE (60)5.4.15 LIST_SYSTEM (60)5.4.16 MERGE_WITH_DA TABASES (61)5.4.17 NEW_DIRECTORY_FILE (61)5.4.18 REJECT (61)5.4.19 RESTORE (62)5.4.20 SET_AUTO_APPEND_DA TABASE (62)5.4.21 SWITCH_DA TABASE (63)5.5 扩展命令 (64)第6部分数据库模块（TDB）——管理指南 (64)6.1 引言 (64)6.2 TDB模块的初始化 (65)6.3 数据库定义文件语法 (66)6.3.1 ELEMENT (67)6.3.2 SPECIES (67)6.3.3 PHASE (67)6.3.4 CONSTITUENT (67)6.3.5 ADD_CONSTITUENT (68)6.3.6 COMPOUND_PHASE (68)6.3.7 ALLOTROPIC_PHASE (68)6.3.8 TEMPERA TURE_LIMITS (68)6.3.9 DEFINE_SYSTEM_DEFAULT (69)6.3.10 DEFAULT_COMMAND (69)6.3.11 DATABASE_INFORMATION (69)6.3.12 TYPE_DEFINITION (69)6.3.13 FTP_FILE (70)6.3.14 FUNCTION (70)6.3.15 PARAMETER (72)6.3.16 OPTIONS (73)6.3.17 TABLE (73)6.3.18 ASSESSED_SYSTEMS (73)6.3.19 REFERENCE_FILE (74)6.3.20 LIST_OF_REFERENCE (75)6.3.21 CASE与ENDCASE (76)6.3.22 VERSION_DA TA (76)6.5 数据库定义文件实例 (77)6.5.1 例1：一个小的钢数据库 (77)6.5.2 例2：Sb-Sn系个人数据库 (78)第7部分制表模块（TAB） (81)7.1 引言 (81)7.2 一般命令 (81)7.2.1 HELP (81)7.2.2 GOTO_MODULE (81)7.2.3 BACK (82)7.2.4 EXIT (82)7.2.5 PATCH (82)7.3 重要命令 (82)7.3.1 TABULATE_SUBSTANCE (82)7.3.2 TABULATE_REACTION (85)7.3.3 ENTER_REACTION (86)7.3.4 SWITCH_DA TABASE (87)7.3.5 ENTER_FUNCTION (88)7.3.6 TABULATE_DERIV A TIVES (89)7.3.7 LIST_SUBSTANCE (91)7.4 其它命令 (92)7.4.1 SET_ENERGY_UNIT (92)7.4.2 SET_PLOT_FORMAT (92)7.4.3 MACRO_FILE_OPEN (92)7.4.4 SET_INTERACTIVE (93)7.5 绘制表 (93)第8部分平衡计算模块（POL Y） (94)8.1 引言 (94)8.2 开始 (95)8.3 基本热力学 (95)8.3.1 体系与相 (95)8.3.2 组元（Species） (95)8.3.3 状态变量 (96)8.3.4 组分 (97)8.3.5 条件 (98)8.4 不同类型的计算 (98)8.4.1 计算单一平衡 (98)8.4.2 性质图的Steping计算 (99)8.4.3 凝固路径模拟 (99)8.4.4 仲平衡与T0温度模拟 (99)8.4.5 相图的Mapping计算 (101)8.4.6 势图计算 (101)8.4.7 Pourbaix图计算 (101)8.4.8 绘制图 (101)8.5.4 更高阶相图 (104)8.5.5 性质图 (104)8.6 普通命令 (104)8.6.1 HELP (104)8.6.2 INFORMA TION (104)8.6.3 GOTO_MODULE (105)8.6.4 BACK (105)8.6.5 SET_INTERACTIVE (105)8.6.6 EXIT (106)8.7 基本命令 (106)8.7.1 SET_CONDITION (106)8.7.2 RESET_CONDITION (107)8.7.3 LIST_CONDITIONS (107)8.7.4 COMPUTE_EQUILIBRIUM (107)8.7.6 DEFINE_MATERIAL (108)8.7.6 DEFINE_DIAGRAM (111)8.8 保存和读取POL Y数据结构的命令 (112)8.8.1 SA VE_WORKSPACES (112)8.8.2 READ_WORKSPACES (113)8.9 计算与绘图命令 (114)8.9.1 SET_AXIS_V ARIABLE (114)8.9.2 LIST_AXIS_V ARIABLE (114)8.9.3 MAP (114)8.9.4 STEP_WITH_OPTIONS (115)8.9.5 ADD_INITIAL_EQUILIBRIUM (117)8.9.6 POST (118)8.10 其它有帮助的命令 (118)8.10.1 CHANGE_STA TUS (118)8.10.2 LIST_STA TUS (119)8.10.3 COMPUTE_TRANSITION (120)8.10.4 SET_ALL_START_V ALUES (121)8.10.5 SHOW_V ALUE (122)8.10.6 SET_INPUT_AMOUNTS (122)8.10.7 SET_REFERENCE_STA TE (122)8.10.8 ENTER_SYMBOL (123)8.10.9 LIST_SYMBOLS (124)8.10.10 EV ALUATE_FUNCTIONS (124)8.10.11 TABULATE (124)8.11 高级命令 (125)8.11.1 AMEND_STORED_EQUILIBRIA (125)8.11.3 DELETE_INITIAL_EQUILIBRIUM (126)8.11.4 LIST_INITIAL_EQUILIBRIA (126)8.11.5 LOAD_INITIAL_EQUILIBRIUM (126)8.11.10 SELECT_EQUILIBRIUM (128)8.11.11 SET_NUMERICAL_LIMITS (128)8.11.12 SET_START_CONSTITUTION (129)8.11.13 SET_START_V ALUE (129)8.11.14 PATCH (129)8.11.15 RECOVER_START_V ALUE (129)8.11.16 SPECIAL_OPTIONS (129)8.12 水溶液 (132)8.13 排除故障 (133)8.13.1 第一步 (133)8.13.2 第二步 (133)8.13.3 第三步 (133)8.14 频繁提问的问题 (134)8.14.1 程序中为什么只得到半行？ (134)8.14.2 在已经保存之后为什么不能绘图？ (134)8.14.3 为什么G.T不总是与-S相同？ (134)8.14.4 如何获得组元偏焓 (135)8.14.5 为什么H(LIQUID) 是零而HM(LIQUID)不是零 (135)8.14.6 即使石墨是稳定的为什么碳活度小于1？ (135)8.14.7 如何获得过剩Gibbs能？ (135)8.14.8 当得到交叉结线而不是混溶裂隙时什么是错的？ (135)8.14.9 怎么能直接计算最大混溶裂隙？ (136)第9部分后处理模块（POST） (136)9.1 引言 (136)9.2 一般命令 (137)9.2.1 HELP (137)9.2.2 BACK (137)9.2.3 EXIT (137)9.3 重要命令 (137)9.3.1 SET_DIAGRAM_AXIS (137)9.3.2 SET_DIAGRAM_TYPE (138)9.3.3 SET_LABEL_CORVE_OPTION (139)9.3.5 MODIFY_LABEL_TEXT (139)9.3.6 SET_PLOT_FORMAT (140)9.3.7 PLOT_DIAGRAM (141)9.3.8 PRINT_DIAGRAM (142)9.3.9 DUMP_DIAGRAM (143)9.3.10 SET_SCALING_STA TUS (144)9.3.11 SET_TITLE (144)9.3.12 LIST_PLOT_SETTINGS (144)9.4 实验数据文件绘图命令 (144)9.4.1 APPEND_EXPERIMENTAL_DA TA (144)9.4.2 MAKE_EXPERIMENTAL_DA TAFILE (145)9.5.3 SET_AXIS_LENGTH (147)9.5.4 SET_AXIS_TEXT_STATUS (147)9.5.5 SET_AXIS_TYPE (147)9.5.6 SET_COLOR (147)9.5.7 SET_CORNER_TEXT (148)9.5.8 SET_FONT (148)9.5.9 SET_INTERACTIVE_MODE (149)9.5.10 SET_PLOT_OPTION (149)9.5.11 SET_PREFIX_SCALING (149)9.5.12 SET_REFERENCE_STA TE (149)9.5.13 SET_TIELINE_STA TE (150)9.5.14 SET_TRUE_MANUAL_SCALING (150)9.5.15 TABULATE (150)9.6 奇特的命令 (150)9.6.1 PATCH_WORKSPACE (150)9.6.2 RESTORE_PHASE_IN_PLOT (150)9.6.3 REINIATE_PLOT_SETTINGS (151)9.6.4 SET_AXIS_PLOT_STATUS (151)9.6.5 SET_PLOT_SIZE (151)9.6.6 SET_RASTER_STATUS (151)9.6.8 SUSPEND_PHASE_IN_PLOT (151)9.7 3D图标是：命令与演示 (151)9.7.1 CREATE_3D_PLOTFILE (153)9.7.2 在Cortona VRML Client阅读器中查看3D图 (154)第10部分一些特殊模块 (155)10.1 引言 (155)10.2 特殊模块生成或使用的文件 (156)10.2.1 POL Y3文件 (156)10.2.2 RCT文件 (156)10.2.3 GES5文件 (156)10.2.4 宏文件 (157)10.3 与特殊模块的交互 (157)10.4 BIN模块 (157)10.4.1 BIN模块的描述 (157)10.4.2 特定BIN模块数据库的结构 (161)10.4.3特定BIN计算的演示实例 (162)10.5 TERN 模块 (162)10.5.1 TERN 模块的描述 (162)10.5.2 特殊TERN模块数据库的结构 (166)10.5.3 TERN模块计算的演示实例 (167)10.6 POT模块 (167)10.7 POURBAIX 模块 (167)10.8 SCHAIL 模块 (167)11.2 热化学 (168)11.2.1 一些术语的定义 (168)11.2.2 元素与物种（Elements and species） (168)11.2.3 大小写模式 (169)11.2.4 相 (169)11.2.5 温度与压力的函数 (169)11.2.6 符号 (170)11.2.7 混溶裂隙 (170)11.3 热力学模型 (170)11.3.1 标准Gibbs能 (171)11.3.2 理想置换模型 (171)11.3.3 规则溶体模型 (171)11.3.4 使用组元而不是元素 (172)11.3.5 亚点阵模型—化合物能量公式 (172)11.3.6 离子液体模型，对具有有序化趋势的液体 (172)11.3.7 缔合模型 (173)11.3.8 准化学模型 (173)11.3.9 对Gibbs能的非化学贡献（如铁磁） (173)11.3.10 既有有序-无序转变的相 (173)11.3.11 CVM方法：关于有序/无序现象 (173)11.3.12 Birch-Murnaghan模型：关于高压贡献 (173)11.3.13 理想气体模型相对非理想气体/气体混合物模型 (173)11.3.14 DHLL和SIT模型：关于稀水溶液 (173)11.3.15 HKF和PITZ模型：对浓水溶液 (173)11.3.16 Flory-Huggins模型：对聚合物 (173)11.4 热力学参数 (173)11.5 数据结构 (175)11.5.1 构造 (175)11.5.2 Gibbs能参考表面 (175)11.5.3 过剩Gibbs能 (175)11.5.4 存储私有文件 (175)11.5.5 加密与不加密数据库 (176)11.6 GES系统的应用程序 (176)11.7 用户界面 (176)11.7.1 模块性和交互性 (177)11.7.2 控制符的使用 (177)11.8 帮助与信息的命令 (177)11.8.1 HELP (177)11.8.2 INFORMATION (177)11.9 改变模块与终止程序命令 (178)11.9.1 GOTO_MODULE (178)11.9.2 BACK (178)11.9.3 EXIT (178)11.10 输入数据命令 (178)11.10.4 ENTER_SYMBOL (180)11.10.5 ENTER_PARAMETER (181)11.11 列出数据的命令 (183)11.11.1 LIST_DATA (183)11.11.2 LIST_PHASE_DA TA (183)11.11.3 LIST_PARAMETER (184)11.11.4 LIST_SYMBOL (185)11.11.5 LIST_CONSTITUENT (185)11.11.6 LIST_STATUS (185)11.12 修改数据命令 (185)11.12.1 AMEND_ELEMENT_DA TA (185)11.12.2 AMEND_PHASE_DESCRIPTION (186)11.12.3 AMEND_SYMBOL (188)11.12.4 AMEND_PARAMETER (189)11.12.5 CHANGE_STATUS (191)11.12.6 PATCH_WORKSPACES (191)11.12.7 SET_R_AND_P_NORM (191)11.13 删除数据的命令 (192)11.13.1 REINITIATE (192)11.13.2 DELETE (192)11.14 存储或读取数据的命令 (192)11.14.1 SA VE_GES_WORKSPACE (192)11.14.2 READ_GES_WORKSPACE (193)11.15 其它命令 (193)11.15.1 SET_INTERACTIVE (193)第12部分优化模块（PARROT） (193)12.1 引言 (193)12.1.1 热力学数据库 (194)12.1.2 优化方法 (194)1 2.1.4 其它优化软件 (195)12.2 开始 (195)12.2.1 试验数据文件：POP文件 (195)12.2.2 图形试验文件：EXP文件 (197)12.2.3 系统定义文件：SETUP文件 (197)12.2.4 工作文件或存储文件：PAR文件 (198)12.2.5 各种文件名与其关系 (198)12.2.6 交互运行PARROT模块 (199)12.2.6.3 绘制中间结果 (199)12.2.6.4 实验数据的选择 (199)12.2.6.6 优化与连续优化 (200)12.2.7 参数修整 (200)12.2.8 交互完成的变化要求编译 (201)12.3 交替模式 (201)12.4 诀窍与处理 (201)12.4.4 参数量 (201)12.5 命令结构 (201)12.5.1 一些项的定义 (201)12.5.2 与其它模块连接的命令 (201)12.5.3 用户界面 (201)12.6 一般命令 (201)12.7 最频繁使用的命令 (202)12.8 其它命令 (203)第13部分编辑-实验模块（ED-EXP） (203)第14部分系统实用模块（SYS） (203)14.1 引言 (203)14.2 一般命令 (203)14.2.1 HELP (203)14.2.2 INFORMA TION (204)14.2.4 BACK (205)14.2.5 EXIT (205)14.2.6 SET_LOG_FILE (205)14.2.7 MACRO+FILE_OPEN (205)14.2.8 SET_PLOT_ENVIRONMENT (206)14.3 Odd命令 (207)14.3.1 SET_INTERACTIVE_MODE (207)14.3.2 SET_COMMAND_UNITS (207)14.3.4 LIST_FREE_WORKSPACE (207)14.3.5 PATCH (207)14.3.6 TRACE (207)14.3.7 STOP_ON_ERROR (208)14.3.8 OPEN_FILE (208)14.3.9 CLOSE_FILE (208)14.3.10 SET_TERMINAL (208)14.3.11 NEWS (208)14.3.12 HP_CALCULATOR (208)14.4 一般信息的显示 (209)第15部分数据绘图语言（DATAPLOT） (215)第1部分一般介绍1.1 计算热力学在近十年内与材料科学与工程相联系的计算机计算与模拟的研究与发展已经为定量设计各种材料产生了革命性的方法，热力学与动力学模型的广泛结合使预测材料成分、各种加工后的结构和性能。

集成光学集成光学、波导理论

(150007)),赵灿//制造技术与机床.―2007,(2).―7175介绍了目前基于三维光栅扫描测量标定技术的国内外发展状况,提出自由置位固定编码多目标跟踪提取连续图像轮廓的全自动标定技术,给出了此技术的具体算法实现。

利用提出的全自动标定技术,得出了相应的实验结果,并将该实验结果与利用通用的M AR R 算法得到的结果对标定精度进行了比较、分析。

结果表明,所提出的全自动标定技术使操作更加便捷,能很好地保证标定图像轮廓边缘的定位精度,使整个系统的标定精度保持在0.005~0.02mm 范围内。

图6表2参11(于晓光)TP3912007043503点阵式测头成像视觉三坐标测量系统建模=M odeling vi sion 3D coordinate measur ing system with a special probe[刊,中]/张雪飞(天津大学精密测试技术与仪器国家重点实验室.天津(300072)),彭凯//制造技术与机床.―2007,(3).―103105在透视投影和坐标变换的基础上,建立测头成像视觉三坐标测量系统的数学模型。

求解了共线、共面和空间三种点阵式测头构成的系统模型,以最简单的共线3点型点阵测头为例建立实测视觉坐标测量系统,实验结果证明了系统模型的正确性。

图4参3(于晓光)TP391TB922007043504基于机器视觉的二维小尺寸精密测量系统=Two dimensional pr ecision measurement systems based on machine vi sion[刊,中]/罗钧(重庆大学光电技术及系统国家教育部重点实验室.重庆(400030)),黄俊//计算机测量与控制.―2007,15(1).―1113设计了一种对万能工具显微镜进行改造而成的二维精密测量系统。

该系统引入了机器视觉技术进行自动测量对准,采用光栅传感器提供坐标系以检测尺寸,测量系统测量范围为25mm !25mm,测量精度可达5m 。

生长厚度对InGaAs_InAs薄膜表面形貌的影响_王一

生长厚度对InGaAs/InAs薄膜表面形貌的影响王　一　郭　祥　刘　珂　黄梦雅　魏文　赵　振　胡明哲　罗子江　丁　召＊(贵州大学理学院电子科学系　贵阳　550025)Effect of Thickness on Microstructures of InGaAs/InAs Thin Films Wang Yi,Guo Xiang,Liu Ke,Huang Mengya,Wei Wenzhe,Zhao Zhen,Hu Mingzhe,Luo Zijiang,Ding Zhao＊(Depar tment of Electro nics,College of Science,Gui zhou Univer sity,Guiyang550025,China) Abstract　The gro wth of In0.86Ga0.14As thin films by molecular bea m epitaxiy(MBE)on InAs substrate was in-situ annealed and monitored with reflection high ener gy diffraction(RHEED).The influence of the growth conditions,the film thickness in particular,on the atomic structures of the film surfaces was characterized with scanning tunneling microscopy (STM).The results show that while slightly affecting the surface reconstruction,the thickness has a major impact on the micr ostructures and mechanical properties of the film surface.For example,as the thickness increased,the tensile inside the epitaxial layerincreased,but the average size of terrace decreased.Moreover,an increased number of steps with sharp irregular zigza g edgeswere observed,possibly because of accumulation of surface tensile and anisotropic preferred diffusion mechanisms. Keywords　MB E,InGaAs/InAs,STM,Gro wth thickness 摘要　利用分子束外延技术,通过反射式高能电子衍射仪实时监控In0.86Ga0.14As/InAs薄膜生长状况,在InAs(001)基片上生长5、10、20原子单层(ML)厚度的InGaAs单晶薄膜。

18 MeV质子辐照对TiNi形状记忆合金R相变的影响

第15卷　第1期强激光与粒子束Vol.15,No.1 2003年1月HIGH POWER LASER AND PAR TICL E B EAMS Jan.,2003　文章编号:100124322(2003)012009720418Me V质子辐照对Ti Ni形状记忆合金R相变的影响Ξ王治国,　祖小涛,　封向东,　刘丽娟,　林理彬(四川大学物理系教育部辐射物理及技术重点实验室,四川成都610064) 摘　要:　研究了用HZ2B串列加速器的18MeV质子辐照对TiNi形状记忆合金R相变的影响,辐照在奥氏体母相状态下进行。

示差扫描量热法(DSC)表明,辐照后R相变开始温度T s R和逆马氏体相变结束温度T f A随辐照注量的增加而降低。

当注量为1.53×1014/cm2时,T s R和T f A分别下降6K和13K,辐照未引起R相变结束温度T f R和逆马氏体相变开始温度T s A的变化。

表明辐照后母相(奥氏体相)稳定。

透射电镜(TEM)分析表明辐照后没有引起合金可观察的微观组织变化。

辐照对R相变开始温度T s R和逆马氏体相变结束温度Af的影响可能是由于质子辐照后产生了孤立的缺陷团,形成了局部应力场,引起晶格有序度的下降所造成的。

关键词:　TiNi形状记忆合金;质子辐照;R相变;示差扫描量热法;TEM 中图分类号:TG139.6 文献标识码:A TiNi形状记忆合金是目前应用最为广泛,也最成功的一种智能材料,集传感功能与驱动功能于一体,在核反应堆和太空等核辐射环境下用作传感与驱动元件已引起了关注[1,2]。

TiNi合金中R相变具有热滞后小,响应速度快的特点,在实际应用中得到了广泛的应用[3]。

在以前的研究中利用R相变得到了具有双向记忆效应的弹簧,伸缩率可达25%[4]。

由于核辐射会对形状记忆合金相变特性产生影响,因而研究其改变规律及机理对形状记忆合金在辐射环境下应用的可靠性和可行性是十分必要的。

MLX90620ESF-BAB-000;中文规格书,Datasheet资料

(1) Supply Voltage B = 3V
(3) Package options: A = reserved B = 60° FOV C = reserved D = 40° FOV
Example: MLX90620ESF-BAB-000-TU
Functional diagram
Digital Active Thermopile Array
16x4 IR array
General Description (continued)
The results of the infrared sensor measurements are stored in RAM: • 16-bit result of IR measurement for each individual sensor (64 words) • 16-bit result of PTAT sensor Depending on the application, the external microcontroller can read the different RAM data and, based on the calibration data stored in the EEPROM memory, compensate for difference between sensors to build up a thermal image, or calculate the temperature at each spot of the imaged scene. These constants are accessible by the user microcontroller through the I2C bus and have to be used for external post processing of the thermal data. This post processing includes: • Ta calculation • Pixel offset cancelling • Pixel to pixel sensitivity difference compensation • Object emissivity compensation • Object temperature calculation The result is an image with NETD better than 0.5K at 1Hz refresh rate. The refresh rate of the array is programmable by means of register settings or directly via I2C command. Changes of the refresh rate have a direct impact on the integration time and noise bandwidth (faster refresh rate means higher noise level). The frame rate is programmable in the range 0,5Hz…512Hz and can be changed to achieve the desired trade off between speed and accuracy. The MLX90620 requires a single 3V supply (±0,6V). The customer can choose between 3 operating modes: • Normal. In this mode the device is free running under control of the internal state machine. Depending on the selected refresh rate Fps (Frame per second) the chip is constantly measuring both IR and PTAT and is refreshing the data in the RAM with specified refresh rate; • Step. This mode is foreseen for synchronization with an external micro-controller. The internal state machine is halted. If the command ‘StartMeas’ is received via the I2C bus, a single measurement of all IR and PTAT sensors will be done, then the chip will return in wait state. When in wait state the data in RAM can be read. The MLX90620 is factory calibrated in wide temperature ranges: • -40…85 ˚C for the ambient temperature sensor • -50…300 ˚C for the object temperature. Each pixel of the array measures the average temperature of all objects in its own Field Of View (called nFOV). It is very important for the application designer to understand that the accuracy of the temperature measurement is very sensitive to the thermal equilibrium isothermal conditions (there are no temperature differences across the sensor package). The accuracy of the thermometer can be influenced by temperature differences in the package induced by causes like (among others): Hot electronics behind the sensor, heaters/coolers behind or beside the sensor or by a hot/cold object very close to the sensor that not only heats the sensing element in the thermometer but also the thermometer package. This effect is especially relevant for thermometers with a small FOV as the energy received by the sensor from the object is reduced

LDAEffectSize分析LEfSe详解全

LDAEffectSize 分析LEfSe 详解LDA Effect Size 分析 LEfSe 详解LEfSe 的作⽤在介绍LEfSe的作⽤前，我们先解释⼀个概念——biomarker，维基百科给出的定义是A bio-marker, or biological marker is a measurable indicator of some biological state or condition. Biomarkers are often measured and evaluated to examine normal biological processes, pathogenic processes, or pharmacologic responses to a therapeutic intervention.⽤我们搞数据的⼈能理解的话讲，biomarker就是⾮常强⼒的⽤来分类的特征，它可以是基因、细胞或者物种分类单元等。

⽐如（瞎编的例⼦，不能当真）某个研究团队发现脚⽓会影响中央后回位于Sylvian Fissure附近的区域，从⽽影响⾆头的知觉，于是这个团队打算进⼀步研究脚⽓怎么通过肠脑轴影响⾆头的知觉，他们随机调查了⼀批志愿者，记录了志愿者的⼀些demographic information以及病理信息，并记录了他们肠道菌群的物种分类信息与物种丰度，于是他们有了⼀张数据表：Group有脚⽓⽆脚⽓界.门.纲.⽬.科.属1丰度丰度界.门.纲.⽬.科.属2丰度丰度。

他们想知道哪个属的细菌的丰度在有脚⽓与⽆脚⽓的志愿者之间是存在显著差异的，这个时候就需要LDA Effect Size分析了。

也就是说LDA Effect Size分析的作⽤是发现不同group之间存在显著差异的biomarker。

下⾯我们介绍LDA Effect Size分析的原理。

LEfSe 的原理⾸先我们写出数据，⽤表⽰第个志愿者，，⽤表⽰第个志愿者所在的group， (⽐如讨论有⽆脚⽓时，我们可以⽤表⽰志愿者有脚⽓，⽤表⽰志愿者⽆脚⽓)，⽤表⽰第个志愿者的肠道菌群物种分类信息，⽐如可以表⽰是Bacteroidaceae(拟杆菌科)、可以表⽰Coriobacteriaceae（红蝽杆菌科）；⽤表⽰第个志愿者的demographic information以及病理信息，⽐如可以表⽰性别、表⽰种族、表⽰BMI等。

温度梯度对冷冻靶微管充气过程的影响

温度梯度对冷冻靶微管充气过程的影响余铭铭;周晓松;彭述明;陈绍华;李海容;温成伟;夏立东;尹剑;王伟伟;陈晓华;张晓安【摘要】The research for gas‐filling technology under control was carried out for cryo‐genic target with fill‐tube ,based on temperature gradient between the fuel reservoir and target cell .The effect of temperature gradient on fuel filling processes with different size targets was studied by theoretical calculation and experiment .The results show that the influence on the initial fill pressure difference is slight with respect to different final temperatures for fuel reservoir .In other words ,it is suitable for gas‐filling process for target with any diameter under control by temperature gradient .One can note that the fuel injectivity is more accuracy for the wide range temperature gradient as the tempera‐ture difference reduction for fuel liquefying ,with the target diameter expansion .The sensitivity within ± 3 μm/K for the reservoir at the operation temperature of 75 Kis quite adequate for layer thickness control purposes ,with regardto the target at the diameter of 2 mm and the fuel reservoir at the volume of 1.6 mL .These results could afford an important foundation for highly accuracy of fuel injectivity of cryogenic target .%采用基于燃料室和靶室独立控温的温度梯度法开展了冷冻靶微管可控充气技术研究。

CdSe量子点滤光片尺寸、温度依赖的光学特性

CdSe量子点滤光片尺寸、温度依赖的光学特性王佳彤黄启章高剑峤马越邢笑雪张宇Size and temperature dependence of spectral transmittance for CdSe colloidal quantum dot film filtersWANG Jia-tong, HUANG Qi-zhang, GAO Jian-qiao, MA Yue, XING Xiao-xue, ZHANG Yu引用本文:王佳彤,黄启章,高剑峤,马越,邢笑雪,张宇. CdSe量子点滤光片尺寸、温度依赖的光学特性[J]. 中国光学, 2021, 14(1): 163-169. doi: 10.37188/CO.2020-0198WANG Jia-tong, HUANG Qi-zhang, GAO Jian-qiao, MA Yue, XING Xiao-xue, ZHANG Yu. Size and temperature dependence of spectral transmittance for CdSe colloidal quantum dot film filters[J].Chinese Optics, 2021, 14(1): 163-169. doi: 10.37188/CO.2020-0198在线阅读 View online: https:///10.37188/CO.2020-0198您可能感兴趣的其他文章Articles you may be interested inCdSe/ZnS量子点下转换膜的红、绿、蓝顶发射有机发光器件Top-emitting red, green and blue organic light-emitting devices with CdSe/ZnS quantum dots down-conversion films中国光学. 2019, 12(6): 1431 https:///10.3788/CO.20191206.1431温度变化对金属Ag膜反射镜偏振特性的影响研究Influence of temperature variation on polarization characteristics of silver thin film mirror中国光学. 2018, 11(4): 604 https:///10.3788/CO.20181104.0604量子点液晶显示背光技术Advances and prospects in quantum dots based backlights中国光学. 2017, 10(5): 666 https:///10.3788/CO.20171005.0666基于量子点的荧光型太阳能聚光器Quantum dots based luminescent solar concentrator中国光学. 2017, 10(5): 555 https:///10.3788/CO.20171005.0555PbSe量子点近红外光源的CH4气体检测CH4 detection based on near-infrared luminescence of PbSe quantum dots中国光学. 2018, 11(4): 662 https:///10.3788/CO.20181104.0662荧光碳量子点的制备与生物医学应用研究进展Advances in preparation and biomedical applications of fluorescent carbon quantum dots中国光学. 2018, 11(3): 431 https:///10.3788/CO.20181103.0431文章编号 2095-1531（2021）01-0163-07CdSe 量子点滤光片尺寸、温度依赖的光学特性王佳彤1，黄启章1，高剑峤1，马　越1，邢笑雪2 *，张　宇1 *（1. 吉林大学电子科学与工程学院，吉林长春 130012；2. 长春大学电子信息工程学院，吉林长春 130022）摘要：为了缩小光谱仪体积使之适用于军事卫星等领域，本文将胶体量子点作为滤光材料，研究了CdSe 胶体量子点滤光片的光学特性。

1、下载文档前请自行甄别文档内容的完整性，平台不提供额外的编辑、内容补充、找答案等附加服务。
2、"仅部分预览"的文档,不可在线预览部分如存在完整性等问题,可反馈申请退款(可完整预览的文档不适用该条件!)。
3、如文档侵犯您的权益，请联系客服反馈,我们会尽快为您处理(人工客服工作时间：9:00-18:30)。

NANO EXPRESSEffect of Size-Dependent Thermal Instability on Synthesis of Zn 2SiO 4-SiO x Core–Shell Nanotube Arrays and Their Cathodoluminescence PropertiesChun Li •Yoshio Bando •Benjamin Dierre •Takashi Sekiguchi •Yang Huang •Jing Lin •Dmitri GolbergReceived:21December 2009/Accepted:28January 2010/Published online:10February 2010ÓThe Author(s)2010.This article is published with open access at Abstract Vertically aligned Zn 2SiO 4-SiO x (x \2)core–shell nanotube arrays consisting of Zn 2SiO 4-nanoparticle chains encapsulated into SiO x nanotubes and SiO x -coated Zn 2SiO 4coaxial nanotubes were synthesized via one-step thermal annealing process using ZnO nanowire (ZNW)arrays as templates.The appearance of different nanotube morphologies was due to size-dependent thermal instability and speciﬁc melting of ZNWs.With an increase in ZNW diameter,the formation mechanism changed from decom-position of ‘‘etching’’to Rayleigh instability and then to Kirkendall effect,consequently resulting in polycrys-talline Zn 2SiO 4-SiO x coaxial nanotubes,single-crystalline Zn 2SiO 4-nanoparticle-chain-embedded SiO x nanotubes,and single-crystalline Zn 2SiO 4-SiO x coaxial nanotubes.The difference in spatially resolved optical properties related to a particular morphology was efﬁciently documented by means of cathodoluminescence (CL)spectroscopy using a middle-ultraviolet emission at 310nm from the Zn 2SiO 4phase.Keywords Nano-template ÁCore–shell nanotube ÁCathodoluminescence ÁZinc Silicate ÁRayleigh instability ÁKirkendall effectIntroductionNanotubes made of carbon and diverse inorganic com-pounds have continuously attracted signiﬁcant attention due to their unique fundamental physical properties and many potential applications [1–3].Inorganic nanotube generation strategy can generally be classiﬁed into two categories:ﬁrst,one-step self-organization such as self-rolling and/or Ostwald ripening;second,two-step tem-plate-based fabrication through either scariﬁcation or recently developed solid-state reaction utilizing the Kir-kendall effect [4–6].From a viewpoint of ﬁnal device integration,aligned nanotube arrays are highly desirable.In addition to the direct epitaxial growth of nanotube arrays on lattice-matched substrates,a solid-state reaction under thermal annealing and using readily available nanostruc-ture arrays as templates could be an efﬁcient way to gen-erate novel chemically complex,multiphase nanotube arrays.However,due to a considerable increase in the surface-to-volume ratio with decreasing a nanomaterial size,size-dependent thermal stability and melting of the nanostructured templates during annealing must be care-fully addressed.ZnO nanowire (ZNW)arrays are one of the most com-mon aligned nanostructures owing to their naturally pref-erable epitaxial growth at a relatively low temperature.They have widely been utilized as templates for the syn-thesis of lattice-matched GaN and SiC nanotubes,[7,8]ZnO-related semiconducting heterojunctions [9,10],and ZnO-based ternary compound nanostructures [11].In most cases the latter have been considered to be even more important than pure-phase ZnO ones [4–6].Although solid-state reactions by thermal annealing based on a ZNW template have been used to synthesize ZnO-based ternary compound nanotubes,[4–6,11]no research has beenElectronic supplementary material The online version of this article (doi:10.1007/s11671-010-9556-7)contains supplementary material,which is available to authorized users.C.Li (&)ÁY.Bando ÁY.Huang ÁJ.Lin ÁD.Golberg World Premier International Center for MaterialsNanoarchitectonics (MANA),National Institute for Materials Science,Namiki 1-1,Tsukuba,Ibaraki 305-004,Japan e-mail:LI.Chun@nims.go.jp;whulic@ B.Dierre ÁT.SekiguchiAdvanced Electronic Materials Center,National Institute for Materials Science (NIMS),Namiki 1-1,Tsukuba,Ibaraki 305-0044,JapanNanoscale Res Lett (2010)5:773–780DOI 10.1007/s11671-010-9556-7carried out on the inﬂuence of size-dependent template thermal stability on theﬁnal characteristics of such tubes.In this study,this phenomenon was thoroughly demon-strated for the case of Zn2SiO4.While employing a one-step solid-state reaction and the ZNW array templates, vertically aligned Zn2SiO4-SiO x core–shell nanotube arrays (ZSO),i.e.,Zn2SiO4-nanoparticle chains encapsulated in SiO x nanotubes,SiO x-coated polycrystalline and single-crystalline Zn2SiO4coaxial nanotubes were simultaneously obtained.Furthermore,a cathodoluminescence(CL)study, as a noninvasive and high spatial-resolution characteriza-tion tool,was employed to detect and analyze local struc-tural and optical properties of the nanostructures.The structural and optical differences were effectively identi-ﬁed by transmission electron microscopy(TEM)paired with CL spectroscopy.Finally,the size-dependent thermal instability induced formation mechanism was proposed. Experimental MethodsSynthesis MethodsThe ZNW templates were grown on a ZnO-ﬁlm-coated silicon substrate using a vapor phase transport,as we pre-viously reported[12].The synthesis of ZSO nanotube arrays was carried out in a vacuum tube furnace with an outer diameter of24mm and a length of1,200mm.A0.2g powder mixture of SiO2(Alfa Aldrich,99.9%),activated carbon(as a reductant),and Si with equal molar ratios was placed at the center of the tube.A piece of the ZNW tem-plate was placed downstream of the tube at*14cm away from the source.The tube was sealed and evacuated to a base pressure of*2Pa.The furnace was then heated to 1,100°C at a rate of24°C min-1and kept at this temperature for2h.The local temperature of the substrate was about 1,000°C due to the temperature gradient along the tube furnace.A constantﬂow of high-purity Ar gas was fed into the tube at aﬂow rate of80sccm(standard cubic centi-meters per minute)and a pressure of100Torr throughout entire heating/cooling.After the furnace was naturally cooled to room temperature,the surface color of ZNW templates changed from black to white–gray. Characterization ToolsThe ZNW templates and ZSO samples were characterized by a powder X-ray diffraction(XRD;Rigaku,Ultima III, 40kV/40mA with Cu K a radiation),a scanning electron microscope(SEM;JEOL,JSM-6700F),a high-resolution ﬁeld-emission transmission electron microscope(TEM; JEOL,JEM-3000F),a high-angle annular dark-ﬁeld scan-ning transmission electron microscope detector(HAADF-STEM)(JEOL JEM-3100FEF),and an energy-dispersive X-ray spectrometer(EDX).CL spectra were recorded at room temperature in an ultrahigh-vacuum SEM with a Gemini electron gun(Omicron,Germany)at an accelerating voltage of10kV and a current of1nA.Results and DiscussionVertically aligned ZNW arrays with an underlayer of interconnected nanowalls were grown on a ZnO buffer layer[12].The ZNWs have typical diameters of*30–150nm and lengths of*10l m.Cross-sectional view SEM images(Fig.1a,c)revealed that the ZSO arrays keep the vertical alignment peculiar to the ZNW templates and protrude from the underlayer nanowall networks.The top-view SEM images(Fig.1b,d)clearly indicate that the nanotube diameters are larger than those of ZNW due to SiO x coatings.The originally separated ZNWs stick to each other after thermal annealing.Hollow tube morphologies can be clearly seen from the top-view SEM image of the ZSO sample after scratching off the surface nanotubes (inset of Fig.1d).The XRD spectrum of the ZNW arrays displays only one strong(002)peak due to their high c-axis orientation growth.A strong diffraction peak at34.02o and other four relatively weak peaks in the XRD spectrum (Fig.2)of the ZSO sample can be indexed to a rhombo-hedral Zn2SiO4crystal structure with the lattice constants of a=b=1.394nm and c=0.9309nm(JCPDS Card: 08-0492).The ZnO(002)peak coming from the ZnO buffer layer can also be seen.Detailed crystal structure and compositional analyses of the ZSO samples were carried out by TEM and EDS.Three types of tube morphologies,i.e.,discrete particles encap-sulated tubes,hollow tubes,and tubes only partiallyﬁlled with nanoparticles,were observed(Fig.3a).All the nano-tubes including their tip-ends have a uniform SiO x(x\2, as measured by EDS,see supporting Figure S1)shell layer (the inset TEM image of Fig.3b).High-magniﬁcation and high-resolution TEM observations(HRTEM)and selected area electron diffraction(SAED)patterns reveal that the as-synthesized ZSO samples consist of three types of tubular morphologies depending on their diameters(Fig.3b),as shown in Fig.3c,f,and i,respectively.The nanotubes (*20%)with diameters*50–100nm have usually an amorphous SiO x(*20nm)outer shell and an inner tube (diameter*30nm)made of polycrystalline Zn2SiO4par-ticles(Fig.3c).The polycrystalline nature of the inner tube was conﬁrmed by HRTEM imaging and SAED patterns (Fig.3d,e).The weak lattice fringes in Fig.3e with a d-spacing of 1.62,2.53, 2.63,and 2.83A˚can also be observed.These correspond to a rhombohedral Zn2SiO4 phase(r-Zn2SiO4)and(315),(042),(410),and(113)planesof it.About 40%of nanotubes with diameters of *90–160nm have a unique morphology of single-crystalline Zn 2SiO 4-nanoparticle chains encapsulated into a SiO x nanotube.The SAED pattern in Fig.3h recorded from an embedded rod-shaped nanoparticle can be indexed to the [2¯21]zone of r -Zn 2SiO 4and conﬁrms the single-crystalline character of the particle.The nanoparticles are partiallysurrounded by a thin amorphous SiO x layer (*5nm thick),as seen in Fig.3g.The other *40%nanotubes with outer diameters larger than *150nm were found to be SiO x -coated single-crystalline Zn 2SiO 4nanotubes.The corre-sponding HRTEM image and SAED pattern (Fig.3j,k)verify that an r -Zn 2SiO 4nanotube is a single crystal.Theadjacent lattice fringes separated by 2.83,4.02,and 7.02A˚and labeled in Fig.3d,g,and j correspond to the inter-planar distances of (113),(300),and (110)planes in Zn 2SiO 4,respectively.This fact is consistent with the XRD results shown in Fig.2.The spatial distribution of Zn,Si,and O species was mapped using corresponding EDX K a emissions in the HAADF-STEM mode.Due to a large difference in atomic numbers between Zn and Si,a HAADF-STEM image in Fig.4a clearly reveals the typical structural features of a Zn 2SiO 4-nanoparticle-encapsulated SiO x nanotube,and a SiO x -coated Zn 2SiO 4coaxial nanotube.As shown in the spatially resolved elemental maps in Fig.4,Zn species are only present within encapsulated nanoparticles and the inner shell of the Zn 2SiO 4-SiO x coaxial nanotube.CL spectroscopy combined with SEM imaging has been proven to be a powerful tool for probing the local char-acteristics of low-dimensional materials due to itshighFig.1SEM images of a ,b a ZNW template and c ,d a ZSO sample.The inset top-view SEM image shows the ZSO sample after scratching surface nanotubes out of the substrate.The scale bars in a –d are 1l m and for the inset is 500nm.a ,c cross-sectional view images,b ,d top-viewimagesspatial resolution (a submicrometer range)[13,14].This makes CL a valuable nondestructive tool for studies of inhomogeneities in nanostructures caused by doping or growth condition variations.We attempted to use CL spectroscopy to detect and analyze the Zn 2SiO 4phase in the ZSO samples.The CL spectrum recorded from the ZSO sample cross-section is shown in the inset of Fig.5b.The emission peak at 310nm (4.0eV)with a shoulder at 280nm (4.4eV,see ﬁtting curve of supporting Figure S2),which is well below the band gap 5.4eV of Zn 2SiO 4,is attributed to the radiation recombination from the Zn 2SiO 4phase,considering that the similar emission peak at 300nm (4.13eV)was previously reported for the Zn-Zn 2SiO 4nanocables [15].The 380nm emission peak can be ascribed to the near-band-edge recombination of a crystalline ZnO buffer layer.These two sharp emission peaks may result from the single-crystalline characters of Zn 2SiO 4and ZnO phases,respectively.From the CL spectra recorded in different spots,indicated in Fig.5b,the intensity of peaks at 380nm shows a dramatic decrease from the ZnO buffer layer toward the nanotubes;the opposite trend is in effect for the emission at 310nm.This indicates that all crystalline ZNWs have been totally con-sumed during annealing.The 380nm emission only comes from the ZnO buffer layer.The weak violet emission at 440nm can be rationally assigned to a SiO x phase [16],and the broad emission band at 550nm can be ascribed to the Zn 2SiO 4phase.Note that there is also report on the assignment of 440and 550nm emission bands to the related defect emission of ZnO [17].The corresponding CL emission images recorded for each peak are shown in Fig.5c–f.The 310nm CL emission image (Fig.5c),combined with the 550nm emission image in Fig.5f,shows the inhomogeneous distribution of the Zn 2SiO 4phase,i.e.,stronger intensities for the underlayer nanowalls (which have a larger size compared to the topnanotubeFig.3a Low-magniﬁcation TEM image of nanotubes;b Histogram of the diameter distribution based on 70randomly selected nanotubes.The inset shows tip-end TEM images for the three types of nanotubes.c High-magniﬁcation TEM images,type I:a SiO x -coated polycrystalline Zn 2SiO 4nanotube,f type II:a SiO x nanotube periodically encapsulated with single-crystalline Zn 2SiO 4nanoparticles,and i type III:a SiO x -coated single-crystalline Zn 2SiO 4nanotube;d ,g ,and j corresponding HRTEM images taken from the framed areas marked in (c ),(f ),and (i ),respectively;(e )and (k )corresponding SAED patterns taken from the nanotubes shown in (c )and (i ),respectively.(h )SAED pattern taken from the encapsulated nanoparticle in (f ).The diameter distributions for type I,II,III nanotubes are labeled in different sectors in (b )portions).The uniform 440nm emission image is consis-tent with the TEM observation of a homogeneous SiO x coating.The assignment of 310and 550nm emissions was further conﬁrmed by the emission from the interface between a ZnO buffer layer and a Si substrate,as indicated by arrows in Fig.5c,f.One can expect that the reaction between the ZnO ﬁlm and the surface SiO x layer on a Si substrate can lead to Zn 2SiO 4phase formation during high-temperature annealing [18].The further high-magniﬁcation CL images were recor-ded from representative nanotubes dispersed on a TEM grid.Owing to the high sensitivity and high spatial reso-lution of the CL spectroscopy mapping,the nanotubes with almost same elemental compositions but different mor-phologies can readily be indentiﬁed at a large scale.As shown in Fig.6c,strong 310nm CL emission comes from both the embedded Zn 2SiO 4nanopaticles (type II)and Zn 2SiO 4nanotubes of larger size (type III),while smaller SiO x -coated Zn 2SiO 4nanotubes (type I)only show hardly detectable emission.These results are in a good agreement with TEM characterization.Also,they further conﬁrm the above-mentioned peak assignment.It will be of interest to investigate the formation mechanism during simultaneous generation of different tube structures,especially regarding the Zn 2SiO 4-nano-particle chains,for the ﬁrst time observed here in a ternary compound nanomaterial.We believe that nanoscale ther-modynamics of ZNW is an important factor.It has been experimentally demonstrated that the melting point of Cu,Zn,and Sn nanowires will signiﬁcantly decrease with a decrease in wire diameter [19–21].Based on molecular dynamics simulations,the same diameter dependency has also been found for GaN nanowires [22].In addition,under heating in air,ZnO nanorods start to melt at 750°C,[23]much lower temperature than the melting point of a bulk form.Therefore,it is reasonable to assume that the melting point of ZNW would also decrease with diameter decreasing.Based on the above-mentioned assumption and taking the results of XRD,TEM,EDS,and CL measurements,we conclude that the factors responsible for the different tube morphologies are the size-dependent melting behavior of ZnO and the competition between surface and volume diffusions.Since the Si atom has nearly the same radius as the Zn atom and the bonding energy of the Si–O bond (185kJ/mol)is about two times higher than that of the Zn–O bond (92kJ/mol),it is easy for ZnO to diffuse into SiO and to form a Zn 2SiO 4phase,as reported by Wang and Zhou [11,25].It is also noted that the ZNW template was totally decomposed in a control experiment under the same heating process but without introducing a source powder.The formation mechanism here is illustrated in Fig.7,and the three predominated effects decomposition of ‘‘etch-ing’’,the Rayleigh instability,and the Kirkendall solid-state reaction are highlighted for the nanotubes of type I,II,and III,respectively.During the temperature increase,the Si–O vapor (sublimated from the source)can quickly be transported and deposited on the nanowire surfaces,pre-venting their decomposition.Rapid heating quickly drives the system to high temperature.The ZNWs withsmallerFig.4a HAADF-STEM image of two typical Zn 2SiO 4nanotube morphologies,a SiO x -shelled Zn 2SiO 4nanochain-encapsulated nanotube (left )and a SiO x -coated Zn 2SiO 4nanotube (right );b ,c ,and d the corresponding spatially resolved elemental EDS maps of Zn,Si,and Odiameters (\60nm,estimated from the above TEM sta-tistical analysis),which possess lower melting tempera-tures,prefer to decompose into Zn vapor and O 2rather than to react with Si forming Zn 2SiO 4.The ZnO deﬁciency and SiO x richness at the interface lead to the formation of a thin polycrystalline Zn 2SiO 4layer on the inner surface oftheFig.5a Cross-sectional view SEM image of a ZSO sample and corresponding room-temperature CL spectrum in (b );c ,d ,e ,and f CL images of the array under the 310,380,440,and 550nm emissions,respectivelyFig.6a ,b ,and c TEM and SEM images,and corresponding 310nm CL emission image of representative nanotubes dispersed on a TEM grid.Three nanotube types are labeled as I,II,and III,and indicatedby arrows ,respectively.Note that type I nanotubes consist of two nanotubes of small diameter.The scale bars are 200nmSiO x shell.The ZNWs with medium diameters (*60–120nm)prefer melting rather than decomposition.In accord with the Rayleigh instability effect,[19,24]the ZnO liquid cylinder of radius r is unstable to radius per-turbations whose wavelength l exceeds the circumference of the cylinder.The cylinder thus decreases its total energy and breaks into nanoparticles under surface tension.As the ZnO liquid starts to shrink,the Si/O surface diffusion will rapidly increase around the curved portions,since the surface diffusion of atoms is proportional to the gradient of the wire surface curvature at a certain temperature [26].Subsequent volume diffusion at the interface and sub-sequent reactions lead to single-crystalline Zn 2SiO 4-nano-particle chains encapsulated into SiO x nanotubes.The ZNWs with larger diameters ([120nm)will be in a ther-mally unstable solid state.Unequal diffusion rates for the outward Zn/O and inward Si/O,known as the Kirkendall diffusion,[4–6]produce a hollow interior along the struc-ture length.The SiO x coating is thicker than necessary for complete consuming of the interior ZnO to form a new phase.As a result,SiO x -coated single-crystalline Zn 2SiO 4coaxial nanotubes are formed.Note that due to the diam-eter nonuniformity of an individual wire,nanotubes with partial nanoparticle encapsulation may also form,as shown in Fig.3a.It is worth mentioning that by using plasma-enhanced chemical vapor deposition (PECVD)of a thin amorphous Si ﬁlm (*10nm)on the top half of a ZNW template followed by vacuum annealing,Zhou et al .[11]fabricated vertically aligned Zn 2SiO 4nanotube/ZnO nanowire het-erojunction arrays.Our results undoubtedly demonstrate that prior to ZnO template utilization for reliable makingternary compound nanotubes,problems associated with their thermal instability must be seriously taken into account and solved.The good news is that due to many diameter-controllable synthesis methods of ZNWs reported to date,ﬁnding templates with uniform diameters for the subsequent unique one speciﬁc ternary nanotube syntheses looks highly plausible.Furthermore,the Rayleigh insta-bility applied to ZNW provides a structuring technique that produces long chains of ZnO-based compound semicon-ducting nanospheres which may ﬁnd potential applications in nanoscale photonic devices [27].In conclusions,vertically aligned Zn 2SiO 4-nanoparticle chains encapsulated into SiO x nanotubes,SiO x -coated polycrystalline and single-crystalline Zn 2SiO 4coaxial nanotubes were simultaneously fabricated by a one-step solid-state reaction using ZNW array templates.It is found that the nanotubes of different morphologies can be easily identiﬁed by CL spectroscopy mapping technique.The apparent size-dependent nonuniform nanotube generation was due to the strong size-dependent thermodynamic behavior of ZNWs.With diameter increasing of starting ZNW,the formation mechanism changes from decompo-sition of ‘‘etching’’,to Rayleigh instability,and then to Kirkendall effect,resulting in polycrystalline Zn 2SiO 4-SiO x coaxial nanotubes,single-crystalline Zn 2SiO 4-nanoparticle-chain-embedded SiO x nanotubes,and single-crystalline Zn 2SiO 4-SiO x coaxial nanotubes,respectively.Acknowledgments This work was supported by the World Premier International Center for Materials Nanoarchitectonics (MANA)Pro-ject tenable at the National Institute for Materials Science (NIMS),Tsukuba,Japan.The authors thank Drs.A.Nukui,M.Mitome,and Mr.K.Kurashima for the continuous technical support and kindhelp.Fig.7Schematics of theformation mechanism for three discovered types of tubular structuresOpen Access This article is distributed under the terms of the Creative Commons Attribution Noncommercial License which per-mits any noncommercial use,distribution,and reproduction in any medium,provided the original author(s)and source are credited. References1.M.S.Dresselhaus,Y.M.Lin,O.Rabin,A.Jorio,A.G.S.Filho,M.A.Pimenta,R.Saito,G.G.Samsonidze,G.Dresselhuas, Mater.Sci.Eng.C23,129(2003)2.Y.Xiong,B.T.Mayers,Y.Xia,mun.5013(2005)3.H.C.Zeng,J.Mater.Chem.16,649(2006)4.H.J.Fan,U.Go¨sele,M.Zacharias,Small3,1660(2007)5.H.J.Fan,Y.Yang,M.Zacharias,J.Mater.Chem.19,885(2009)6.C.Yan,D.Xue,Adv.Mater20,1055(2008)7.J.Goldberger,R.R.He,Y.F.Zhang,S.W.Lee,H.Q.Yan,H.J.Choi,P.D.Yang,Nature422,599(2003)8.J.Zhou,J.Liu,J.Yang,o,C.S.Gao,R.Tummala,N.S.Xu,Z.L.Wang,Small2,1344(2006)9.C.Yan,D.Xue,J.Phys.Chem.B110,25850(2006)10.K.Wang,J.Chen,W.Zhou,Y.Zhang,Y.Yan,J.Pern,A.Mascarenhas,Adv.Mater20,3248(2008)11.J.Zhou,J.Liu,X.D.Wang,J.H.Song,R.Tummala,N.S.Xu,Z.L.Wang,Small4,622(2007)12.C.Li,G.J.Fang,J.Li,L.Ai,B.Z.Dong,X.Z.Zhao,J.Phys.Chem.C112,990(2008)13.S.I.Maximenko,L.Mazeina,Y.N.Picard,J.A.Freitas Jr.,V.M.Bermudez,S.M.Prokes,Nano Lett.9,3245(2009)14.X.L.Yuan,B.Dierre,J.B.Wang,B.P.Zhang,T.Sekiguchi,J.Nanosci.Nanotech.7,3323(2007)15.X.Feng,X.Yuan,T.Sekiguchi,W.Lin,J.Y.Kang,J.Phys.Chem.B.109,15786(2005)16.D.P.Yu,Q.L.Hang,Y.Ding,H.Z.Zhang,Z.G.Bai,J.J.Wang,Y.H.Zou,W.Qian,G.C.Xiong,S.Q.Feng,Appl.Phys.Lett.73, 3076(1998)17.A.B.Djurisic,Y.H.Leung,Small2,944(2006)18.X.L.Xu,P.Wang,Z.M.Qi,H.Ming,J.Xu,H.T.Liu,C.S.Shi,G.Lu,W.K.Ge,J.Phys.:Condens.Matter.15,L607(2003)19.M.E.T.Molares,A.G.Balogh,T.W.Cornelius,R.Neumann,C.Trautmann,Appl.Phys.Lett.85,5337(2004)20.X.W.Wang,G.T.Fei,K.Zheng,Z.Jin,L.D.Zhang,Appl.Phys.Lett.88,173114(2006)21.H.S.Shin,J.Yu,J.Y.Song,Appl.Phys.Lett.91,173106(2007)22.W.H.Moon,H.J.Kim,C.H.Choi,Scr.Mater.56,345(2007)23.X.Su,Z.Zhang,M.Zhu,Appl.Phys.Lett.88,061913(2006)24.Y.Qin,S.M.Lee,A.Pan,U.Go¨sele,M.Knez,Nano Lett.8,114(2008)25.H.Q.Wang,G.Z.Wang,L.C.Jia,C.J.Tang,G.H.Li,J.Phys.Chem.C.111,14307(2007)26.F.A.Nichols,W.W.Mullins,Trans.Metall.Soc.AIME233,1840(1965)27.Y.Zhu,H.I.Elim,Y.-L.Foo,T.Yu,Y.Liu,W.Ji,J.-Y.Lee,Z.A.Shen,T.S.Wee,J.T.L.Thong,C.H.Sow,Adv.Mater.5, 587(2006)。