In$uence of annealing treatment on Laves phase compound containing a V-based BCC solid solution p

1 表面活性剂研究表明，选用有机改性硅共聚物L-520作为表面活性剂处理含α结晶相PVDF 粉末，再经溶剂抽提处理，使得PVDF 粉末表面含有<1.8 wt%的表面活性剂[1]。

处理后的粉末在经过热压缓慢升温至熔点以上，然后将淬火后的样品浸泡于甲苯中除去表面活性剂。

预浸表面活性剂的PVDF 粉末在DSC 慢速升温测试过程中，在低温处出现对应α结晶相的熔融吸热峰；在高温处出现了第二个熔融吸热峰，且随着升温速率的下降该峰越发明显，该峰对应的是γ结晶相的熔融。

研究表明：有机改性硅共聚物表面活性剂的引入可诱导PVDF 中大量γ结晶相的形成。

形成机理为表面活性剂诱导α结晶相向γ结晶相发生转变，表面活性剂干扰在熔融过程中α结晶相内分子的扩散，使分子链以松散的片晶形式存在，转变为更稳定的γ结晶相的分子链构象。

2 离子液体离子液体具有低熔点、低蒸气压、高的化学和热稳定性、高的离子导率和宽的电化学势能范围等特点，广泛用于润滑剂、绿色溶剂、电池电解质等领域。

离子液体可用作无机填料与聚合物树脂基体的增溶剂，增强两者的界面相容性，改善无机填料的分散性，尤其是纳米填料如碳纳米管和蒙脱土的分散。

离子液体也可以促进PVDF 中极性结晶相的形成。

离子液体修饰改性的碳纳米管可改变PVDF 的结晶行为。

离子液体的引入不仅有利于碳纳米管的分散，也可诱导非极性晶型向极性晶型的转变，同时离子液体与PVDF 分子链间会形成特定的相互作用。

利用1-丁基-3-甲基咪唑六氟磷酸盐([BMIM][PF 6])离子液体改性PVDF ，将PVDF 与不同比例的[BMIM][PF 6]在190 ℃转速为50 r/min 条件下混合5 min ，将混合后的样品在200 ℃热压成300 μm 厚的薄膜[2]。

红外测试结果表明：纯PVDF 样品中0 引言聚偏氟乙烯(PVDF)是一种电活性的多晶型氟聚合物，分子主链中-CH 2-和-CF 2-链节交替排列。

Effects of annealing treatment on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIPLijuan Zhou a ,⇑,Hongbo Li b ,Yongting Zheng ba School of Materials Science and Engineering,Shandong University of Technology,ZiBo 255049,PR China bCenter for Composite Materials and Structures,Harbin Institute of Technology,Harbin 150001,PR Chinaa r t i c l e i n f o Article history:Received 14July 2012Received in revised form 10November 2012Accepted 12November 2012Available online 20November 2012Keywords:AlN–SiC–TiB 2ceramic Annealing treatment MicrostructureMechanical propertiesa b s t r a c tAnnealing treatments were carried out at different temperatures (600–1200°C)for various holding times (2–10h)to evaluate the microstructural and mechanical properties changes of the AlN–SiC–TiB 2ceramic composites prepared by SHS–HIP.The experimental results show that the annealing treatment is bene-ficial to improve the mechanical properties.In the meantime,the desolution of SiC-rich solid solution and precipitation of fine TiB 2grains occurred.Due to the precipitation strengthening by the SiC-rich solid solution and the grain boundary strengthening by fine TiB 2particles,the improvement of the mechanical properties was obvious with higher annealing temperatures and longer holding times.The highest HRA,bending strength and fracture toughness were 94.1,517.5MPa and 5.79MPa m 1/2,respectively,after the annealing treatment at 1200°C for 10h.Ó2012Elsevier B.V.All rights reserved.1.IntroductionAlN–SiC solid solution has a wide range of engineering appli-cations due to their excellent mechanical properties at high tem-peratures and has been investigated considerably in the past decades years [1,2].Many preparation methods,such as hot pressing [3],reaction synthesis [4],spark plasma synthesis sinter-ing [5]and combustion synthesis [6],were adopted to fabricate AlN–SiC ceramics.TiB 2is another important material for high temperature applications because of its high melting point,mechanical properties and electrical conductivity,and relatively low coefficient of thermal expansion.So AlN–SiC–TiB 2ceramics have wide potential applications in the field of electrical conduc-tivity at high temperatures.It is known that combustion synthesis (self-propagating high-temperature synthesis,SHS)is a novel method for preparation of AlN–SiC–TiB 2ceramics with many advantages,such as high efficiency and energy saving.However,combustion synthesis is a high temperature and velocity reaction process,which often result in large amount of non-equilibrium phases and high residual stress,and this deteriorates the proper-ties of the combustion synthesized products.So annealing treat-ment at elevated temperatures (800–1200°C)has been introduced to reduce the residual thermal stress and improve the mechanical properties [7,8].In this paper,the effects of annealing treatment on properties of the AlN–SiC–TiB 2ceramics prepared by self-propagating high tem-perature synthesis and hot isostatic pressing (SHS–HIP)were stud-ied.The influences of the annealing temperature and holding time on the composition,microstructure and mechanical properties were discussed in detail.2.Experimental procedureThe AlN–SiC–TiB 2ceramics were prepared by SHS–HIP reported in reference [9].The reaction for preparing AlN–SiC–TiB 2ceramics is as follows:Al þSiC þTiB 2þN 2!AlN þSiC þTiB 2ð1ÞThe samples with Al 35wt%–SiC 35wt%–TiB 230wt%were chosen to carry out the annealing treatment tests.The annealing treatments were carried out on the box-type electric furnace (SXK-8-16,Longjiang electrical furnace works,Harbin,China)in vacuum atmosphere.The specimens were enclosed and vacuumed in quartz glass tube.During the annealing treatment test,four different temperatures (600,800,1000and 1200°C)and five different holding times (2,4,6,8and 10h)were adopted to study the effects of annealing temperature and holding time on the microstructure and mechanical properties of the AlN–SiC–TiB 2ceramics.The composition of the annealing treated samples was analyzed by X-ray dif-fraction (D/max-rB,Rigaku,Japan).The microstructures were observed by scanning electron microscope (SEM,HITACHI S-4700),energy dispersive X-ray spectrometer,and transmission electron microscopy (TEM,Philips CM12/STEM,Holland).The specimen of 3mm Â4mm Â36mm and 2mm Â4mm Â22mm in dimension for bending strength and fracture toughness tests were sliced respectively.The three-point bending strength and fracture toughness were tested on electronic uni-versal test machine with crosshead speed of 0.5mm/min (Instron5569,USA).The Rockwell hardness test was carried out with a ½00ball and 60kg load applied for 30s on the Rockwell hardness test machine (HR-150A,China).0925-8388/$-see front matter Ó2012Elsevier B.V.All rights reserved./10.1016/j.jallcom.2012.11.080Corresponding author.Tel./fax:+865332781357.E-mail address:zhoulijuan@ (L.Zhou).3.Results and discussion3.1.MicrostructureFig.1shows the XRD patterns of theafter the annealing treatment for10h atAs thefigure indicated,oxidationnot occurred and the main phases wereTiB2after the annealing treatment.temperature and nitrogen pressure duringcontent of hexagonal boron nitride(h-BN)ucts.With the increasing of the annealingtion intensity of h-BN phase decreased,TiN phase appeared at the annealing10h,and slightly increases as the annealingto1200°C.The formation of the smallfavorable to the improvement of thermalfracture toughness because of theirand micro-cracking interactionkind of structural ceramic with excellenttion of h-BN phase during preparation anding treatment will not decrease the properties of the products,onthe contrary,it is helpful to the improvement of mechanical properties.Fig.2shows the XRD patterns of AlN–SiC–TiB2ceramics after annealing treatment at1200°C for different holding times.It can be found that the composition of the annealing treated ceramics was almost unchanged at1200°C for8h.When the holding time increased to10h,a small quantity of TiN was formed in the prod-uct,which was consistent with the XRD results shown in Fig.1. Furthermore,the X-ray diffraction peaks of2H–AlN–SiC phase were becoming narrower and sharper with the increasing of the annealing holding time.Meantime,the XRD peaks of the2H–AlN–SiC phase appeared to shift to larger angles,which indicated 3.2.Strengthening and toughening mechanismFig.3shows the BSE images of AlN–SiC–TiB2samples annealed for10h at600and1200°C.AlN–SiC matrix shows gray color and the dispersed white particles are TiB2phase.Fig.3a shows the sam-ple surface annealed at600°C for10h.White columnar TiB2were dispersed inhomogenously in the matrix,and also the sample sur-face presented many pores with different sizes and shapes.When the annealed temperature was1200°C,the dispersion of TiB2 was relatively homogenous andfiner,and the porosity of the sur-face was relatively decreased,as shown in Fig.3b.It is noted that some light gray phase exists between gray ma-trix and white particles,as shown in Fig.4.Line scanning for the light gray region was conducted to determine the phase composi-tion,as shown in Fig.4.The results show that the light gray phase contained large amount of Si element and tiny content of Al,and no Ti element was detected,which can be indicated that the light gray phase was the SiC-rich phase precipitated during the annealing treatment.After the precipitation of SiC,the AlN-rich and SiC-rich solid solutions were formed correspondingly,which was consis-tent with the phase diagram analysis[10].Phase separation leads to grain size refinement,and then results in the improvement of the yield strength.According to Hall–Petch equation,the relationship between the yield strength and grain size is described as follows[11]:ry¼r0þkdÀ1=2ð2Þwhere r y is the yield strength,r0is a materials constant for the starting stress for dislocation movement,k is the strengthening coefficient,and d is the average grain diameter.As the equation indicated,the grain size refinement is in favor of the improvement of yield strength of the materials.Fig.5shows the TEM image and SAED pattern of AlN–SiC–TiB2 ceramic annealed at1200°C for10h.Light gray matrix was AlN–SiC and the black particles embedded in the matrix were TiB2 phase,as shown in Fig.5a.The SAED pattern of the TiB2particles and the calibration were shown in Fig.5b and c,respectively.The average size of these TiB2particles was much less than those in reactant,and also,these particles were completely embedded in the AlN–SiC matrix,which indicated that thefine TiB2particles were precipitated during the annealing treatment.Since combustion synthesis is a high temperature and velocity reaction process,many dislocations,grain boundaries andFig.1.XRD patterns of AlN–SiC–TiB2ceramics annealed at different temperatures for10h.Fig.2.XRD patterns of AlN–SiC–TiB2ceramics annealed at1200°C for different holding times.500L.metastructures can be formed in the products.In addition,boron and titanium could dissolve in the AlN–SiC matrix or in the grain boundaries because of high reaction temperature.What is more,the standard enthalpy of formation (D H )of TiB 2is relatively low (D H =À293kJ mol À1),and the fine TiB 2particles can be precipi-tated from the supersaturated and non-equilibrium AlN–SiC ma-trix during the high temperature annealing treatment.Precipitating of these fine TiB 2particles decreased the lattice distortion,which also released the inner stress and improved the order degree of the lattice.Meantime,composition gradient formed by the precipitation between the matrix and TiB 2phases enhanced the interface bonding strength.In summary,the improvement of the strength and toughness of the AlN–SiC–TiB 2ceramics can be attributed to the particle dis-persion strengthening effect of the precipitated SiC-rich solid solution.And also the fine TiB 2particles precipitated uniformly at the phase interface enhanced the strength of the grain boundaries.nano-particles in AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)TEM image of AlN–SiC–TiB 2composite;(b)SAED of light gray phase in the AlN–SiC–TiB 2ceramic annealed at 1200°C for 10h:(a)SEM image of the surface;(b)line Fig.3.Backscattered electron images of AlN–SiC–TiB 2samples annealed at (a)600°C,10h;(b)1200°C,10Fig. 6.Typical SEM fractograph in annealing treated AlN–SiC–TiB2ceramic composite.Fig.7.Rockwell hardness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.8.Bending strength of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.Fig.9.Fracture toughness of AlN–SiC–TiB2ceramics annealed at different temper-atures for various holding times.annealing temperature reached800and1200°C.The highest HRA,bending strength and fracture toughness were94.1, 517.5MPa and 5.79MPa m1/2,respectively,after the annealing treatment at1200°C for10h.4.ConclusionsThe effects of annealing treatments on the microstructure and mechanical properties of AlN–SiC–TiB2ceramics prepared by SHS–HIP were studied.The annealing treatment can promote the precipitation SiC-rich solid solution in AlN–SiC matrix andfine TiB2particles on the grain boundaries,which improve the mechan-ical properties obviously.With the annealing temperature increas-ing from600to1200°C and holding times increasing from2to 10h,the strengthening effect was more evident.The highest HRA,bending strength and fracture toughness of AlN–SiC–TiB2 ceramics were94.1,517.5MPa and5.79MPa m1/2,respectively, after the annealing treatment at1200°C for10h.AcknowledgmentsThis work has beenfinancially supported by‘‘A Project of Shandong Province Higher Educational Science and Technology Program(J09LD01)’’and‘‘Shandong Province Natural Science Foundation(ZR2010EM045)’’.References[1]L.J.Zhou,Y.T.Zheng,S.Y.Du,H.B.Li,Mater.Sci.Forum546–549(2007)1505–1508.[2]vrenko,J.Desmaison,A.D.Panasyuk,M.Desmaison-Brut,E.Fenard,J.Eur.Ceram.Soc.25(2005)1781–1787.[3]K.Strecker,M.J.Hoffmann,J.Eur.Ceram.Soc.25(2005)801–807.[4]X.M.Yue,G.J.Zhang,Y.M.Wang,J.Eur.Ceram.Soc.19(1999)293–298.[5]M.Hotta,J.Hojo,J.Eur.Ceram.Soc.30(2010)2117–2122.[6]L.Mei,J.T.Li,Acta Mater.56(2008)3543–3549.[7]A.Varma,A.S.Rogachev,A.S.Mukasyan,S.Hwang,Adv.Chem.Eng.24(1998)79–226.[8]D.Sciti,S.Guicciardi,A.Bellosi,J.Eur.Ceram.Soc.21(2001)621–632.[9]L.J.Zhou,Y.T.Zheng,S.Y.Du,Key Eng.Mater.353–358(2007)1517–1520.[10]A.Zangvil,R.Ruh,J.Am.Ceram.Soc.71(1988)884–890.[11]H.Conrad,J.Narayan,Scripta Mater.42(2000)1025–1030.L.Zhou et al./Journal of Alloys and Compounds552(2013)499–503503。

Influence of interface mobility on the evolution of austenite–martensite grain assemblies during annealingM.J.Santofimia a,b,*,J.G.Speer c ,A.J.Clarke d ,L.Zhao a,b ,J.Sietsma baMaterials Innovation Institute (M2i),Mekelweg 2,2628CD Delft,The NetherlandsbDepartment of Materials Science and Engineering,Delft University of Technology,Mekelweg 2,2628CD Delft,The NetherlandscAdvanced Steel Processing and Products Research Center,Colorado School of Mines,Golden,CO 80401,USAdMaterials Science and Technology Division,Mail Stop G770,Los Alamos National Laboratory,Los Alamos,NM 87545,USAReceived 9March 2009;accepted 17June 2009Available online 13July 2009AbstractThe quenching and partitioning (Q&P)process is a new heat treatment for the creation of advanced high-strength steels.This treatment consists of an initial partial or full austenitization,followed by a quench to form a controlled amount of martensite and an annealing step to partition carbon atoms from the martensite to the austenite.In this work,the microstructural evolution during annealing of martensite–austenite grain assemblies has been analyzed by means of a modeling approach that considers the influence of martensite–austenite interface migration on the kinetics of carbon partitioning.Carbide precipitation is precluded in the model,and three different assumptions about interface mobility are considered,ranging from a completely immobile interface to the relatively high mobility of an incoherent ferrite–austenite interface.Simulations indicate that different interface mobilities lead to profound differ-ences in the evolution of microstructure that is predicted during annealing.Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.Keywords:Quenching;Annealing;Steels;Diffusion;Thermodynamics1.IntroductionCurrent demands on fuel consumption and safety have led the automotive industry to search for new advanced steels with enhanced strength and ductility.One of the ideas being explored is the development of low-carbon steels with a microstructure consisting of martensite and a consider-able fraction of retained austenite.This combination of phases can lead to a high strength,because of the presence of martensite,and considerable formability.Although these microstructures have been observed in the past in quenched martensitic steels,the amount and stability of the retainedaustenite found were usually low [1,2].In addition,during subsequent tempering,reduction of carbon in the martens-ite occurred via carbide precipitation,whereas austenite was usually decomposed into ferrite and carbides.Knowledge of the effect of some elements,e.g.silicon and aluminum,in inhibiting cementite precipitation has opened the possibility for obtaining carbon-enriched aus-tenite by partitioning of carbon from supersaturated mar-tensite.The recently proposed [3,4]‘‘quenching and partitioning ”(Q&P)process makes use of this idea.This new heat treatment consists of a partial martensite trans-formation (quenching step)from a fully or partially austen-itized condition,followed by an annealing treatment (partitioning step)at the same or higher temperature to promote carbon partitioning from the supersaturated martensite to the austenite.During the partitioning step it is intended that the austenite be enriched with carbon,thus allowing its stabilization at room temperature.The1359-6454/$36.00Ó2009Acta Materialia Inc.Published by Elsevier Ltd.All rights reserved.doi:10.1016/j.actamat.2009.06.024*Corresponding author.Present address:Instituto Madrilen ˜o de Estu-dios Avanzados en Materiales (IMDEA-Materiales),ETS de Ingenierı´a de Caminos 28040Madrid,Spain.Tel.:+34915493422;fax:+34915503047.E-mail address:mariajesus.santofimia@ (M.J.Santofimia)./locate/actamatAvailable online at Acta Materialia 57(2009)4548–4557resulting microstructure after the whole thermal cycle con-sists of ferrite(in the case of an initial partial austenitiza-tion),martensite and retained austenite.In this paper,the partitioning step will be further referred to as annealing, to avoid confusion with the process of carbon migration (partitioning).From the above,it is clear that the essential mechanism of the Q&P process is the transfer of carbon from the supersaturated martensite to the austenite.Given that this mechanism of carbon partitioning was not considered in detail in the past,the conditions under which it takes place are now under debate.Some authors[5–8]have postulated a‘‘constrained carbon equilibrium”(CCE)condition gov-erning the carbonflux from the martensite to the austenite. The CCE takes into account that iron and substitutional atoms are less mobile at temperatures at which carbon dif-fusion takes place and that the martensite–austenite inter-face can be assumed immobile or stationary.Therefore, only carbon equilibrates its chemical potential.There are some experimental observations that question whether the martensite–austenite interface remains station-ary during annealing.Zhong et al.[9]have reported the apparent migration of these interfaces in a low-carbon steel after annealing at480°C.Although the direction of migra-tion has not been established,this observation indicates the importance of understanding the transfer of iron atoms in relation to the partitioning of carbon.Another interesting observation that contradicts the simplifying assumption of a stationary interface is the reported expansion of the material during the annealing(partitioning step)observed by dilatometry[10],probably indicating changes in the fractions of phases.However,a definitive explanation of the causes of this expansion is not yet available,since it is unclear if the expansion is a result of the continued growth of already present athermal martensite,the nucle-ation of new isothermal martensite or bainite reaction [11].Another interesting unexplained feature is the pres-ence of two peaks in the representation of retained austen-ite fraction vs.annealing time,which has been attributed to the competition between carbon partitioning and carbide precipitation[12].Given these contradictions,Speer et al.[13]recently con-sidered the implications of iron atom movement on the evolution of the martensite–austenite interface during annealing.According to that work[13],‘‘the difference in iron potential between the ferrite and the austenite creates a driving force for iron to move from one structure to the other,which is accomplished via migration of the existing interface,assuming that nucleation of new crystals does not occur”.Under these considerations,Santofimia et al.[14]quantitatively analyzed the motion of the martensite–austenite interface in a model based on thermodynamics and diffusion,assuming the same chemical potential of car-bon in martensite and austenite at the interface and allow-ing motion of the phase interface when a free-energy difference occurs.Simulations corresponding to a particu-lar realistic microstructure were presented,showing a sig-nificant bidirectional movement of the martensite–austenite interface.These calculations were made assuming an activation energy for the migration of iron atoms corre-sponding to data on austenite to ferrite transformation in steels(140kJ molÀ1)[15,16],which implies the assumption of an incoherent martensite–austenite interface.In princi-ple,the use of this activation energy could seem inconsis-tent with the well-known semicoherent character of the martensite–austenite interface created during martensite formation[17].However,a treatment of annealing at the transformation temperature or at higher temperatures(that can be identified as the annealing or partitioning tempera-ture of the Q&P process,typically between250and500°C) can affect the character and thus the mobility of the mar-tensite–austenite interface.In any case,there is a significant lack of studies in this area.Therefore,the theoretical anal-ysis of phases and carbon behavior during annealing of martensite–austenite microstructures assuming different interface characters is an alternative way to study mecha-nisms and provide an insight into the above-mentioned experimental observations.In this work,microstructural evolution during annealing of martensite–austenite grain assemblies has been analyzed by means of a modeling approach that considers the influ-ence of the coupling between martensite–austenite interface migration and the kinetics of carbon partitioning.Assum-ing that the character of the martensite–austenite interface influences the activation energy for iron migration from one phase to the other,three different activation energies are considered in this study:(i)‘‘infinite”(i.e.immobile interface)which corresponds to CCE conditions;(ii) 140kJ molÀ1from data on the austenite to ferrite transfor-mation involving incoherent interfaces[15,16];and(iii)a higher value(180kJ molÀ1)which represents an estimated value for semicoherent interfaces.Carbon profiles and vol-ume fraction of phases predicted as a function of the quenching temperature,annealing temperature and mar-tensite–austenite interface mobility are analyzed.For sim-plicity,carbide precipitation is assumed to be suppressed completely.2.ModelThe interaction between carbon partitioning and inter-face migration is analyzed using the model presented by Santofimia et al.[14].Some aspects of this model are reviewed here for a proper understanding of the analysis presented in the following sections.For modeling purposes,martensite is considered to have a body-centered cubic(bcc)structure supersaturated in carbon,whereas austenite is a face-centered cubic(fcc) phase.The model considers the same chemical potential of carbon in bcc and in fcc at the bcc–fcc interface because of the high atomic mobility of interstitial carbon,which is one of the CCE conditions.This condition is expressed in terms of carbon concentration by Eq.(1)presented in Ref.[14].M.J.Santofimia et al./Acta Materialia57(2009)4548–45574549The motion of interfaces in a microstructure is a result of the repositioning of atoms from lattice positions in one grain to projected lattice positions in a neighboring grain.At a given temperature,the equilibrium concentra-tions of carbon in fcc,x fcc-eqC ,and bcc,x bcc-eqC,are given bythe metastable equilibrium phase diagram,excluding car-bide formation.If the carbon concentrations at the inter-face are different from the equilibrium values,the phases will experience a driving pressure,D G,for a phase transfor-mation towards the equilibrium phase composition.This local driving pressure is experienced at the interface and results in an interface velocity,v,which is proportional to the driving pressure according to:v¼M D G;ð1Þwhere M is the interface mobility.In this work,the driving pressure is considered proportional to the difference be-tween the equilibrium concentration of carbon in fcc and the interface carbon concentration in fcc,for which the proportionality factor is calculated from Thermo-Calc[18].The driving pressure can be positive or negative, depending on the relative difference between the equilib-rium carbon content of the austenite and the actual carbon concentration in austenite at the interface.The relationship between the carbon content in the austenite at the fcc–bccinterface,x fcc–bccC ,and the interface migration behavior,according to the present model,is schematically repre-sented in Fig.1.If the interface is enriched in carbon rela-tive to equilibrium,then the chemical potential of iron is higher in martensite than in austenite and the driving pres-sure for the movement of the interface promotes interface migration from the austenite to the martensite(Fig.1a), whereas the interface would be promoted to move in the opposite direction if the interface is depleted in carbon relative to equilibrium(Fig.1b).The interface mobility,which is temperature dependent, can be expressed as a product of a pre-exponential factor and an exponential term:M¼M0expÀQM RT;ð2Þwhere Q M is the activation energy for iron atom motion at the interface.The pre-exponential factor,M0,can be expressed as[19]:M0¼d4m Dk B T;ð3Þwhere d is the average atomic spacing in the two phases separated by the interface in question,m D is the Debye fre-quency and k B is the Boltzmann constant.The value of d has been estimated to be2.55A˚for a martensite–austenite interface[20].The diffusion of carbon in martensite and austenite is modeled by solving Fick’s second law using a standard finite-difference method[21].Diffusion coefficients are cal-culated referring to the carbon content in martensite[22] and austenite[23].3.Simulation conditionsIn order to study the influence of the martensite–austen-ite interface character on the interaction between carbon partitioning and iron migration during annealing,it is assumed that modifications to the interface character lead to different values of the activation energy for iron migra-tion.This is a reasonable qualitative approximation,since the mobility of a martensite–austenite interface during annealing is related to the coherency of the interface.For example,iron atoms migrate more easily in incoherent interfaces.Although it is now not possible to exactly relate the value of the activation energy to the specific character of the interface,approximations can lead to insightful results,as will be shown in the following sections.In this work,three different activation energies are assumed in the calculations.3.1.Case1:infinite activation energyUsing the described model,it is possible to check that a very high value of the activation energy(higher than 300kJ molÀ1)leads to an interface mobility low enough to be considered nonexistent over any reasonabletimescale Fig.1.Schematic diagram illustrating the austenite interface composition under CCE conditions(dashed lines)and under equilibrium(dotted lines).(a) Carbon concentration in the austenite at the interface higher than the equilibrium concentration,and(b)carbon concentration in the austenite at the interface lower than the equilibrium concentration.4550M.J.Santofimia et al./Acta Materialia57(2009)4548–4557(up to days)during annealing at temperatures up to 500°C.For simplicity,the simulations have been done assuming an infinite value of the activation energy by set-ting the interface mobility equal to zero.This assumption leads to an immobile interface and to results corresponding to CCE conditions.3.2.Case2:Q M=180kJ molÀ1An activation energy for iron migration equal to 180kJ molÀ1was selected for this case in order to simulate the situation of limited martensite–austenite mobility, slower than for austenite to ferrite transformations.This value of the activation energy should be considered illustra-tive for coherent or semicoherent interfaces rather thanquantitatively accurate,since currently there is no basis for an accurate estimation of the activation energy for movement of iron atoms at the martensite–austenite interface.3.3.Case3:Q M=140kJ molÀ1In this case,the activation energy for interface migration was set equal to140kJ molÀ1,which is the value used by Krielaart and Van der Zwaag in a study on the austenite to ferrite transformation behavior of binary Fe–Mn alloys [15]and by Mecozzi et al.[16]to study the same phase transformation in a Nb microalloyed CMn steel.The resulting mobility can be seen as an upper limit,applying to incoherent interfaces.Model predictions are sensitive to the alloy used in the calculations.In this work,simulations have been per-formed assuming a binary Fe–0.25wt.%C system and a martensite–austenitefilm morphology(also used in Ref.[14]).The corresponding martensite start temperature (M s)was calculated to be433°C[24].Simulations consid-ered two annealing temperatures(350and400°C)and dif-ferent quenching temperatures ranging from220to400°C. Values of the martensite–austenite interface mobility M, calculated according to Eqs.(2)and(3)for both annealing temperatures studied,are presented in Table1.Variations in the quenching temperatures lead to different amounts of martensite and austenite prior to annealing.The volume fractions of phases present after the quenching step are esti-mated by the Koistinen–Marburger equation[25],leading to the values shown in Table2.The volume fractions of phases present at each quenching temperature and the lath widths of martensite and austenite are related using a‘‘con-stant ferrite width approach”[26].This approach is based on the transmission electron microscopy observations of Krauss and co-workers,indicating that most martensitic lath widths range approximately from0.15to0.2l m [27,28].Additionally,Marder[29]reported that a lath width of0.2l m was most frequently observed for 0.2wt.%C martensite.Therefore,a constant martensite lath width equal to0.2l m has been assumed for the initial conditions in the simulations.Corresponding austenite dimensions are obtained based on the appropriate austen-ite fraction predicted for every quenching temperature,and the results are shown in Table2.The volume fraction of martensite during annealing can be estimated from the size of the martensite domain at every annealing time.The local fraction of austenite that is stable upon quenching to room temperature is estimated by calculation of the M s temperature using Eq.(5),pre-sented in Ref.[24],across the austenite carbon profile and by further use of the Koistinen–Marburger[25]rela-tionship to estimate the volume fraction of stable austenite at each point[30].Final retained austenite fractions are cal-culated by integration of the area under each local fraction of stable austenite curve for different annealing times[31].Simulations of the interaction between carbon partition-ing and interface migration under the conditions explained above are presented and discussed with respect to the evo-lution of the phase fractions and phase compositions.4.Results and discussion4.1.Carbon profiles in martensite and austenite during annealingFigs.2and3show the evolution of carbon profiles in martensite and austenite during annealing at350and 400°C,respectively,assuming a quenching temperature of300°C and the three activation energies considered to describe interface mobility.The samefigures also show the estimation of the local retained austenite fraction when the material isfinally quenched to room temperature after annealing.A general observation is,in all cases,a sharp increase in the carbon content in the austenite close to the martensite–austenite interface at short annealing times.Table1Mobility(m4JÀ1sÀ1)corresponding to two activation energies and annealing temperatures studied.Annealing temperature(°C)Mobility forQ M=180kJ molÀ1Mobility forQ M=140kJ molÀ1350 2.45Â10À20 5.53Â10À17 400 2.99Â10À19 3.81Â10À16Table2Calculated martensite and austenite fractions present at each quenching temperature and corresponding martensite and austenite widths using the constant ferrite width approach.Quenchingtemperature(°C)Approximate fractionat quench temperatureLath orfilmwidth(l m)Austenite Martensite Austenite Martensite 2200.100.900.020.20 2500.130.870.030.20 2700.170.830.040.20 2890.200.800.050.20 3000.230.770.060.20 3200.290.710.080.20 3500.400.600.130.20 4000.690.310.450.20M.J.Santofimia et al./Acta Materialia57(2009)4548–45574551Starting with the results corresponding to annealing at 350°C,it is observed that,under stationary interface con-ditions (Fig.2a–c ),the sharp carbon profiles observed in the austenite at short annealing times are progressively reduced.After about 100s,the carbon concentration in both phases is equilibrated according to the conditions established by CCE,i.e.the same chemical potential of car-bon in the martensite and the austenite but with the limita-tion of an immobile interface.Fig.2c shows estimations of the local fraction of retained austenite,indicating that the final state corresponds to the retention of about half of the austenite available during annealing.When the activation energy is assumed equal to 180kJ mol À1(Fig.2d–f ),the interface mobility is not high enough to produce interface migration during the time-frame in which carbon partitioning occurs from the mar-tensite to the austenite.This behavior results in evolution of carbon profiles similar to that obtained with a stationary interface for annealing times lower than 100s (the time necessary to obtain the final profiles in the case of an immobile interface).However,at longer annealing times,there is interface migration from the martensite into the austenite until the establishment of full equilibrium in both phases,with a substantial reduction of the austenite frac-tion in this instance.The final profiles are obtained after annealing for about 10,000s ($3h).In this case,the vol-ume fraction of retained austenite at the end of the process (Fig.2f)is less than half of the austenite available after the first quench because of the reduction of the austenite thickness.When the activation energy is 140kJ mol À1(Fig.2g–i ),the interface mobility is high enough to produce migration of the martensite–austenite interface during carbon trans-fer between the two phases.Initially,the carbon content at the interface is higher than the equilibrium value and migration of the interface from the austenite into themar-Fig.2.Calculated carbon profiles in martensite (left column)and austenite (middle column)together with local austenite volume fraction that is stable to the final quench (right column)during annealing at 350°C after quenching to 300°C:(a–c)immobile interface;(d–f)Q M =180kJ mol À1;(g–i)Q M =140kJ mol À1.Arrows in the upper part of the figures and dashed lines indicate the movement of the martensite–austenite interface.According to Table 2,the combined thickness of one martensite plus one austenite film is 0.26l m when quenching at 300°C,but because of symmetry the calculation domain includes only the half-thickness,which is 0.13l m.4552M.J.Santofimia et al./Acta Materialia 57(2009)4548–4557tensite takes place.However,carbon diffusion causes a reduction of this peak in the time interval between 0.1and 1s,to carbon levels at the interface that are lower than the equilibrium value.Consequently,the interface then migrates from the martensite to the austenite.The homog-enization of carbon in the austenite leads to further movement of the interface until the carbon content corre-sponding to full equilibrium in both phases is reached after annealing for about 100s.The time taken to attain the final carbon profiles is similar to the that required in the case of an immobile interface,but considerably lower than for an activation energy of 180kJ mol À1.The final fraction of local retained austenite (Fig.2i)is the same as that obtained in the previous case.From the above results,it is clear that the interface mobility has an important influence on the kinetics of the carbon partitioning process.In the case of a stationary interface or when the interface mobility corresponds to the value determined for reconstructive austenite to ferrite transformations (Q M =140kJ mol À1),the final carbon profiles are obtained after annealing for a similar length of time (about 100s).However,in the case of an interme-diate interface mobility (Q M =180kJ mol À1),as might apply to a lower-energy semicoherent interface,the devel-opment of the carbon profiles is essentially similar to those obtained in the case of an immobile interface for times shorter than about 100s.However,longer annealing times lead to slow migration of the interface until full equilibrium conditions are reached after annealing for about 10,000s.In the case of annealing at 400°C (Fig.3),the evolution of the carbon profiles in martensite and austenite and local fractions of retained austenite is similar to those obtained for annealing at 350°C,but takes place on a different time-scale.For example,uniform carbon concentration profiles in the case of an immobile interface (Fig.3a and b)and Q M =140kJ mol À1(Fig.3g and h)are obtained inbothFig.3.Calculated carbon profiles in martensite (left column)and austenite (middle column)together with local austenite volume fraction that is stable to the final quench (right column)during annealing at 400°C after quenching to 300°C:(a–c)immobile interface;(d–f)Q M =180kJ mol À1;(g–i)Q M =140kJ mol À1.Arrows in the upper part of the figures and dashed lines indicate the movement of the martensite–austenite interface.According to Table 2,the combined thickness of one martensite plus one austenite film is 0.26l m when quenching at 300°C,but because of symmetry the calculation domain includes only the half-thickness,which is 0.13l m.M.J.Santofimia et al./Acta Materialia 57(2009)4548–45574553phases after annealing for about 10s.However,in the case of Q M =180kJ mol À1(Fig.3d and e),the time required to reach full equilibrium is substantially longer,in the range between 100and 1000s.This is a consequence of the low mobility of the interface.4.2.Evolution of the interface position during annealing Fig.4a and b shows the evolution of the interface posi-tion with annealing time for the case of quenching to 300°C and annealing at 350and 400°C,respectively.Fig.4c and d shows the corresponding evolution of the car-bon content in the austenite at the interface.The three curves give results for the three martensite–austenite inter-face mobilities considered in this work.Examination of these figures leads to the observations described below.In the case of an immobile interface,the carbon content in the austenite increases fast very early in the process (although this rapid increase in carbon is not represented in the timescale of Fig.4)and then decreases before reach-ing the value given by the constrained carbon equilibrium condition.For Q M =180kJ mol À1,the interface does not significantly change its position for annealing times lower than about 10s in the case of annealing at 350°C and about 1s for annealing at 400°C.During this time,the car-bon content in the austenite at the interface reaches the value corresponding to CCE,i.e.evolves identically to the case of an immobile interface.However,longer anneal-ing times lead to the initiation of interface migration from the martensite into the austenite and the progressive enrichment of carbon at the interface until full equilibriumconditions are reached.Finally,consideringQ M =140kJ mol À1,the evolution of the interface position and the carbon concentration in the austenite at the inter-face largely occur simultaneously during the annealing pro-cess.In this case,carbon partitioning starts with an increase of the carbon content in the austenite at the inter-face,which is compensated by the movement of the inter-face from the austenite into the martensite.Once the carbon content of the austenite is lower than the equilib-rium value,the motion of the interface reverses its direc-tion,from the martensite into the austenite.This migration ends when full equilibrium conditions are reached.4.3.Evolution of the volume fraction of martensite during annealingThe predicted evolution of the volume fraction 1of mar-tensite during annealing for the case of quenching to 300°C and annealing at 350or 400°C is shown in Fig.5.As expected,the volume fraction of martensite for the case of an immobile interface is constant.In the case of Q M =180kJ mol À1,the volume fraction of martensite is constant for annealing times below about 100s (annealing at 350°C)or 10s (annealing at 400°C).Afterwards,the vol-ume fraction of martensite increases by about 0.16.The evo-lution of the martensite volume fraction with annealing time is more complex for the case of Q M =140kJ mol À1.First,Fig.4.(a and b)Position of the martensite–austenite interface for quenching to 300°C and annealing at (a)350°C and (b)400°C.Position 0.00refers to the initial position of the interface and any decrease or increase of the position represents a decrease or increase,respectively,of the martensite width.(c and d)Carbon content in the austenite at the interface for quenching at 300°C and annealing at (c)350°C and (d)400°C.1Predicted volume fractions ignore any slight changes in the phase densities associated with carbon partitioning.4554M.J.Santofimia et al./Acta Materialia 57(2009)4548–4557。

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HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationFLUOXETINE HClC17H18F3NO•HClM.W. = 345.79CAS — 59333-67-4STABILITY INDICATINGA S S A Y V A L I D A T I O NMethod is suitable for:ýIn-process controlþProduct ReleaseþStability indicating analysis (Suitability - US/EU Product) CAUTIONFLUOXETINE HYDROCHLORIDE IS A HAZARDOUS CHEMICAL AND SHOULD BE HANDLED ONLY UNDER CONDITIONS SUITABLE FOR HAZARDOUS WORK.IT IS HIGHLY PRESSURE SENSITIVE AND ADEQUATE PRECAUTIONS SHOULD BE TAKEN TO AVOID ANY MECHANICAL FORCE (SUCH AS GRINDING, CRUSHING, ETC.) ON THE POWDER.ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationTABLE OF CONTENTS INTRODUCTION........................................................................................................................ PRECISION............................................................................................................................... System Repeatability ................................................................................................................ Method Repeatability................................................................................................................. Intermediate Precision .............................................................................................................. LINEARITY................................................................................................................................ RANGE...................................................................................................................................... ACCURACY............................................................................................................................... Accuracy of Standard Injections................................................................................................ Accuracy of the Drug Product.................................................................................................... VALIDATION OF FLUOXETINE HCl AT LOW CONCENTRATION........................................... Linearity at Low Concentrations................................................................................................. Accuracy of Fluoxetine HCl at Low Concentration..................................................................... System Repeatability................................................................................................................. Quantitation Limit....................................................................................................................... Detection Limit........................................................................................................................... VALIDATION FOR META-FLUOXETINE HCl (POSSIBLE IMPURITIES).................................. Meta-Fluoxetine HCl linearity at 0.05% - 1.0%........................................................................... Detection Limit for Fluoxetine HCl.............................................................................................. Quantitation Limit for Meta Fluoxetine HCl................................................................................ Accuracy for Meta-Fluoxetine HCl ............................................................................................ Method Repeatability for Meta-Fluoxetine HCl........................................................................... Intermediate Precision for Meta-Fluoxetine HCl......................................................................... SPECIFICITY - STABILITY INDICATING EVALUATION OF THE METHOD............................. FORCED DEGRADATION OF FINISHED PRODUCT AND STANDARD..................................1. Unstressed analysis...............................................................................................................2. Acid Hydrolysis stressed analysis..........................................................................................3. Base hydrolysis stressed analysis.........................................................................................4. Oxidation stressed analysis...................................................................................................5. Sunlight stressed analysis.....................................................................................................6. Heat of solution stressed analysis.........................................................................................7. Heat of powder stressed analysis.......................................................................................... System Suitability stressed analysis.......................................................................................... Placebo...................................................................................................................................... STABILITY OF STANDARD AND SAMPLE SOLUTIONS......................................................... Standard Solution...................................................................................................................... Sample Solutions....................................................................................................................... ROBUSTNESS.......................................................................................................................... Extraction................................................................................................................................... Factorial Design......................................................................................................................... CONCLUSION...........................................................................................................................ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationBACKGROUNDTherapeutically, Fluoxetine hydrochloride is a classified as a selective serotonin-reuptake inhibitor. Effectively used for the treatment of various depressions. Fluoxetine hydrochloride has been shown to have comparable efficacy to tricyclic antidepressants but with fewer anticholinergic side effects. The patent expiry becomes effective in 2001 (US). INTRODUCTIONFluoxetine capsules were prepared in two dosage strengths: 10mg and 20mg dosage strengths with the same capsule weight. The formulas are essentially similar and geometrically equivalent with the same ingredients and proportions. Minor changes in non-active proportions account for the change in active ingredient amounts from the 10 and 20 mg strength.The following validation, for the method SI-IAG-206-02 , includes assay and determination of Meta-Fluoxetine by HPLC, is based on the analytical method validation SI-IAG-209-06. Currently the method is the in-house method performed for Stability Studies. The Validation was performed on the 20mg dosage samples, IAG-21-001 and IAG-21-002.In the forced degradation studies, the two placebo samples were also used. PRECISIONSYSTEM REPEATABILITYFive replicate injections of the standard solution at the concentration of 0.4242mg/mL as described in method SI-IAG-206-02 were made and the relative standard deviation (RSD) of the peak areas was calculated.SAMPLE PEAK AREA#15390#25406#35405#45405#55406Average5402.7SD 6.1% RSD0.1ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::PRECISION - Method RepeatabilityThe full HPLC method as described in SI-IAG-206-02 was carried-out on the finished product IAG-21-001 for the 20mg dosage form. The method repeated six times and the relative standard deviation (RSD) was calculated.SAMPLENumber%ASSAYof labeled amountI 96.9II 97.8III 98.2IV 97.4V 97.7VI 98.5(%) Average97.7SD 0.6(%) RSD0.6PRECISION - Intermediate PrecisionThe full method as described in SI-IAG-206-02 was carried-out on the finished product IAG-21-001 for the 20mg dosage form. The method was repeated six times by a second analyst on a different day using a different HPLC instrument. The average assay and the relative standard deviation (RSD) were calculated.SAMPLENumber% ASSAYof labeled amountI 98.3II 96.3III 94.6IV 96.3V 97.8VI 93.3Average (%)96.1SD 2.0RSD (%)2.1The difference between the average results of method repeatability and the intermediate precision is 1.7%.HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationLINEARITYStandard solutions were prepared at 50% to 200% of the nominal concentration required by the assay procedure. Linear regression analysis demonstrated acceptability of the method for quantitative analysis over the concentration range required. Y-Intercept was found to be insignificant.RANGEDifferent concentrations of the sample (IAG-21-001) for the 20mg dosage form were prepared, covering between 50% - 200% of the nominal weight of the sample.Conc. (%)Conc. (mg/mL)Peak Area% Assayof labeled amount500.20116235096.7700.27935334099.21000.39734463296.61500.64480757797.52000.79448939497.9(%) Average97.6SD 1.0(%) RSD 1.0ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::RANGE (cont.)The results demonstrate linearity as well over the specified range.Correlation coefficient (RSQ)0.99981 Slope11808.3Y -Interceptresponse at 100%* 100 (%) 0.3%ACCURACYACCURACY OF STANDARD INJECTIONSFive (5) replicate injections of the working standard solution at concentration of 0.4242mg/mL, as described in method SI-IAG-206-02 were made.INJECTIONNO.PEAK AREA%ACCURACYI 539299.7II 540599.9III 540499.9IV 5406100.0V 5407100.0Average 5402.899.9%SD 6.10.1RSD, (%)0.10.1The percent deviation from the true value wasdetermined from the linear regression lineHPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::ACCURACY OF THE DRUG PRODUCTAdmixtures of non-actives (placebo, batch IAG-21-001 ) with Fluoxetine HCl were prepared at the same proportion as in a capsule (70%-180% of the nominal concentration).Three preparations were made for each concentration and the recovery was calculated.Conc.(%)Placebo Wt.(mg)Fluoxetine HCl Wt.(mg)Peak Area%Accuracy Average (%)70%7079.477.843465102.27079.687.873427100.77079.618.013465100.0101.0100%10079.6211.25476397.910080.8011.42491799.610079.6011.42485498.398.6130%13079.7214.90640599.413080.3114.75632899.213081.3314.766402100.399.618079.9920.10863699.318079.3820.45879499.418080.0820.32874899.599.4Placebo, Batch Lot IAG-21-001HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::VALIDATION OF FLUOXETINE HClAT LOW CONCENTRATIONLINEARITY AT LOW CONCENTRATIONSStandard solution of Fluoxetine were prepared at approximately 0.02%-1.0% of the working concentration required by the method SI-IAG-206-02. Linear regression analysis demonstrated acceptability of the method for quantitative analysis over this range.ACCURACY OF FLUOXETINE HCl AT LOW CONCENTRATIONThe peak areas of the standard solution at the working concentration were measured and the percent deviation from the true value, as determined from the linear regression was calculated.SAMPLECONC.µg/100mLAREA FOUND%ACCURACYI 470.56258499.7II 470.56359098.1III 470.561585101.3IV 470.561940100.7V 470.56252599.8VI 470.56271599.5(%) AverageSlope = 132.7395299.9SD Y-Intercept = -65.872371.1(%) RSD1.1HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationSystem RepeatabilitySix replicate injections of standard solution at 0.02% and 0.05% of working concentration as described in method SI-IAG-206-02 were made and the relative standard deviation was calculated.SAMPLE FLUOXETINE HCl AREA0.02%0.05%I10173623II11503731III10103475IV10623390V10393315VI10953235Average10623462RSD, (%) 5.0 5.4Quantitation Limit - QLThe quantitation limit ( QL) was established by determining the minimum level at which the analyte was quantified. The quantitation limit for Fluoxetine HCl is 0.02% of the working standard concentration with resulting RSD (for six injections) of 5.0%. Detection Limit - DLThe detection limit (DL) was established by determining the minimum level at which the analyte was reliably detected. The detection limit of Fluoxetine HCl is about 0.01% of the working standard concentration.ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::VALIDATION FOR META-FLUOXETINE HCl(EVALUATING POSSIBLE IMPURITIES)Meta-Fluoxetine HCl linearity at 0.05% - 1.0%Relative Response Factor (F)Relative response factor for Meta-Fluoxetine HCl was determined as slope of Fluoxetine HCl divided by the slope of Meta-Fluoxetine HCl from the linearity graphs (analysed at the same time).F =132.7395274.859534= 1.8Detection Limit (DL) for Fluoxetine HClThe detection limit (DL) was established by determining the minimum level at which the analyte was reliably detected.Detection limit for Meta Fluoxetine HCl is about 0.02%.Quantitation Limit (QL) for Meta-Fluoxetine HClThe QL is determined by the analysis of samples with known concentration of Meta-Fluoxetine HCl and by establishing the minimum level at which the Meta-Fluoxetine HCl can be quantified with acceptable accuracy and precision.Six individual preparations of standard and placebo spiked with Meta-Fluoxetine HCl solution to give solution with 0.05% of Meta Fluoxetine HCl, were injected into the HPLC and the recovery was calculated.HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::META-FLUOXETINE HCl[RECOVERY IN SPIKED SAMPLES].Approx.Conc.(%)Known Conc.(µg/100ml)Area in SpikedSampleFound Conc.(µg/100mL)Recovery (%)0.0521.783326125.735118.10.0521.783326825.821118.50.0521.783292021.55799.00.0521.783324125.490117.00.0521.783287220.96996.30.0521.783328526.030119.5(%) AVERAGE111.4SD The recovery result of 6 samples is between 80%-120%.10.7(%) RSDQL for Meta Fluoxetine HCl is 0.05%.9.6Accuracy for Meta Fluoxetine HClDetermination of Accuracy for Meta-Fluoxetine HCl impurity was assessed using triplicate samples (of the drug product) spiked with known quantities of Meta Fluoxetine HCl impurity at three concentrations levels (namely 80%, 100% and 120% of the specified limit - 0.05%).The results are within specifications:For 0.4% and 0.5% recovery of 85% -115%For 0.6% recovery of 90%-110%HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::META-FLUOXETINE HCl[RECOVERY IN SPIKED SAMPLES]Approx.Conc.(%)Known Conc.(µg/100mL)Area in spikedSample Found Conc.(µg/100mL)Recovery (%)[0.4%]0.4174.2614283182.66104.820.4174.2614606187.11107.370.4174.2614351183.59105.36[0.5%]0.5217.8317344224.85103.220.5217.8316713216.1599.230.5217.8317341224.81103.20[0.6%]0.6261.3918367238.9591.420.6261.3920606269.81103.220.6261.3920237264.73101.28RECOVERY DATA DETERMINED IN SPIKED SAMPLESHPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::REPEATABILITYMethod Repeatability - Meta Fluoxetine HClThe full method (as described in SI-IAG-206-02) was carried out on the finished drug product representing lot number IAG-21-001-(1). The HPLC method repeated serially, six times and the relative standard deviation (RSD) was calculated.IAG-21-001 20mg CAPSULES - FLUOXETINESample% Meta Fluoxetine % Meta-Fluoxetine 1 in Spiked Solution10.0260.09520.0270.08630.0320.07740.0300.07450.0240.09060.0280.063AVERAGE (%)0.0280.081SD 0.0030.012RSD, (%)10.314.51NOTE :All results are less than QL (0.05%) therefore spiked samples with 0.05% Meta Fluoxetine HCl were injected.HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationED. N0: 04Effective Date:APPROVED::Intermediate Precision - Meta-Fluoxetine HClThe full method as described in SI-IAG-206-02 was applied on the finished product IAG-21-001-(1) .It was repeated six times, with a different analyst on a different day using a different HPLC instrument.The difference between the average results obtained by the method repeatability and the intermediate precision was less than 30.0%, (11.4% for Meta-Fluoxetine HCl as is and 28.5% for spiked solution).IAG-21-001 20mg - CAPSULES FLUOXETINESample N o:Percentage Meta-fluoxetine% Meta-fluoxetine 1 in spiked solution10.0260.06920.0270.05730.0120.06140.0210.05850.0360.05560.0270.079(%) AVERAGE0.0250.063SD 0.0080.009(%) RSD31.514.51NOTE:All results obtained were well below the QL (0.05%) thus spiked samples slightly greater than 0.05% Meta-Fluoxetine HCl were injected. The RSD at the QL of the spiked solution was 14.5%HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationSPECIFICITY - STABILITY INDICATING EVALUATIONDemonstration of the Stability Indicating parameters of the HPLC assay method [SI-IAG-206-02] for Fluoxetine 10 & 20mg capsules, a suitable photo-diode array detector was incorporated utilizing a commercial chromatography software managing system2, and applied to analyze a range of stressed samples of the finished drug product.GLOSSARY of PEAK PURITY RESULT NOTATION (as reported2):Purity Angle-is a measure of spectral non-homogeneity across a peak, i.e. the weighed average of all spectral contrast angles calculated by comparing all spectra in the integrated peak against the peak apex spectrum.Purity Threshold-is the sum of noise angle3 and solvent angle4. It is the limit of detection of shape differences between two spectra.Match Angle-is a comparison of the spectrum at the peak apex against a library spectrum.Match Threshold-is the sum of the match noise angle3 and match solvent angle4.3Noise Angle-is a measure of spectral non-homogeneity caused by system noise.4Solvent Angle-is a measure of spectral non-homogeneity caused by solvent composition.OVERVIEWT he assay of the main peak in each stressed solution is calculated according to the assay method SI-IAG-206-02, against the Standard Solution, injected on the same day.I f the Purity Angle is smaller than the Purity Threshold and the Match Angle is smaller than the Match Threshold, no significant differences between spectra can be detected. As a result no spectroscopic evidence for co-elution is evident and the peak is considered to be pure.T he stressed condition study indicated that the Fluoxetine peak is free from any appreciable degradation interference under the stressed conditions tested. Observed degradation products peaks were well separated from the main peak.1® PDA-996 Waters™ ; 2[Millennium 2010]ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationFORCED DEGRADATION OF FINISHED PRODUCT & STANDARD 1.UNSTRESSED SAMPLE1.1.Sample IAG-21-001 (2) (20mg/capsule) was prepared as stated in SI-IAG-206-02 and injected into the HPLC system. The calculated assay is 98.5%.SAMPLE - UNSTRESSEDFluoxetine:Purity Angle:0.075Match Angle:0.407Purity Threshold:0.142Match Threshold:0.4251.2.Standard solution was prepared as stated in method SI-IAG-206-02 and injected into the HPLC system. The calculated assay is 100.0%.Fluoxetine:Purity Angle:0.078Match Angle:0.379Purity Threshold:0.146Match Threshold:0.4272.ACID HYDROLYSIS2.1.Sample solution of IAG-21-001 (2) (20mg/capsule) was prepared as in method SI-IAG-206-02 : An amount equivalent to 20mg Fluoxetine was weighed into a 50mL volumetric flask. 20mL Diluent was added and the solution sonicated for 10 minutes. 1mL of conc. HCl was added to this solution The solution was allowed to stand for 18 hours, then adjusted to about pH = 5.5 with NaOH 10N, made up to volume with Diluent and injected into the HPLC system after filtration.Fluoxetine peak intensity did NOT decrease. Assay result obtained - 98.8%.SAMPLE- ACID HYDROLYSISFluoxetine peak:Purity Angle:0.055Match Angle:0.143Purity Threshold:0.096Match Threshold:0.3712.2.Standard solution was prepared as in method SI-IAG-206-02 : about 22mg Fluoxetine HCl were weighed into a 50mL volumetric flask. 20mL Diluent were added. 2mL of conc. HCl were added to this solution. The solution was allowed to stand for 18 hours, then adjusted to about pH = 5.5 with NaOH 10N, made up to volume with Diluent and injected into the HPLC system.Fluoxetine peak intensity did NOT decrease. Assay result obtained - 97.2%.ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationSTANDARD - ACID HYDROLYSISFluoxetine peak:Purity Angle:0.060Match Angle:0.060Purity Threshold:0.099Match Threshold:0.3713.BASE HYDROLYSIS3.1.Sample solution of IAG-21-001 (2) (20mg/capsule) was prepared as per method SI-IAG-206-02 : An amount equivalent to 20mg Fluoxetine was weight into a 50mL volumetric flask. 20mL Diluent was added and the solution sonicated for 10 minutes. 1mL of 5N NaOH was added to this solution. The solution was allowed to stand for 18 hours, then adjusted to about pH = 5.5 with 5N HCl, made up to volume with Diluent and injected into the HPLC system.Fluoxetine peak intensity did NOT decrease. Assay result obtained - 99.3%.SAMPLE - BASE HYDROLYSISFluoxetine peak:Purity Angle:0.063Match Angle:0.065Purity Threshold:0.099Match Threshold:0.3623.2.Standard stock solution was prepared as per method SI-IAG-206-02 : About 22mg Fluoxetine HCl was weighed into a 50mL volumetric flask. 20mL Diluent was added. 2mL of 5N NaOH was added to this solution. The solution was allowed to stand for 18 hours, then adjusted to about pH=5.5 with 5N HCl, made up to volume with Diluent and injected into the HPLC system.Fluoxetine peak intensity did NOT decrease - 99.5%.STANDARD - BASE HYDROLYSISFluoxetine peak:Purity Angle:0.081Match Angle:0.096Purity Threshold:0.103Match Threshold:0.3634.OXIDATION4.1.Sample solution of IAG-21-001 (2) (20mg/capsule) was prepared as per method SI-IAG-206-02. An equivalent to 20mg Fluoxetine was weighed into a 50mL volumetric flask. 20mL Diluent added and the solution sonicated for 10 minutes.1.0mL of 30% H2O2 was added to the solution and allowed to stand for 5 hours, then made up to volume with Diluent, filtered and injected into HPLC system.Fluoxetine peak intensity decreased to 95.2%.ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationSAMPLE - OXIDATIONFluoxetine peak:Purity Angle:0.090Match Angle:0.400Purity Threshold:0.154Match Threshold:0.4294.2.Standard solution was prepared as in method SI-IAG-206-02 : about 22mg Fluoxetine HCl were weighed into a 50mL volumetric flask and 25mL Diluent were added. 2mL of 30% H2O2 were added to this solution which was standing for 5 hours, made up to volume with Diluent and injected into the HPLC system.Fluoxetine peak intensity decreased to 95.8%.STANDARD - OXIDATIONFluoxetine peak:Purity Angle:0.083Match Angle:0.416Purity Threshold:0.153Match Threshold:0.4295.SUNLIGHT5.1.Sample solution of IAG-21-001 (2) (20mg/capsule) was prepared as in method SI-IAG-206-02 . The solution was exposed to 500w/hr. cell sunlight for 1hour. The BST was set to 35°C and the ACT was 45°C. The vials were placed in a horizontal position (4mm vials, National + Septum were used). A Dark control solution was tested. A 2%w/v quinine solution was used as the reference absorbance solution.Fluoxetine peak decreased to 91.2% and the dark control solution showed assay of 97.0%. The difference in the absorbance in the quinine solution is 0.4227AU.Additional peak was observed at RRT of 1.5 (2.7%).The total percent of Fluoxetine peak with the degradation peak is about 93.9%.SAMPLE - SUNLIGHTFluoxetine peak:Purity Angle:0.093Match Angle:0.583Purity Threshold:0.148Match Threshold:0.825 ED. N0: 04Effective Date:APPROVED::HPLC ASSAY with DETERMINATION OF META-FLUOXETINE HCl.ANALYTICAL METHOD VALIDATION10 and 20mg Fluoxetine Capsules HPLC DeterminationSUNLIGHT (Cont.)5.2.Working standard solution was prepared as in method SI-IAG-206-02 . The solution was exposed to 500w/hr. cell sunlight for 1.5 hour. The BST was set to 35°C and the ACT was 42°C. The vials were placed in a horizontal position (4mm vials, National + Septum were used). A Dark control solution was tested. A 2%w/v quinine solution was used as the reference absorbance solution.Fluoxetine peak was decreased to 95.2% and the dark control solution showed assay of 99.5%.The difference in the absorbance in the quinine solution is 0.4227AU.Additional peak were observed at RRT of 1.5 (2.3).The total percent of Fluoxetine peak with the degradation peak is about 97.5%. STANDARD - SUNLIGHTFluoxetine peak:Purity Angle:0.067Match Angle:0.389Purity Threshold:0.134Match Threshold:0.8196.HEAT OF SOLUTION6.1.Sample solution of IAG-21-001-(2) (20 mg/capsule) was prepared as in method SI-IAG-206-02 . Equivalent to 20mg Fluoxetine was weighed into a 50mL volumetric flask. 20mL Diluent was added and the solution was sonicated for 10 minutes and made up to volume with Diluent. 4mL solution was transferred into a suitable crucible, heated at 105°C in an oven for 2 hours. The sample was cooled to ambient temperature, filtered and injected into the HPLC system.Fluoxetine peak was decreased to 93.3%.SAMPLE - HEAT OF SOLUTION [105o C]Fluoxetine peak:Purity Angle:0.062Match Angle:0.460Purity Threshold:0.131Match Threshold:0.8186.2.Standard Working Solution (WS) was prepared under method SI-IAG-206-02 . 4mL of the working solution was transferred into a suitable crucible, placed in an oven at 105°C for 2 hours, cooled to ambient temperature and injected into the HPLC system.Fluoxetine peak intensity did not decrease - 100.5%.ED. N0: 04Effective Date:APPROVED::。

Effect of annealing treatment on electroluminescence from GaN/Si nanoheterostructure arrayChang Bao Han, Chuan He, Xiao Bo Meng, Ya Rui Wan, Yong Tao Tian, Ying Jiu Zhang, and Xin Jian Li*Department of Physics and Laboratory of Materials Physics, Zhengzhou University, Zhengzhou 450052, China * lixj@Abstract: A GaN/Si nanoheterostructure array was prepared by growing GaN nanostructures on silicon nanoporous pillar array (Si-NPA). Based on as-grown and annealed GaN/Si-NPA, two light-emitting diodes (LEDs) were fabricated. It was found that after the annealing treatment, both the turn-on voltage and the leakage current density of the nanoheterostructure varied greatly, together with the electroluminescence (EL) changed from a yellow band to a near infrared band. The EL variation was attributed to the radiative transition being transformed from a defect-related recombination in GaN to an interfacial recombination of GaN/Si-NPA. 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Introduction In the past decade, gallium nitride (GaN) has been widely used in fabricating ultraviolet, blue and green light-emitting diodes (LEDs) or laser diodes (LDs) because of the merits of its direct and wide bandgap (3.4 eV), high carrier mobility, and good thermal and chemical stability [1–3]. Although GaN/Si heterostructures were also deemed to be promising candidates for making integrated high-speed or high-power photoelectronic devices [4–6], the practical course was badly baffled by the inferior interface quality resulted from the large lattice mismatch between the two semiconductors [7], because an inferior interface quality would bring serious damage to both the inner quantum efficiency (IQE) and the performances of as-constructed devices. To reduce the lattice mismatch and thereby improve the interface#161177 - $15.00 USD (C) 2012 OSAReceived 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5637quality, two main approaches have been developed in the past years [8, 9]. One was the usage of an intermediate layer to accommodate most of the lattice mismatch and enhance the metallurgical compatibility between GaN and silicon substrates. This was usually realized by introducing a specially designed multilayer with a lattice constant gradient or an N-ion implanted substrate surface with partially distorted crystal lattices [10, 11]. The other was the adoption of nanoheteroepitaxy method, which was often carried out by utilizing various nanopatterned substrates to release the interfacial stress, reduce the dislocation density and improve the IQE [12–14]. For example, the GaN-Si interfacial residual stress could be greatly reduced by growing GaN film on a silicon nanopore array [15], and good rectifying properties could be obtained by growing p-GaN nanowires on n-Si crystal wafers [16]. These results greatly promoted the confidence for making GaN/Si-based LEDs or LDs with high efficiency and broad emission band range. Encouraged by these experiments and utilizing silicon nanoporous pillar array (Si-NPA) [17] as functional substrates, we have constructed a GaN/Si nanoheterostructure array (GaN/Si-NPA) by growing GaN nanograins onto Si-NPA, in which an effective yellow or infrared (NIR) electroluminescence (EL) tuned by the applied voltages was obtained [18]. This indicates that GaN/Si-NPA might be a promising material system for fabricating practical GaN/Si-based LEDs. According to the basic theory of luminescence, the adjustability of the EL wavelength inferred that there might have different radiative recombination paths in GaN/Si-NPA, such as the band-band transition or the transitions relating with the high-density defects formed in GaN or at the interfaces. Clearly, the coexistence of multi-recombination paths would produce strong effect on the EL qualities, both the EL intensity and monochromaticity. On the other hand, thermal treatments have been proved to be an effect approach to promote the EL properties of a semiconductor heterojunction through improving the interfacial quality and changing the electronic structures. For instance, the carrier concentration of n-ZnO/p-Si could be changed through annealing treatment and the J-V curve as well as the EL properties could be adjusted notably [19]. The EL intensity and peak position of n-ZnO nanorods/p-GaN LED could be tuned through controlling the concentration and sorts of the defect states by performing annealing treatments at different temperatures and in different atmospheres [20]. As a result, a systematic study of the annealing effect on the EL properties is necessary for both clarifying the luminescent mechanism and promoting the emission qualities of GaN/Si-NPA. In this paper, two GaN/Si LEDs were prepared based on as-grown and annealed GaN/SiNPA. The structural and physical properties, including the X-ray diffraction (XRD) patterns, surface morphologies, current density-voltage (J-V) curves, EL and photoluminescence (PL) spectra, were measured and comparatively studied. Based on the experimental results, the EL mechanisms of the LEDs were put forward through building up the corresponding electronics structures. Our results might indicate a novel approach for designing and fabricating highperformance LEDs based directly on GaN/Si nanoheterostructures. 2. Experimental details Si-NPA was prepared by hydrothermally etching (111) oriented, boron-doped single crystal Si (sc-Si) wafers in a solution of hydrofluoric acid containing ferric nitrate [17]. A thin layer of platinum (~3 nm), which acted as catalyst in the subsequent GaN growing process, was predeposited on freshly prepared Si-NPA samples by a magnetron sputtering technique. Using high-purity metal Ga (99.999%, 0.8 g) and NH3 gas (99.999%, introduced with a flow rate of 20 sccm) as the sources for the two elements, GaN were grown on Si-NPA by a chemical vapor deposition (CVD) method. The deposition was carried out in a vacuum tube furnace equipped with multichannel gas inlets and a gas mixing chamber at 1050 °C for 20 min. Here two kinds of GaN/Si-NPA were prepared, one was the as-grown sample and the other was annealed at 800 °C for 3 hours in nitrogen atmosphere afterwards. Layers of indium tin oxide (ITO, ~100 nm) acting as top electrode and Al (~500 nm) acting as back electrode were deposited by magnetron sputtering and vacuum evaporation methods, respectively. Asconstructed LEDs have a device structure of ITO/n-GaN/p-Si-NPA/sc-Si/Al. For the#161177 - $15.00 USD (C) 2012 OSAReceived 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5638convenience of narration, the two LEDs here were named as as-grown LED and annealed LED, respectively. The two LEDs were both annealed at 300 °C for 1 hour in Ar atmosphere to realize ohmic contact between the electrodes and the semiconductors. The active areas of the diodes were specified as 10 mm × 10 mm. The surface morphology and the crystal structure of GaN/Si-NPA were characterized by a field emission scanning electron microscope (FESEM, JSM 6700F) and an X-ray diffractometer (Panalytical X' Pert Pro). The electrical and luminescent properties of the devices were measured at room temperature through an electrical group system consisted of Sourcemeter-2400 (Keithley) and a fluorescence spectrometer (Spex Fluorolog-3), respectively. 3. Results and discussion The XRD patterns of as-grown and annealed GaN/Si-NPA are shown in Part A of Fig. 1(a), in which all the diffraction peaks were indexed to crystalline hexagonal wurtzite GaN (JCPDS card: No. 50-0792). The obvious difference between the two curves is the reduction of the full width at half maximum (FWHM) for all the corresponding diffraction peaks after annealing treatment, as could be seen more obviously in Part B of Fig. 1(a). The typical cross-sectional FESEM image of as-grown GaN/Si-NPA is given in Fig. 1(b), in which GaN layers characterized by two different morphologies were observed. The upper layer was composed of two kinds of quasi one-dimensional GaN nanostructures, straight nanowires with an average diameter of ~30 nm and pencil-like nanorods with an average diameter of ~300 nm. Both the nanowires and the nanorods were well separated and nearly aligned locally perpendicular to the substrate surface, with an average length of ~1.5 µm. Between the nanowire/nanorod layer and Si-NPA substrate was a granular layer consisted of large quantities of GaN nanocrystallites (nc-GaN). The layer thickness and the average grain size were ~150 nm and ~20 nm, respectively. No apparent morphological variation was found by comparing the FESEM images of the samples before and after annealing treatment. Therefore, it was reasonable to think that the reduction of the FWHM of the diffraction peaks observed in Fig. 1(a) should result from the growing up of nc-GaN, which might have been formed in the GaN granular layer. This indicates that the crystallinity of as-deposited nc-GaN might have been greatly improved after the annealing treatment. As a consequence, the density of crystal defects should have been largely reduced.Fig. 1. (a) Part A: the XRD patterns of as-grown and annealed GaN/Si-NPA; Part B: the comparison of the FWHM variation for all the corresponding XRD peaks before and after annealing treatment. (b) The cross-sectional FESEM image of as-grown GaN/Si-NPA.The dark J-V curves of as-grown and annealed LEDs measured at room temperature are depicted in Fig. 2. The inset of Fig. 2 shows the schematic structure of the LEDs. Both of the J-V curves exhibited rectifying characteristic. Because the contact between Si-NPA and sc-Si has been proved to be ohmic [21], the observed rectification behaviors confirmed the formation of heterojunctions for both as-grown and annealed GaN/Si-NPA. But all the junction parameters for annealed LEDs, including the turn-on voltage, breakdown reverse#161177 - $15.00 USD (C) 2012 OSAReceived 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5639voltage and leakage current density, have changed largely compared with those for as-grown ones. For example, the turn-on voltage (for obtaining a current density of 1 mA/cm2) increased from ~1.6 V to ~3.9 V, and the leakage current density (at an applied voltage of −4 V) reduced from ~3.2 mA/cm2 to ~0.04 mA/cm2. The rectiðcation ratios for the two LEDs were calculated to be ~9 (at ± 3.9 V) and ~36 (at ± 4.8 V), respectively. According to the basic theory of heterojunctions [22], the leakage current density of a heterojunction was generally attributed to the defect-mediated tunneling effect caused by a high defect or trap concentration at the interface. Therefore, the distinct reduction of the leakage current density for the annealed LED might indicate an improvement of interfacial quality and a decrease of the defect state density, just as what occurred in the annealing process of ZnO nanorods/Si heterojunctions [23].Fig. 2. The room-temperature J-V curves of as-grown and annealed GaN/Si-NPA. Inset: the schematic diagram of the LEDs.For clarifying the underlying transportation mechanism of the variation, the log-log plot of the J-V data is presented in Fig. 3. It was found that both the curves for as-grown and annealed LEDs could be fitted by two straight lines. For as-grown LED, the J-V curve exhibits firstly a linear relation at a low forward voltage region (V < 0.9 V, region I). This indicates that the transportation of the carriers obeying the Ohmic law. With the applied voltage increased over 0.9 V (region II), the J-V curve exhibits an exponential relationship (J~V3.3), which infers a typical space charge limited current (SCLC) mechanism [24]. The SCLC mechanism was usually observed in wide bandgap p-n diodes, such as ZnO/Si [25, 26] and ZnO/SiC [27]. As for annealed LED, the J-V curve also exhibits a linear relation before the inflection point of ~1.9 V (region I′), but the current density is about three orders of magnitude lower than that of as-grown device. With the applied voltage increased beyond ~1.9 V (region II′), the transportation mechanism also transferred to the SCLC model, but with a relationship of J~V10. Clearly, the exponent varied largely from ~3.3 for as-grown LED to ~10 for annealed LED, and the increment of the exponent in SCLC model indicated a narrowed distribution of the localized states and a lowered defect state density in the annealed LED [28].#161177 - $15.00 USD (C) 2012 OSAReceived 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5640Fig. 3. The room-temperature log-log plots for the forward J-V of as-grown and annealed GaN/Si-NPA.The EL spectra of the two LEDs are presented in Fig. 4(a). Under an applied forward bias of 10 V, the current densities for as-grown and annealed LEDs were ~205 mA/cm2 and ~185 mA/cm2, respectively, and both devices exhibited efficient EL with relatively good monochromaticity, a yellow band peaked at ~567 nm and with a FWHM of ~23.5 nm for asgrown LED, and a NIR band peaked at ~830 nm and with a FWHM of ~18.5 nm for annealed LED. To clarify the origins of the EL for the two LEDs, the PL spectra of as-grown and annealed GaN/Si-NPA were measured, as is given in Fig. 4(b). Under the excitation of an ultraviolet with a wavelength of 320 nm (using a 300 W Xe lamp as the light source and operated with a slit width of 3 nm), both samples showed a strong ultraviolet PL band centered at 366 nm. These bands were due to the near-band-edge (NBE) emission of crystal GaN. A weak but distinct yellow PL band peaked at ~560 nm was observed in as-grown GaN/Si-NPA, but it almost disappeared in the annealed samples. The yellow PL band was usually attributed to the radiative recombination related with the deep-level defects in undoped GaN [29–31], such as gallium vacancies (VGa), which has a relatively small formation energy and low migration barrier [29, 30, 32]. Therefore it is rational to deduce that the yellow PL and EL band for as-grown GaN/Si-NPA have the same origin, i. e., both of them originate from a defect-related radiative recombination process. The presentation and disappearance of the yellow PL band before and after annealing treatment reflected a reduction of defect density in GaN.Fig. 4. The room-temperature (a) EL spectra of as-drown and annealed GaN/Si-NPA LEDs operating with a dc voltage of 10 V, and (b) PL spectra of as-grown and annealed GaN/SiNPA.The mechanism of the EL could be explained through analyzing the energy diagram of GaN/Si-NPA (Fig. 5). As has been reported, the electron affinity ψ of GaN and Si-NPA was ~4.1 eV and ~3.6 eV, and their bandgap Eg were ~3.4 eV and ~2.0 eV, respectively [33–35]. The barrier heights at the interfaces for the conduction bands and valence bands, ∆EC and ∆EV, were calculated to be ~0.5 eV and ~1.9 eV, respectively. As a result, the interfacial band#161177 - $15.00 USD (C) 2012 OSAReceived 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5641off-set ∆E = Eg1(Si-NPA) - ∆EC = Eg2(GaN) - ∆EV = 1.5 eV, which is equal to the difference of their quasi-Fermi levels after applying bias. In terms of the depletion region of a heterojuction, the distribution of the depletion region is proportional to the built-in field but inversely proportional to the doping concentration. The low defect density for annealed GaN/Si-NPA would surely lead to a low doping concentration. Compared with as-grown LED, the built-in field for annealed LED is wider, which will result in a larger turn-on voltage (Fig. 2 and Fig. 5). When a larger forward bias was applied, the yellow EL band peaked at ~567 nm (~2.2 eV) observed in as-grown LED most probably originated from the radiative recombination of the deep-level defect states, such as VGa, in GaN (Fig. 5(a)), just as what occurred in the PL process. The VGa, which would form deep acceptor level in GaN, can accept the electrons transited from conduction band and give yellow emission under the excitation of electric field [32]. Furthermore, the low energy barrier (∆E’ = 2.2 - 1.5 = 0.7 eV) between the valence bands of Si-NPA and deep acceptor levels will be favorable to the injection of holes for realizing the yellow emission. On the other hand, for annealed LED, the NIR EL peaked at ~830 nm (1.5 eV) is attributed to interfacial transition between electrons in the conduction band of GaN and holes in the valence band of Si-NPA for the higher ∆EV (1.9 eV), as is shown in Fig. 5(b) [18]. As-grown LED exhibits a high leakage current density (high defect density), in which nonradiative recombination would dominate the interfacial recombination. So only the yellow EL originating from the GaN defect states could be observed. After the annealing treatment, both the defect density in GaN and at the interface would be reduced notably, thus the radiative recombination would mainly occur through band-band transition at the interface. In addition, the ∆E (1.5 eV) is much lower than the energy bandgap of Si-NPA (~2.0 eV) and GaN (~3.4 eV), so the transition probability at the interface would be much higher than the NBE transition within either Si-NPA or GaN according to quantum theory [36]. As a result, the annealing process could effectively tune the EL of GaN/Si-NPA from yellow band to NIR band.Fig. 5. The mechanism illustration of the yellow EL from as-grown (a) and NIR EL from annealed GaN/Si-NPA LEDs based on the energy band diagram.Just as discussed above, the origins of the yellow and NIR luminescence from as-grown and annealed GaN/Si-NPA LEDs were attributed to the defect-related radiative transition in GaN and the rediative recombination at the interface of GaN/Si, respectively. Clearly, for the LEDs based on two semiconductors with large lattice mismatch, the promotion on the relatively low IQE is crucial for its practical device application. In addition to further improvement of the material quality of GaN/Si-NPA through optimizing the preparing conditions, such as controlling the microstructure and surface chemical status of Si-NPA, changing the CVD preparing and post-treating parameters, adopting different LED fabrication arts or procedures, some recently developed approaches could also be used for references. The representative demonstrations include the fabrication of nonpolar InGaN quantum well (QW)#161177 - $15.00 USD (C) 2012 OSA Received 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5642LEDs [37] or InGaN QW LEDs with large optical matrix elements [38] and surface plasmon coupling [39, 40], through which the radiative recombination rate was greatly improved. This might indicate a promising path for achieving high IQE in the LEDs and might be utilized in preparing GaN/Si-NPA-based LEDs. 4. Conclusions In conclusion, a GaN/Si nanoheterostructure array was prepared by depositing n-GaN on p-SiNPA. Two prototype GaN/Si LEDs with a device structure of ITO/n-GaN/p-Si-NPA/sc-Si/Al were prepared based on as-grown and annealed GaN/Si-NPA, from which sharp yellow and NIR EL bands were obtained correspondingly. The origins for the yellow and NIR EL were attributed to the defect-related radiative recombination in GaN and the interfacial band-band rediative recombination of GaN/Si-NPA, respectively. Our research might have provided a doable approach for fabricating visible and NIR LEDs directly based on GaN/Si heterostructures. Acknowledgments This work was supported by the National Natural Science Foundation of China (No. 61176044, 11074224), the Sci.-Tech. Project for Innovative Scientist of Henan Province (No. 114200510017) and the Science and Technology Project on Key Problems of Henan Province (No. 082101510007).#161177 - $15.00 USD (C) 2012 OSAReceived 10 Jan 2012; revised 9 Feb 2012; accepted 9 Feb 2012; published 22 Feb 2012 27 February 2012 / Vol. 20, No. 5 / OPTICS EXPRESS 5643。

焊接专业英语词汇最全版本焊接专业英语词汇（共10部分）焊接专业英语词汇（共10部分）1、熔化焊熔接fusion welding压接pressure welding焊接过程welding process焊接技术welding technique焊接工艺welding technology/procedure焊接操作welding operation焊接顺序welding sequence焊接方向direction of welding焊接位置welding position熔敷顺序build-up sequence/deposition seq uence焊缝倾角weld slope/inclination of weld ax is焊缝转角weld rotation/angle of rotation平焊位置flat position of welding横焊位置horizontal position of welding立焊位置vertical position of welding仰焊位置overhead position of welding平焊downhand welding/flat position welding横焊horizontal position welding立焊vertical position welding仰焊overhead position welding全位置焊all position welding：熔焊时，焊件接逢所处空间位置包括平焊、横焊、仰焊等位置所进行的焊接。

如水平固定管所进行的环缝焊接向下立焊vertical down welding/downward welding in the vertical position向上立焊vertical up welding/upward weldi ng in the vertical position倾斜焊inclined position welding上坡焊upward welding in the inclined po sition下坡焊downward welding in the inclined position对接焊butt welding角焊fillet welding搭接焊lap welding船形焊fillet welding in the downhand pos ition/fillet welding in the flat position平角焊horizontal fillet welding立角焊fillet welding in the vertical positio n仰角焊fillet welding in the overhead posit ion坡口焊groove weldingI形坡口对接焊square butt welding喇叭形坡口焊flare groove welding卷边焊flanged edge welding纵缝焊接welding of longitudinal seam横缝焊接welding of transverse seam环缝焊接girth welding/ circumferential螺旋缝焊接welding of spiral seam/welding of helical seam环缝对接焊butt welding of circumferential seam定位焊tack welding单面焊welding by one side双面焊welding by both sides单道焊single pass welding/single run wel ding多道焊multi-pass welding单层焊single layer welding多层焊multi-layer welding分段多层焊block sequence/ block welding 分层多道焊multi-layer and multi-pass weld ing连续焊continuous welding断续焊intermittent welding打底焊backing weld封底焊back sealing weld盖面焊cosmetic welding深熔焊deep penetration welding摆动焊welding with weaving/weave bead welding前倾焊foreward welding (英国)/ forehand welding (美国)后倾焊backward welding(英国)/ backhand welding(美国)分段退焊backstep welding跳焊skip welding对称焊balanced welding/ balanced weldin g sequence左焊法leftward welding forehand welding 右焊法rightward welding/backhand welding挑弧焊whipping method自动焊automatic welding手工焊manual welding/hand welding车间焊接shop welding工地焊接site welding(英国)/ field welding (美国)拘束焊接restraint welding堆焊surfacing/building up/overlaying隔离层堆焊buttering端部周边焊boxing/end return返修焊rewelding补焊repair welding塞焊plug welding槽焊slot welding衬垫焊welding with backing焊剂垫焊welding with flux backing窄间隙焊narrow-gap welding强制成形焊enclosed welding脉冲电弧焊pulsed are welding电弧点焊arc spot welding螺柱焊stud welding热风焊hot gas welding高能焊high grade energy welding固态焊接solid-state welding单面焊双面成形one-side welding with bac k formation焊接条件welding condition焊接工艺参数welding parameter极性polarity正接electrode negative/straight polarity 反接electrode positive/reversed polarity 运条方式manipulation of electrode焊接电流welding current焊接电流增加时间welding current upslope time焊接电流衰减时间welding current downslo pe time电流密度current density短路电流short circuit current脉冲电流pulse level/pulse current level 脉冲电流幅值pulse current amplitude基值电流background level脉冲频率pulse frequency脉冲焊接电流占空比duty cycle of pulse du ration电弧电压arc voltage再引弧电压reignition voltage焊接速度welding speed行走速度rate of travel/travel speed送丝速度wire feed rate线能量heat input/energy input热输入heat input预热preheat后热postheat焊后热处理posweld heat treatment/posthe at treatment预热温度preheat temperature层间温度interpass temperature焊接终了温度finishing temperature后热温度postheating temperature焊丝伸出长度wire extension弧长arc length熔化速度melting rate熔化时间melting time熔化系数melting coefficient熔敷速度rate of deposition/deposition rate 熔敷系数deposition coefficient熔敷效率deposition efficiency损失系数loss coefficient飞spatter飞溅率spatter loss coefficient融合比fusion ratio稀释dilution稀释率rate of dilution合金过度系数transfer efficiency/recovery (of an element)坡口groove坡口面groove face坡口面角度angle of bevel (英国)/ bevel an gle (美国)坡口角度included angle(英国)/groove angle(美国)坡口高度groove depth钝边root face钝边高度thickness of root face/width of r oot face根部间隙root gap(英国)/root opening (美国)根部半径root radius/groove radius根部锐边root edge卷边高度height of flange卷边半径radius of flange单面坡口single groove双面坡口double groove坡口形式groove typeI形坡口square grooveV形坡口single V grooveY形坡口single V groove with root face双Y形坡口double Vgroove with root face 带钝边U形坡口single U groove带钝边双U形坡口double U grooveVY形坡口single compound angle groove 带钝边J形坡口single J groove带钝边双J形坡口double J groove单边V形坡口single bevel groove双V形坡口double V groove不对称双V形坡口asymmetric double V gr oove双单边V形坡口double bevel groove/K gr oove带垫板V形坡口V groove with backing/ si ngle V groove with backing喇叭形坡口flare groove锁底坡口single bevel groove with backing locked坡形板边tapered edge焊缝weld接逢seam焊缝符号welding symbol焊缝金属weld metal填充金属filler metal熔敷金属deposited metal焊缝表面weld face/ face of weld焊缝背面back of weld焊缝轴线axis of weld焊缝尺寸size of weld焊缝宽度weld width/ width of weld焊缝长度weld length/ length of weld焊缝有效长度effective length of weld焊缝厚度throat depth/ throat thickness焊缝计算厚度theoretical throat焊缝实际厚度actual throat熔深penetration/ depth of penetration焊缝成形appearance of weld焊缝成形系数form factor of weld余高reinforcement/ excess weld metal背面余高root reinforcement削平焊缝flush weld/ weld machined flush 对接焊缝butt weld角焊缝fillet焊脚leg/ fillet weld leg角焊缝断面形状profile of fillet weld平形角焊缝flat fillet凸形角焊缝convex fillet weld凹形角焊缝concave fillet weld角焊缝凹度concavity侧面角焊缝side fillet weld/ fillet weld in p arallel shear正面角焊缝front fillet weld/ fillet weld in normal shear立角焊缝fillet weld in the vertical position 横角焊缝fillet weld in the horizontal positi on平角焊缝fillet weld in the flat position斜角焊缝oblique fillet weld连续焊缝continuous weld断续焊缝intermittent weld连续角焊缝continuous fillet weld断续角焊缝intermittent fillet weld交错断续角焊缝staggered intermittent fillet weld并列断续角焊缝chain intermittent fillet wel d端接焊缝edge weld卷边焊缝flanged edge weld塞焊焊缝plug weld纵向焊缝longitudinal weld横向焊缝transverse weld环行焊缝girth weld/ circumferential weld 螺旋形焊缝spiral weld/ helical weld密封焊缝seal weld承载焊缝strength weld联系焊缝connective weld定位焊缝tack weld焊道bead/ run/ pass焊波ripple焊根weld root/ root of weld焊趾weld toe/ toe封底焊道sealing run (after making main weld)/ back weld打底焊道backing weld (before making mai n weld)/ back weld根部焊道root pass/ root run填充焊道filling bead盖面焊道cosmetic bead/ cover pass回火焊道temper bead/ annealing bead熔透焊道penetration bead焊层layer焊接接头welded joint接头形状joint geometry等强匹配接头equalmatching weld joint低强匹配接头undermatching weld joint超强匹配接头overmatching weld joint接头根部root of joint对接接头butt jointI形对接接头square butt jointV形对接接头single V butt jointU形对接接头single U butt jointJ形坡口接头single J butt joint双V形对接接头double V butt joint双单边V形对接接头double bevel butt joint / K groove butt joint带钝边U形对接接头double U butt joint带钝边J形坡口接头double J joint角接接头corner jointT形接头T joint斜T形接头inclined T joint十字接头cruciform joint/ cross-shaped joi nt三联接头joint among three members搭接接头lap joint套管接头muff joint/ sleeve joint双盖板接头double strapped joint盖板接头strapped joint端接接头edge joint卷边接头flanged edge joint锁底对接接头lock butt joint斜对接接头oblique butt joint混合接头mixed joint/ composite joint有间隙接头open joint无间隙接头closed joint焊接电弧welding arc电弧形态arc shape电弧物理行为arc physics behaviour引弧striking arc引弧电压striking voltage电弧气氛arc atmosphere阴极cathode热阴极hot cathode冷阴极cold cathode阴极斑点cathode spot阴极区cathode region阴极区电场强度intensity of the electric fiel d in the cathode region阴极压降cathode drop阳极anode阳极斑点anode spot斑点压力spot pressure阳极区anode region阳极区电场强度intensity of the electric fiel d in the anode region阳极压降anode drop弧柱arc column/ arc stream弧柱压降voltage drop in arc column弧柱电位梯度potential gradient in the arc column弧焰arc flame弧心arc core硬电弧forceful arc/ hard arc软电弧soft arc旋转电弧rotating arc脉冲电弧pulsed arc脉冲喷射电弧pulsed spray arc起皱现象puckering phenomena起皱电弧puckering arc起皱临界电流puckering critical current间接电弧indirect arc压缩电弧compressive arc磁控电弧magnetic controlling arc电弧力arc force电磁力electromagnectic force电磁收缩效应pinch effect电弧飘移wandering of arc电弧稳定性arc stability电弧静特性static characteristic of arc电弧动特性dynamic characteristic of arc 最小电压原理principle of minimum voltag e电弧挺度arc stiffness电弧偏吹arc blow磁偏吹magnetic blow阴极清理作用cleaning action of the catho de电弧自身调节arc self-regulation挖掘作用digging action极性效应polarity effect熔滴droplet熔滴比表面积specific surface of droplet 熔滴过渡metal transfer过度频率transition frequency粗滴过渡globular transfer; drop transfer 短路过渡short circuiting transfer喷射过渡spray transfer旋转喷射过渡rotating spray transfer脉冲喷射过渡pulsed spray transfer爆炸过渡explosive transfer渣壁过渡flux wall guided transfer熔池molten pool沸腾状熔池boiling molten pool弧坑crater熔渣slag渣系slag system渣系相图slag system diagram碱性渣basic slag酸性渣acid slag碱度basicity酸度acidity长渣long slag短渣short slag粘性熔渣viscous slag氧化物型熔渣oxide melting slag盐型熔渣salt melting slag盐-氧化物型熔渣salt-oxide melting slag 熔渣流动性fluidity of the slag; slag fluidi ty熔渣solidified slag多孔焊渣porous slag玻璃状焊渣vitreous slag自动脱落焊渣self-releasing slag脱渣性slag detachability焊接设备welding equipment; welding set 焊机welding machine; welder电焊机electric welding machine; electric welder焊接电源welding power source焊接热循环weld thermal cycle焊接温度场field of weld temperature; wel d temperature field准稳定温度场quasi-stationary temperature field焊接热源welding heat source点热源point heat source线热源linear heat source面热源plane heat source瞬时集中热源instantaneous concentration heat source热效率thermal efficiency热能集中系数coefficient of heat flow con centration峰值温度peak temperature瞬时冷却速度momentary cooling rate冷却时间cooling time置换氧化substitutionary oxydation扩散氧化diffusible oxydation脱氧desoxydation先期脱氧precedent desoxydation扩散脱氧diffusible desoxydation沉淀脱氧precipitation desoxydation扩散氢diffusible hydrogen初始扩散氢initial diffusible hydrogen 100℃残余扩散氢diffusible hydrogen rema ined at 100℃残余氢residual hydrogen去氢dehydrogenation去氢热处理heat treatment for dehydrogen ation脱硫desulphurization脱磷dephosphorization渗合金alloying微量合金化microalloying一次结晶组织primary solidification structu re二次结晶组织secondary solidification stru cture联生结晶epitaxial solidification焊缝结晶形态solidification mode in weld-bead结晶层状线ripple多边化边界polygonization boundary结晶平均线速度mean solidification rate针状铁素体acicular ferrite条状铁素体lath ferrite侧板条铁素体ferrite side-plate晶界欣素体grain boundary ferrite; polygo nal ferrite; pro-entectoid ferrite粒状贝氏体granular bainite板条马氏体lath martensite过热组织overheated structure魏氏组织Widmannst?tten structureM-A组元martensite-austenite constituent 焊件失效分析failure analysis of weldment s冷裂判据criterion of cold cracking冷裂敏感系数cold cracking susceptibity coefficient脆性温度区间brittle temperature range氢脆hydrogen embrittlement层状偏析lamellar segregation愈合healing effect断口金相fractography断口fracture延性断口ductile fracture韧窝断口dimple fracture脆性断口brittle fracture解理断口cleavage fracture准解理断口quasi-cleavage fracture氢致准解理断口hydrogen-embrittlement in duced沿晶断口intergranular fracture穿晶断口transgranular fracture疲劳断口fatigue fracture滑移面断口glide plane fracture断口形貌fracture apperance断口试验fracture test宏观断口分析macrofractography放射区radical zone纤维区fibrous zone剪切唇区shear lip aone焊接性weldability使用焊接性service weldability工艺焊接性fabrication weldability冶金焊接性metallurgical weldability热焊接性thermal weldability母材base metal; parent metal焊接区weld zone焊态as-welded (AW)母材熔化区fusion zone半熔化区partial melting region未混合区unmixed zone熔合区bond area熔合线weld junction (英)；bond line (美) 热影响区heat-affected zone (HAZ)过热区overheated zone粗晶区coarse grained region细晶区fine grained region过渡区transition zone硬化区hardened zone碳当量carbon equivalent铬当量chromium equivalent镍当量nickel equivalent舍夫勒组织图Schaeffler's diagram德龙组织图Delong’s diagram连续冷却转变图（CCT图）continuous cooli ng transformation裂纹敏感性cracking sensibility焊接裂纹weld crack焊缝裂纹weld metal crack焊道裂纹bead crack弧坑裂纹crater crack热影响区裂纹heat-affected zone crack纵向裂纹longitudinal crack横向裂纹transverse crack微裂纹micro-crack; micro-fissure热裂纹hot crack凝固裂纹solidification crack晶间裂纹intercrystalline crack穿晶裂纹transcrystalline crack多边化裂纹polygonization crack液化裂纹liquation crack失延裂纹ductility-dip crack冷裂纹cold crack延迟裂纹delayed crack氢致裂纹hydrogen-induced crack焊道下裂纹underbead crack焊根裂纹root crack焊趾裂纹toe crack锯齿形裂纹chevron cracking消除应力处理裂纹stress relief annealing c rack (SR crack)再热裂纹reheat crack焊缝晶间腐蚀weld intercryctalline corrosi on刀状腐蚀knife line attack敏化区腐蚀weld decay层状撕裂lamellar tearing焊接性试验weldability裂纹试验cracking testIIW裂纹试验IIW cracking testY形坡口裂纹试验slit type cracking test 分块形槽热裂纹试验segmented circular gr oove cracking testH形裂纹试验H-type cracking test鱼骨形裂纹试验fishbone cracking test指形裂纹试验finger (cracking) testT形裂纹试验Tee type cracking test环形槽裂纹试验circular-groove cracking t est可调拘束裂纹试验varestraint testBWRA奥氏体钢裂纹试验BWRA cracking t est for austenitie steel圆棒裂纹试验bar type cracking test; roun d bar cracking test里海裂纹试验Lehigh restraint cracking te st圆形镶块裂纹试验circular-path cracking te st十字接头裂纹试验cruciform cracking test Z向窗口拘束裂纹试验Z-direction window t ype restraint cracking testG-BOP焊缝金属裂纹试验G-BOP weld met al crack test巴特尔焊道下裂纹试验Battelle type underb ead cracking testU形拉伸试验U-tension test缪雷克期热裂纹试验Murex hot cracking te st菲斯柯裂纹试验FISCO (type) cracking test CTS裂纹试验controlled thermal severity 拉伸拘束裂纹试验（TRC试验）tensile restra int cracking test刚性拘束裂纹试验（RRC试验）rigid restrain t cracking test插销试验implant testTigamajig 薄板焊接裂纹试验Tigamajing thi n plate cracking test焊道纵向弯曲试验longitudinal-bead test柯麦雷尔弯曲试验Kommerell bead bend t est肯泽尔弯曲试验Kinzel test缺口弯曲试验notch bend test热朔性试验hot-ductility test热影响区冲击试验impact test of HAZ热影响区模拟试验synthetic heat-affected z one test最高硬度试验maximum hardness test落锤试验NRL (Naval Research Laboratory)测氢试验Hydrogen test焊接材料电极焊接材料welding consumables电极electrode熔化电极consumable electrode不熔化电极nonconsumable electrode钨电极tungsten electrode焊丝welding wire. Welding rod实心焊丝solid wire渡铜焊丝copper-plating welding wire自保护焊丝self-shielded welding wire药芯焊丝flux-cored wire复合焊丝combined wire堆焊焊丝surfacing welding rod填充焊丝filler wire焊条electrode/ covered electrode焊芯core wire药皮coating (of an electrode)/ covering (of an electrode)涂料coating flux/coating material造气剂gas forming constituents造渣剂slag forming constituents合金剂alloying constituent脱氧剂dioxidizer稳弧剂arc stabilizer粘接剂binder水玻璃water glass水玻璃模数modules of water glass酸性焊条acid electrode高钛型焊条high titania (type) electrode钛钙型焊条lime titania type electrode钛铁矿形焊条ilmenite type electrode氧化铁型焊条iron oxide type electrode/ hi gh iron oxide type electrode高纤维素型焊条high cellulose (type) electr ode石墨型焊条graphite type electrode碱性焊条basic electrode/ lime type covere d electrode低氢型焊条low hydrogen type electrode 高韧性超低氢焊条high toughness super lo w hydrogen electrode奥氏体焊条austenitic electrode铁素体焊条ferritic electrode不锈钢焊条stainless steel electrode珠光体耐热钢焊条pearlitic heat resistant st eel electrode低温钢焊条low temperature steel electrod e/ steel electrode for low temperature铝合金焊条aluminum alloy arc welding el ectrode铜合金焊条copper-alloy arc welding electr ode铜芯铸铁焊条cast iron electrode with stee l core纯镍铸铁焊条pure nickel cast iron electro de球墨铸铁焊条electrode for welding sphero idal graphite cast iron铸芯焊条electrode with cast core wire镍基合金焊条nickel base alloy covered el ectrode蒙乃尔焊条Monel electrode纯铁焊条pure iron electrode渗铝钢焊条alumetized steel electrode高效率焊条high efficiency electrode铁粉焊条iron powder electrode底层焊条backing welding electrode深熔焊条deep penetration electrode重力焊条gravity electrode立向下焊条electrode for vertical down po sition welding节能焊条saving energy electrode水下焊条underwater welding electrode耐海水腐蚀焊条seawater corrosion resista nt steel electrode低尘低毒焊条low-fume and harmfulless el ectrode/low-fume and low-toxic electrode堆焊焊条surfacing electrode耐磨堆焊焊条hardfacing electrode钴基合金堆焊焊条cobalt base alloy surfaci ng electrode碳化钨堆焊焊条tungsten carbide surfacing electrode高锰钢堆焊焊条high manganese steel surfacing electrode双芯焊条twin electrode绞合焊条stranded electrode编织焊条braided electrode双层药皮焊条double coated electrode管状焊条flux-cored electrode气渣联合保护型药皮semi-volatile covering 焊条工艺性usability of the electrode/ tech nicality of the electrode焊条使用性running characteristics of an e lectrode/ operating characteristics of an electrode焊条熔化性melting characteristics of an el ectrode焊条直径core diameter焊条偏心度eccentricity (of an electrode) 药皮重量系数gravity coefficient of coating 焊条药皮含水量percentage of moisture for covering焊条夹吃持端bare terminal (of an electrod e)焊条引弧端striking end (of an elcetrode)焊剂welding flux/ flux熔炼焊剂fused flux粘结焊剂bonded flux烧结焊剂sintered flux/ agglomerated flux 窄间隙埋弧焊焊剂flux for narrow-gap sub merged arc welding低氢型焊剂low hydrogen type flux高速焊剂high speed welding flux无氧焊剂oxygen-free flux低毒焊剂low poison flux磁性焊剂magnetic flux电弧焊arc welding直流电弧焊direct current arc welding交流电弧焊alternating current arc weldin g三相电弧焊three phase arc welding熔化电弧焊arc welding with consumable 金属极电弧焊metal arc welding不熔化极电弧焊arc welding with noncons umable碳弧焊carbon arc welding明弧焊open arc welding焊条电弧焊shielded metal arc welding (S MAW)重力焊gravity welding躺焊fire cracker welding电弧堆焊arc surfacing自动堆焊automatic surfacing躺板极堆焊surfacing by fire cracker weld ing带极堆焊surfacing with band-electrode振动电弧堆焊vibratory arc surfacing耐磨堆焊hardfacing埋弧焊submerged arc welding (SAW)多丝埋弧焊multiple wire submerged arc w elding纵列多丝埋弧焊Tandem sequence (subme rged-arc welding)横列多丝埋弧焊series submerged arc wel ding (SAW-S)横列双丝并联埋弧焊transverse submerged arc welding热丝埋弧焊hot wire submerged-arc weldi ng窄间隙埋弧焊narrow-gap submerged arc welding弧压反馈电弧焊arc voltage feedback cont rolling arc welding自调节电弧焊self-adjusting arc welding适应控制焊接adaptive control welding焊剂层burden; flux layer气体保护电弧焊gas shielded arc welding 保护气体protective atmosphere惰性气体inert-gas活性气体active-gas惰性气体保护焊inert-gas (arc) welding氩弧焊argon arc welding熔化极惰性气体保护电弧焊metal inert-gas arc welding钨极惰性气体保护电弧焊tungsten inert-gas arc welding钨极氢弧焊argon tungsten arc welding脉冲氢弧焊pulsed argon arc welding熔化极脉冲氢弧焊argon metal pulsed arc welding钨极脉冲氢弧焊argon tungsten pulsed arcwelding热丝MIG焊hot wire MIG welding热丝TIG焊hot wire TIG welding氨弧焊helium-arc welding活性气体保护电弧焊metal active-gas arc welding混合气体保护电弧焊mixed gas arc weldin g二氧化碳气体保护电弧焊carbon-dioxide arc welding; CO2 arc welding细丝CO2焊CO2 arc welding with thin wi re粗丝CO2焊CO2 arc welding with thick w ire磁性焊剂CO2焊unionarc welding药芯焊丝CO2焊arcos arc process; dual shield arc welding气电立焊electrogas (arc) welding氮弧焊nitrogen-arc welding水蒸气保护电弧焊water vapour arc weldin g原子氢焊atomic hydrogen welding冲器室中电弧焊controlled atmosphere arc welding旋转电弧焊rotating arc welding短路过渡电弧焊short circuiting arc weldin g焊丝横摆频率weaving speed of wire焊丝停摆时间electrode keep time of slide r等离子弧焊plasma arc welding (PAW)等离子弧plasma arc等离子流plasma jet转移弧transferred arc非转移弧nontransferred arc联合型等离子弧combined plasma arc主弧main arc维弧pilot arc维弧电流pilot arc surrent双弧现象double arcing双弧临界电流critical current of double ar cing等离子弧焊枪plasma (welding) torch压缩喷嘴constricting nozzle单孔喷嘴single port nozzle多孔喷嘴multiport nozzle压缩喷嘴孔径orifice diameter孔道长度orifice throat length孔道比orifice throat ratio等离子气plasma gas; orifice gas电极内缩长度electrode setback小孔效应keyhole effect小孔型等离子弧焊keyhole-mode welding 熔透型等离子弧焊fusion type plasma arc welding大电流等离子弧焊high-current plasma arc welding中电流等离子弧焊intermediate-current plas ma arc welding小电流等离子弧焊low-current plasma arc welding微束等离子弧焊micro-plasma arc welding 交流等离子弧焊AC plasma arc welding 脉冲等离子弧焊pulsed plasma arc weldin g等离子弧堆焊plasma arc surfacing热丝等离子弧堆焊hot wire plasma arc sur facing粉末等离子弧堆焊plasma arc powder surf acing等离子-熔化极惰性气体保护电弧焊plasma M IG welding转移弧电源transferred arc power supply 非转移弧电源nontransferred arc power su pply电弧焊设备arc welding equipment电弧焊机arc welding machine直流弧焊机DC arc welding machine交流弧焊机AC arc welding machine交直流两用弧焊机AC/DC arc welding mac hine单站弧焊机single operator arc welding m achine多站弧焊机multi-operator arc welding set 固定式弧焊机stationary arc welding mach ine移动式弧焊机portable arc welding machine台式弧焊机bench arc welding machine内燃机驱动式弧焊机combustion engine dri ven arc welding set电动机驱动式弧焊机motor driven arc weld ing set熔化极弧焊机arc welding machine using a consumable electrode不熔化极弧焊机arc welding machine usin g a non-consumable electrode脉冲弧焊机pulsed arc welding machine气体保护弧焊机gas shielded arc welding machine氩弧焊机argon arc welding machine二氧化碳弧焊机CO2 arc welding machine 钨极惰性气体保护弧焊机tungsten inert-gas welding machine熔化仍惰性气体保护弧焊机metal inert-gas welding machine气电立焊机electrogas (arc) welding mach ine等离子弧焊机plasma arc welding machine微束等离子弧焊机micro-plasma welding e quipment原子氢焊机atomic hydrogen welding app aratus埋弧焊机submerged arc welding machine 弧焊电源arc welding power source直流弧焊电源DC arc welding power sour ce交流弧焊电源AC arc welding power sour ce交直流两用弧焊电源AC/DC arc welding po wer source脉冲弧焊电源pulsed arc welding power s ource上升特性弧焊电源rising characteristic arc welding power source平特性弧焊电源constant –voltage arc wel ding power source下降特性弧焊电源dropping characteristic arc welding power source垂降特性弧焊电源constant-current arc wel ding power source多特性弧焊电源slope-controlled arc weldi ng power source逆变式焊接电源inverter welding power so urce晶体管弧焊电源transistor arc welding po wer source电源动特性dynamic characteristic电源外特性external characteristic弧焊变压器arc welding transformer弧焊整流器arc welding rectifier硅弧焊整流器silicon arc welding rectifier 晶闸管弧焊整流器SCR arc welding rectifie r; arc welding silicon controlled rectifier 脉冲弧焊整流器pulsed arc welding rectifie r弧焊发电机arc welding generator焊车welding tractor焊接机头welding head行走机构traveller送丝机构wire feeder等速送丝方式constant wire-feed system 变速送丝方式alternate wire-feed system跟踪装置tracer焊丝盘wire reel焊钳electrode holder焊枪welding gun电极夹electrode holder导电嘴tip; contact tube喷嘴nozzle焊剂漏斗flux-hopper高频振荡器oscillator; HF unit脉冲引弧器pulsed arc starter; surge injec tor脉冲稳弧器pulsed arc stabilizer脉冲激弧器pulsed arc exciter输出电抗器out put reactor镇定变阻器ballast rheostat直流分量抑制器direct current suppressor 焊接回路welding circuit额定焊接电流rated welding current焊接电流调节范围range of welding curren t regulation空载电压open circuit voltage(no load vol tage)约定负载电压conventional load voltage负载持续率duty cycle额定负载持续率rated duty cycle; standard service手工弧焊机manual arc welding machine 电焊渣electroslag welding (ESW)手工电渣焊manual electroslag welding丝极电渣焊electroslag welding with wire electrode板极电渣焊electroslag welding with plate electrode熔嘴电渣焊electroslag welding with cons umable nozzle管极电渣焊electroslag welding with tube electrode窄间隙电渣焊narrow-gap electroslag weldi ng电渣堆焊electroslag surfacing电渣焊机electrosalg welding machine熔嘴consumable nozzle; consumable wir e钢档板steel shoe (钢冷却板Cu-cooling plate铜滑板copper shoe渣池slag bath渣池深度depth of slag bath渣池电压voltage of slag bath电渣过程稳定性electroslag process stabilit y焊丝间距distance between welding wires 电子束焊electron beam welding (EBW)脉冲电子束焊pulsed electron beam weldin g加速电压acceleration voltage/ operating v oltage电子束电流beam current电子束功率beam power电子束功率密度beam power density焦点focal spot焦距focal length工作距离work distance电子束焊机electron beam welding machin e高真空电子束焊机full vacuum electron beam welder低真空电子束焊机partial vacuum electron beam welder非真空电子束焊机nonvacuum electron bea m welder真空度vacuum电子枪electron gun二极电子枪diode gun三极电子枪triode gun偏压电极bias electrode电磁透镜electromagnetic lens电子束偏转线圈electron beam deflection c oils导流系数perveance钉尖spiking激光焊laser welding/ laser beam welding 连续激光焊continuous laser welding脉冲激光焊impulsed laser welding激光焊机laser welding equipment气体激光器gas laser固体激光器solid laser焦斑直径focussed diameter of the beam离焦量clearance between focal point and (plate) surface焊缝深宽比weld seam depth-to-width ratio3、气焊、热剂焊、水下焊gas welding氧乙炔焊oxy-acetylene welding氢氧焊oxy-hydrogen welding空气乙炔焊air-acetylene welding氧乙炔焊oxy-acetylene flame氢氧焰oxy-hydrogen flame氧煤气焰oxy-coal gas flame焊接火焰welding flame混合比mixing ratio混合气体可燃范围inflammable limit of thegaseous一次燃烧primary combustion二次燃烧secondary combustion燃烧速度combustion rate燃烧强度combustion intensity火焰热效率flame heating efficiency焰芯inner cone; flame cone内焰internal flame外焰flame envelope中性焰neutral flame氧化焰oxidizing flame碳化焰carburizing flame回火flashback逆火backfire回烧flashback气体发生速度gasification speed焊炬torch; blow pipe等压式焊炬balanced pressure torch射吸式焊炬injector torch氧乙炔焊炬oxy-acetylene torch焊割两用炬combined cutting and welding torch混合室mixing chamber喷射器injector焊嘴welding nozzle; welding tip液氧气化器oxygen evaporator气瓶gas cylinder乙炔瓶acetylene cylinder阀罩cylinder cap气瓶阀cylinder valve汇流排cylinder manifold减压器pressure regulator; gas regulator单级减压器single stage regulator两级减压器two stage regulator回火防止器flashback arrestor干式回火防止器dry flashback arrestor水封式回火防止器water-closing type arrestor净化器purifier乙炔发生器acetylene generator低压乙炔发生器low pressure acetylene generator热剂补焊thermit repair welding钢轨热剂焊thermit rail welding。

Multi-step slow annealing perovskitefilms for high performanceplanar perovskite solar cellsLike Huang,Ziyang Hu n,Jie Xu,Ke Zhang,Jing Zhang,Yuejin Zhu nDepartment of Microelectronic Science and Engineering,Ningbo Collabrative Innovation Center of Nonlinear Harzard System of Ocean and Atmosphere,Ningbo University,Ningbo315211,Chinaa r t i c l e i n f oArticle history:Received24February2015Received in revised form31May2015Accepted9June2015Available online25June2015Keywords:Perovskite solar cellsOne-step direct annealingMulti-step slow annealingTightly-distributed performance parametersa b s t r a c tThe morphology,structure,optical and electrical properties of perovskitefilms treated by two differentannealing methods with different annealing temperature ramp and their corresponding device perfor-mance have been studied and compared.Annealing temperature ramp significantly influences the sur-face morphology and optical properties of perovskitefilms which determines the performance of solarcells which determines the performance of solar cells.The perovskitefilms treated by one-step directannealing method tend to exhibit irregular and weak ultraviolet–visibleabsorption spectrum,which caneasily result in great variation in thefinal performance of solar cells.While multi-step slow annealing isbeneficial for preparing highly uniform and well-crystallized perovskitefilms,and thus these devicespresent tightly-distributed performance parameters.The best device treated by multi-step slowannealing method showed a short circuit current density of21.49mA/cm2,an open circuit voltage of0.988V,afill factor of64.86%,and a power conversion efficiency(PCE)of13.58%,which is a57%enhancement of the overall PCE relative to8.65%of the device treated by one-step annealing method.Thesefindings suggest that optimized slow temperature ramp is necessary to prepare high-efficient andwell-reproducible perovskite solar cells.&2015Elsevier B.V.All rights reserved.1.IntroductionHybrid organic/inorganic perovskite(e.g.,CH3NH3PbI3and themixed-halide perovskite CH3NH3PbI3Àx Cl x)based solar cells(PVKSCs)are considered to be economically low cost and have potentialcompetitiveness with any other type of solar cells because of theremarkably high power conversion efficiencies(PCEs)combinedwith simple low-temperature and solution-processed capacity[1–6].Perovskite materials have almost all the excellent semi-conducting characteristics required for solar cell,such as smallband-gaps,high extinction coefficients,high carrier mobility,andlarge charge diffusion length etc.[2,3].Wherein,CH3NH3PbI3Àx Cl xis the preferred materials due to its large carrier lifetime and dif-fusion length that makes it very suitable for planar heterojunctionsolar cells(PH PVK SCs)where the perovskitefilm is sandwichedbetween the hole and electron selective contacts.Currently,inorder to simplify the preparation process of PVK SCs the simpleplanar heterojunction structure is widely adopted[4–6].To realizehigh efficient PH PVK SCs,the morphology of the perovskitefilm isconsidered to be one of the most critical issues[7].There are manyefforts reported to control the morphology of the perovskite toobtain high quality perovskitefilm,including controlling theannealing condition[8–10],adjusting the organic/inorganic pre-cursor ratio[11],adding additives into the precursor solution[12],and solvent engineering[13,14].As important process parameters,annealing temperature and time have a great effect on perovskitefilm morphology thus the structural,electrical and optical prop-erties in perovskite material[7,8,15].For typical one-step solution-processing of perovskitefilm,annealing temperature and time of100°C/45min is adopted since the work reported by Lee et al.[16].For instance,Yang et al.demonstrated a highest efficiency of$19.3%in PH PVK SCs by annealing the perovskite precursorsolution directly at90°C for1h and100°C for25min[6].Meanwhile,Snaith et al.investigated influence of the thermalprocessing protocol upon the crystallization and photovoltaicperformance of perovskites[17].Shen et al.observed a60%increase in PCE for optimized device processed with two-stepthermal annealing method relative to that of the device preparedusing a one-step process(90°C for30min)[18].Kim et al.pro-posed a stepwise ramp annealing method to control the solventevaporation rate to obtain high surface coverage perovskitefilms[19].According to the previously reported results,we can concludethat the quality of the perovskitefilms closely interrelates with theannealing temperature and duration,and hence the efficiency ofContents lists available at ScienceDirectjournal homepage:/locate/solmatSolar Energy Materials&Solar Cells/10.1016/j.solmat.2015.06.0180927-0248/&2015Elsevier B.V.All rights reserved.n Corresponding authors.Tel.:þ8657487600770;fax:þ8657487600744.E-mail addresses:huziyang@(Z.Hu),zhuyuejin@(Y.Zhu).Solar Energy Materials&Solar Cells141(2015)377–382the device has a great varied distribution.In practice,we havefound that the as spin-coated perovskite precursors on the sub-strate can easily generate perovskitefilms with abnormal colorvariation if being annealed directly at temperature of90–100°C.Moreover,such perovskitefilms with abnormal light absorptioncan result in large deviation in thefinal device performance.Thiscan result in a low chance of preparing high quality perovskite films,which will hamper the fabrication of large-area perovskite solar cell modules[20].Therefore,exploring a universal annealingprocess to realize high-efficient and well-reproducible perovskitesolar cells is the main motivation in this paper.In this regard,two different annealing processes including one-step(OS)direct annealing method and multi-step(MS)slowannealing method were introduced to treat the perovskitefilms.By comparing the corresponding morphology,structure,opticaland electrical properties of the treated perovskitefilms,we foundthat the MS method is an universal annealing process to realizehigh-efficient and well-reproducible PH PVK SCs.We demon-strated a maximum PCE of13.58%in MS annealing PVK SCs,accompanied by a57%enhancement of the overall PCE relative to8.65%of the OS annealing device.2.Experimental section2.1.MaterialsLead chloride(PbCl2,99.999%),Diethanolamine(98%),4-tert-Butylpyridine and TiCl4were purchased from Sigma-Aldrich,CH3NH3I from Shanghai Materwin New Materials Co.Ltd.,Titanium(IV)isopropoxide(98þ%)and Li-bis(trifluoromethanesulfonyl)imide(Li-TFSI)from Acros,spiro-OMeTAD from Lumtec,dimethyl-formamide(DMF),acetonitrile,isopropanol,ethanol and chlor-obenzene from Shanghai Chemical Agent Ltd.,China(Analysis puritygrade).All materials were used as received.2.2.Synthesis of CH3NH3I and TiO2nanocrystalsMethyl ammonium iodide(CH3NH3I)was synthesized by theavailable process as reported in literature[6].First,18.7mL(0.15mol)methylamine(Sigma-Aldrich,33wt%in absolute etha-nol)and19.8mL(0.15mol)hydroiodic acid(Sigma-Aldrich,99.99%,57wt%in water)at a1:1equimolar ratio were stirred in anice bath for2h.Then the precipitate was collected by evaporatingat50°C for2h.Finally,a white powder was received by washingwith diethyl ether and ethanol three times and then drying at100°C in a vacuum oven for24h.The TiO2nanocrystals were synthesized from a sol–gel methodin the ambient air[21].Simply,0.675mL of titanium(IV)iso-propoxide was added to18mL of isopropanol and0.25g of die-thanolamine;18μL of deionized water were added before stirring for5min at room temperature,then the sol was left to age for halfan hour before using.2.3.Solar cell device fabricationFTO coated glass substrates(Nippon,14Ω/□)were cleaned with deionized water,ethanol and acetone and then subjected to an ozone-ultraviolet treatment for15min.A40nm thick TiO2compact (c-TiO2)layer was spin-coated on the substrates using the sol–gel solution synthesized above with a spin speed and time of4000rpm/ 25s.The c-TiO2layer was then annealed at450°C for30min.Before using,the substrates were treated in a aqueous solution of TiCl4 (0.04M)at70°C for30min,then rinsed with deionized water and dried at120°C for15min.To deposit the perovskite layer,a1:3ratio of PbCl2/CH3NH3I(0.8M and2.4M)was mixed in DMF.The solution was spin-coated on the FTO/c-TiO2substrates at2000rpm for40–50s and then treated with two different annealing methods descri-bed below.Once perovskite thinfilms grow well,a hole transportlayer(HTL)solution was spin-coated at2800rpm for30s,in which1mL spiro-OMeTAD/chlorobenzene(72.3mg/mL)solution wasemployed with addition of18m L Li-TFSI/acetonitrile(520mg/mL),and29m stly,a120nm thick silver layer wasthermally evaporated on top of the device under a pressure of5Â10À6Torr to form the back contact.Apart from the c-TiO2layer,allfunctional layers were prepared in a nitrogenfilled glove box.Thedevices fabricated have a layered p–i–n configuration glass/FTO/c-TiO2/CH3NH3PbI3Àx Cl x(PVK)/spiro-MeOTAD/Ag,which consists of $450nm thick FTO electrode,$40nm of TiO2,$400nm of per-ovskite,$400nm of spiro-MeOTAD,and$120nm of Ag,as shownin Fig.1(a and b).The schematic diagram for fabricating such a deviceis depicted in Fig.1(c).2.4.Characterizations and measurementsThe UV–vis absorption spectra of the perovskitefilms wererecorded with a VARIAN Cary5000UV–vis–NIR spectrophotometer.The X-ray diffraction(XRD)pattern(2θscans)were obtained from perovskitefilms deposited on the FTO/c-TiO2substrates using aBruker AXS D8Advance X-ray diffractometer using Cu-Kαradiation (λ¼1.54050Å).The top view images and thicknesses of the depos-ited perovskitefilms were confirmed by a Hitachi SU-70scanning electron microscope(SEM).Atomic force microscope(AFM)topo-graphic images of the perovskitefilms deposited on the FTO/c-TiO2 substrates were taken using a Bruker Dimension5000Scanning Probe Microscope(SPM)in tapping mode.Steady-state photo-luminescence spectroscopy(PL)measurements were acquired using an Edinburgh Instruments FLS920fluorescence spectrometer with an excitation wavelength of460nm.The current density–voltage(J–V)measurements were con-ducted under simulated AM1.5G sunlight of100mW/cm2usingan AM1.5G typefilter(Newport,81904,USA).The light intensitywas adjusted by using a standard Si cell.J–V curves were obtainedby applying an external bias to the cell,and measurements wererecorded with a Keithley model2400digital source meter at roomtemperature in the ambient air.The incident photo-current con-version efficiency(IPCE)spectrum was measured by a IPCE system(Newport2936-c power meter)in the300–800-nm wavelengthrange under the irradiation of a300W xenon light source with anOriel Cornerstone2601/4monochromator in DC mode at roomtemperature.The effective area of the cell was defined to be0.07cm2by a non-reflective metal mask.3.Results and discussionTwo different annealing processes,OS and MS,were adoptedfor perovskitefilm treatment in fabrication of PH PVK SC devicesas shown in Fig.1.The insets in Fig.1(d)correspond to0min(1),1min(2),10min(3),20min(4),30min(5),45min(6)whilethose of Fig.1(e)correspond to0min(1),15min(2),40min(3),60min(4),80min(5),105min(6).Generally,a transparent yel-lowishfilm forms originally once deposition of the precursorsolution on the FTO/c-TiO2substrate is completed.Leaving thefilmat room temperature for several minutes,the color changes to red,then deep yellow andfinally deep black[22].This change in colorcorresponds to the solvent evaporation,reaction of the precursorand sublimation of excess organic CH3NH3Cl as reaction byproduct[7].Obviously,as shown in Fig.1(d,e),the color change of per-ovskitefilms treated with the OS method is faster than that ofperovskitefilms treated with the MS method.It is reasonable thatprecursor reaction and byproducts sublimation occurred fasterL.Huang et al./Solar Energy Materials&Solar Cells141(2015)377–382 378due to elevated temperature since the beginning of annealing.Because the melting point of precursor CH 3NH 3I (about 70°C)is signi ficantly lower than the boiling point of DMF (about 150°C),rapid annealing such as the OS method may cause too fast and uncontrollable evaporation of solvent accompanied with CH 3NH 3I,which may result in reaction stoichiometry ratio mismatch of the0102030405060708090100102030405060708090100Time (min)Time (min)T e m p e r a t u r e C )654321Fig.1.Device structure (a),scanning electron microscope cross-sectional image of the device (b),and process steps for the fabrication of the perovskite solar cells (c),and two different annealing methods adopted for the treatment of perovskite films (d),OS and MS methods (e).The insets are the photographs of the perovskite films.Fig.2.(a)UV –vis absorption spectra of two representative perovskite films annealed with different methods,(b)XRD patterns for perovskite films annealed with different methods,(c)the photovoltaic performance of their corresponding devices,and (d)the corresponding incident photo-current conversion ef ficiency spectra.L.Huang et al./Solar Energy Materials &Solar Cells 141(2015)377–382379precursor.Actually,we didfind that the use of OS method is dif-ficult to obtain high quality perovskitefilms compared with the MS method.The color of a typical high quality perovskitefilm is reddish-brown as shown in the inset of Fig.1(e),which is also reported by the previous reports[6,7,22].However,in our experiments thefilms treated with the OS method exhibit grayish-purple or grayish-green as shown in the inset of Fig.1(d).Even under a prolonged annealing time,the color was not changed.We characterized thefilms prepared using the two different methods by UV–vis spectroscopy as shown in Fig.2(a).Compared to the absorption spectrum curves of perovskitefilms treated with the MS method,the absorption curve shape and intensity of the per-ovskitefilms treated by the OS method is irregular and weak.The perovskitefilms exhibit the760nm absorption peak corresponding to the direct band gap transition from thefirst valence band(VB1)to conduction band(CB)[16],regardless of the different annealing methods adopted.While the absorption spectrum of these grayish-purple and grayish-green perovskitefilms treated by the OS method scarcely contains the480nm absorption peak corresponding to the transition from the second valence band(VB2)to CB[3].Note that the absorption spectra curves of these grayish-purple and grayish-green perovskitefilms are very similar with those of the initial annealingfilms(annealed at95°C/65min)with the MS method.The stable non-red-brown color regardless of annealing time implies that thesefilms suffered irreversible changes at a direct annealing tem-perature of95°C,possibly resulting from too fast CH3NH3I eva-poration accompanied with the solvent DMF.Since light absorption is thefirst step for a solar cell as a photoelectric conversion device,one can anticipate that such a great variation in UV–vis absorption can lead to significant deviation in thefinal device performance.However,the XRD measurements(Fig.2(b))do not show any difference in diffraction peaks position between the perovskite films treated by the two annealing methods.The peaks at14.2°, 28.5°,43.3°and59.0°are assigned to the(110),(220),(330)and (440)planes[6,16],respectively,similar to those reported for the tetragonal perovskite structured CH3NH3PbI3crystals.While the intensity of each diffraction peak of the perovskitefilms treated by the OS method is weaker than that of perovskitefilms treated by the MS method and no new peaks or peak shifts were observed. This reveals that the crystal structure of the two differently annealed perovskitefilms is the same and the only difference between them is the degree of crystallinity.The slow solvent evaporation in MS method facilitates atomic rearrangement which leads to increase in crystalline fraction and hence the enhanced peak intensities.According to the different light absorption spectra and crys-tallinity of thefilms treated by the two methods,we can expect different photovoltaic performance.Fig.2(c)shows the current density–voltage(J–V)measurements of their corresponding PVK SC devices.For device OS,it contains perovskitefilm treated by the OS method,thefilm is abnormally grayish-purple,the typical device shows PCE¼3.60%,J SC¼8.42mA/cm2,V OC¼0.847V,and FF¼49.63%.For the device MS,it contains perovskitefilm treated by the MS method,thefilm is normally reddish-brown,the typical device shows PCE¼10.14%,J SC¼17.56mA/cm2,V OC¼0.931V,and FF¼60.97%.The dark J–V curve of a solar cell device generally reflects its rectifying characteristics as a pared to the device OS,the device MS behaved as well defined diodes with better rectifying characteristics as its dark J–V curve exhibits larger turn-on voltage.Additional,for the device OS,although the per-ovskitefilm UV–vis absorption spectrum contains only one absorption peak of760nm,it still has a large V OC of0.847V,as if the V OC is decided only by thefirst valence band(VB1)and is independent of the second valence band(VB2),regardless of other factors that affect V OC[23].Fig.2(d)presents the incident photo-current conversion efficiency(IPCE)spectra.Taking into account the fact that the smaller IPCE value of the device OS relative to that of the device MS in the entire visible wavelength,we can conclude that these grayish-purple and grayish-green perovskitefilms are not the desirablefilms for high performance.We further examined the influence of the different annealing processing on the morphology and structure of the perovskite layer formed on the c-TiO2layer by SEM as shown in Fig.3.It is believed that for efficient PVK SCs,aflat,highly uniform,pin-hole-free and high surface coverage CH3NH3PbI3Àx Cl xfilm is of extreme significance,as required in other thinfilm solar cells.We can see that thefilms produced by the OS method(Fig.3(a))contain dif-ferent size grains with an incomplete coverage on the substrate, while thefilms produced by the MS method(Fig.3(b))are very uniform.A pinhole with a feature size of$2m m has been high-lighted in Fig.3(a).It should be noted that the presence of pinholes thus the lower coverage of the perovskitefilm could have a sig-nificant impact on thefinal UV–vis absorption spectra measured, as light may go directly through the pinholes without being absorbed by perovskite.This may be one reason that the above measured perovskite thinfilms prepared by the OS method showed low absorption intensity.Otherwise,we can expect that the coverage of perovskite would have a significant effect on J SC of thefinal devices.It is the reason that the device MS presents the double J SC of the device OS(Fig.2(c)).In addition,we notice that many small particulates are gathered at the step edges of the layered perovskite crystals(Fig.3S).We suppose that these par-ticulates were precipitated from the perovskitefilms because of the rapid annealing process.Actually,we can reduce these pin-holes and particles by surface solvent treatment[14].The surface morphology of the perovskitefilms was also char-acterized by the AFM measurement as shown in Fig.4(a,b).The AFM images(8Â8μm2)show that thefilms produced by theMS Fig.3.SEM images for perovskitefilms annealed with OS method(a),MS method(b).The insets are high-magnification SEM images.L.Huang et al./Solar Energy Materials&Solar Cells141(2015)377–382380method has low roughness with a root-mean-square (RMS)of 35nm while the film prepared by the OS method has a high roughness with a RMS of 72nm.Consistent with the SEM obser-vations,the crystal arrangement in the film produced by the MS method is more compact as no pinholes can be found,which ensures suf ficient light absorption and avoids possible short-circuit occurrence.It is believed that high temperature and rapid annealing treatment can result in the rapid growth of perovskite crystals and lead to the formation of large crystalline islands and the associated large gaps.Lower temperature and slow annealing treatment allow the perovskite crystals to grow slowly and uniformly from a large number of nucleation sites and result in uniform crystal structure with a small number of internal voids or pinholes.Thus,slow annealing may be necessary to fabricate highly uniform perovskite films without pinholes.The film quality may also have a great impact on the electron transport properties of perovskite film.As shown by SEM images,these two films contain varying degrees of surface defects which will have effects on the device performance.In order to compare the crystal defects inside the films,steady-state photoluminescence spectroscopy (PL)was conducted as shown in Fig.4(c).The enhanced PL spectroscopy of the perovskite films treated by the MS method con firms the improved crystallinity of the perovskite films.This can be attributed to the number of crystal defects is well reduced,therefore the non-radiative recombination paths are greatly elimi-nated.Thus,we can anticipate that the carrier can be collected by the corresponding selective contacts ef ficiently without too much recombination induced by crystal defects .The perovskite films treated by the OS method exhibit weak fluorescence intensity,which suggests that these films contains many physical defects or bulk traps that act as carrier recombination centers.Fig.5(a)displays the histograms of PCE of the 48devices from 4different batches.It is shown that the PCEs were distributed within 8–14%for the MS method processed PVK SCs and 0–9%for the OS method processed PVK SCs,respectively.Obviously the PCE values of the former exhibits a narrow distribution with a small deviation about 12%from the average values.While the PCE values of the later show a relatively wide distribution with a large deviation of about 48%from the average values.The origin of this variation in PCE values of the device OS here is primarily attributed to the abnormal characteristics of the perovskite films that are treated by the OS methods.Furthermore,the distribution of PCE values of the solar cells was fabricated using the MS method is fairly narrow compared to that of the OS method processed PVK SCs.Finally,we compare the stability of the two kinds of devices treated with different annealing methods.The tested OS and MS devices without encapsulation were stored at a glove box full with dry nitrogen and tested outside at every 24h for 7days.These results are shown in Fig.5(b).The devices OS retain 20–25%of the initial performance after 72h,and 10–15%after 7days,respec-tively.While the devices MS retains more than 60%of the initial performance after 72h,and 30–35%after 7days,respectively.Here,the enhanced stability of the devices MS was attributed to the slower decomposition of the MS annealing perovskite films induced by fewer molecular defects as the PL measurement reveals.The growth of a more stable perovskite material is of virtual signi ficance for photovoltaic applications.Even in the absence of humidity,a decomposition of the perovskite structure can take place through the statistical formation of molecular defects with a non-ionic character,whose volatility at surfaces should break the thermodynamic defect equilibriums [24].The strategies that can substantially prolong the lifetime of the material were also reported [14,25,26].Here,by comparingtheFig.4.AFM images for perovskite films annealed with OS method (a),MS method (b),and PL spectroscopy of the perovskite films(c).Fig.5.(a)Histograms of the photoelectric conversion performance of 48devices with perovskite films annealed with two different methods from 4different batches,and (b)stability investigation on the unencapsulated devices OS and MS as a function of time.L.Huang et al./Solar Energy Materials &Solar Cells 141(2015)377–382381device performance with the perovskitefilms treated by two dif-ferent methods wefind the origin of the variation in the perfor-mance,which will result in better reproducibility from solution processes without any other sophisticated optimization.With the high efficiency and stability,we can expect efficient and repro-ducible preparation of PVK SC modules.4.ConclusionsIn conclusion,the morphology,structure,optical and electrical properties of perovskitefilms and devices performance with per-ovskitefilms treated by two different annealing methods have been compared.The annealing temperature ramp significantly influences the surface morphology and optical properties of per-ovskitefilms thus the performance of solar cells.The optimized annealing method ensures reproducible preparation of perovskite thinfilms with high crystal quality and good stability.The best performance of the PVK SC treated by the MS method showed a PCE of13.58%,a J SC of21.49mA/cm2,a V OC of0.988V,and an FF of 64.86%.The optimized annealing method is not only suitable for the preparation of highly efficient PVK SCs but also suitable for the preparation of other perovskite based device such as light-emit-ting diodes,lasers,photodetectors,thinfilm transistors and so on. AcknowledgmentsThis work was supported by the National Science Foundation of China(Grant nos.11374168,11304170,and51302137),the Natural Science Foundation of Zhejiang Province(Grant no.LQ13F050007), and the K.C.Wong Magna Fund in Ningbo University.Appendix A.Supplementary materialSupplementary material associated with this article can be found in the online version at /10.1016/j.solmat. 2015.06.018.References[1]P.Gao,M.Grätzel,M.K.Nazeeruddin,Organohalide lead perovskites for pho-tovoltaic applications,Energy Environ.Sci.7(2014)2448–2463.[2]S.D.Stranks,G.E.Eperson,G.Grancini,C.Menelaou,M.J.P.Alcocer,T.Leijtens,L.M.Herz,A.Petrozza,H.J.Snaith,Electron–hole diffusion lengths exceeding 1μm in organometal trihalide perovskite absorber,Science342(2013)341–344.[3]G.C.Xing,N.Mathews,S.Y.Sun,S.S.Lim,m,M.Gratzel,S.Mhaisalkar,T.C.Sum,Long-range balanced electron-and hole-transport lengths in organic–inorganic CH3NH3PbI3,Science342(2013)344–347.[4]M.Liu,M.B.Johnston,H.J.Snaith,Efficient planar heterojunction perovskitesolar cells by vapour deposition,Nature501(2013)395–398.[5]D.Liu,T.L.Kelly,Perovskite solar cells with a planar heterojunction structureprepared using room-temperature solution processing techniques,Nat.Pho-tonics8(2013)133–138.[6]H.P.Zhou,Q.Chen,G.Li,S.Luo,T.B.Song,H.S.Duan,Z.R.Hong,J.B.You,Y.S.Liu,Y.Yang,Interface engineering of highly efficient perovskite solar cells,Science 345(2014)542–546.[7]G.E.Eperon,V.M.Burlakov,P.Docampo,A.Goriely,H.J.Snaith,Morphologicalcontrol for high performance,solution-processed planar heterojunction per-ovskite solar cells,Adv.Funct.Mater.24(2014)151–157.[8]A.Dualeh,N.Tétreault,T.Moehl,P.Gao,M.K.Nazeeruddin,M.Grätzel,Effect ofannealing temperature onfilm morphology of organic–inorganic hybrid per-voskite solid-state solar cells,Adv.Funct.Mater.24(2014)3250–3258.[9]F.X.Xie,D.Zhang,H.Su,X.Ren,K.S.Wong,M.Gratzel,W.C.H.Choy,Vacuum-assisted thermal annealing of CH3NH3PbI3for highly stable and efficient per-ovskite solar cells,ACS Nano9(2015)639–646.[10]J.Xiao,Y.Yang,X.Xu,J.Shi,L.Zhu,S.Lv,H.Wu,Y.Luo,D.Li,Q.Meng,Pressure-assisted CH3NH3PbI3morphology reconstruction to improve the high perfor-mance of perovskite solar cells,J.Mater.Chem.A3(2015)5289–5293. 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Appl.Phys.A74,19–25(2002)/Digital Object Identifier(DOI)10.1007/s003390100893Applied Physics AMaterialsScience&ProcessingFemtosecond laser ablation of silicon–modification thresholds and morphologyJ.Bonse∗,S.Baudach,J.Krüger,W.Kautek,M.LenznerLaboratory for Thin Film Technology,Federal Institute for Materials Research and Testing(BAM),Unter den Eichen87,12205Berlin,Germany Received:4December2000/Revised version:29March2001/Published online:20June2001– Springer-Verlag2001Abstract.We investigated the initial modification and abla-tion of crystalline silicon with single and multiple Ti:sapphire laser pulses of5to400fs duration.In accordance with earlier established models,we found the phenomena amorphization, melting,re-crystallization,nucleated vaporization,and abla-tion to occur with increasing laserfluence down to the short-est pulse durations.We noticed new morphological features (bubbles)as well as familiar ones(ripples,columns).A nearly constant ablation thresholdfluence on the order of0.2J/cm2 for all pulse durations and multiple-pulse irradiation was ob-served.For a duration of≈100fs,significant incubation can be observed,whereas for5fs pulses,the ablation threshold does not depend on the pulse number within the experimental error.For micromachining of silicon,a pulse duration of less than500fs is not advantageous.PACS:79.20D;42.70.QMicromachining with ultrashort laser pulses has attracted growing interest even in industry and medicine since the ap-propriate lasers were made readily available for a wide set of parameters[1,2].It has been demonstrated that ultrashort pulses bear the potential for precise micromachining(later-ally and vertically)in transparent dielectrics[3].In the course of investigations with femtosecond pulses,it became obvious that the detailed mechanisms of damage to solids caused by laser light are far from fully understood.A number of phenomena concerning photo-induced mod-ification of silicon surfaces have been explored in different ranges of wavelength,intensity and duration of the applied laser pulses.In this paper,we want to extend the existing investigations on laser-induced surface damage in silicon to pulse durations as short as5fs.We also observed several different phenomena;we try to methodically“file”these ob-servations into a physical overview.We will demonstrate that the so-far-assumed sequence of physical processes,namely amorphization[4],melting[5, 6],re-crystallization[4,7],nucleated vaporization[8],andfi-nally ablation[9],can also account for these experimental re-∗Corresponding author.(Fax:+49-30/8104-1827,E-mail:joern.bonse@bam.de)sults.Various well-known features,for example,ripples[10] and columns[11],could be realized and appropriately ex-plained as well.In Sect.1,the current knowledge about the interaction between laser pulses and silicon is reviewed. Our experimental results are shown and compared to this in Sect.2.1Physical considerationsThe deposition of the laser energy into a solid is usually viewed in the quantum-mechanical formalism of particle in-teraction.The incident pulse energy is absorbed by the elec-trons,dependent on the peak intensity,by one-,two-or more-photon absorption.Absorption by free carriers(sometimes called inverse bremsstrahlung)depends on the number of al-ready existing carriers and is therefore a subsequent process. The same applies to collisional ionization,which utilizes part of the energy of highly excited carriers to generate new free electrons.These carriers then thermalize to a Fermi–Dirac distribution while transferring their excess energy to phonons, typically on a time scale of100fs.These phonons afterwards recombine to a Bose–Einstein distribution in a few picosec-onds[12].During the detailed exploitation of pulsed-laser annealing(PLA,typically done with nanosecond pulses),a “plasma-annealing”model was established,which stated that a non-thermal“bond softening”was responsible for the loss of the crystal structure[13,14].Recently,this non-thermal model was shown to be applicable for several semiconductors irradiated with femtosecond pulses[15–18].So far,no spatial transport of energy out of the excitation region has been considered.In order to treat the subsequent processes,including melting,boiling,and ablation of mate-rial,one usually uses either a two-temperature model[19,20], which distinguishes between electron and lattice(ion)tem-peratures,or a completely classical model of thermal trans-port in a continuum[8,21].The latter one describes phase changes from the molten phase to a gas,considering the ex-istence of transient thermodynamical states(such as super-heated liquids)due to the rapid action of the ultrashort laser pulses.The physical mechanisms that are involved in photo-excitation of the solid are manifested also in irreversible20changes of the irradiated surface.These changes can be used for identification of some of the processes and also for deter-mination of their thresholdfluences.After irradiation with short laser pulses,re-solidification of molten material was observed to happen in two stages: amorphization and re-crystallization[4].The difference was simply attributed to the amount of energy deposited in the ma-terial(the temperature)and the consequent cooling velocity. At lower temperatures,the material has not enough time to re-crystallize from the melt,thus leaving the semiconductor in an amorphous state.In regions with higher temperatures, cooling is sufficiently slow to allow re-crystallization.Already in previous experiments,a rather mysterious phe-nomenon has been discovered after the solids have been ir-radiated with multiple subsequent pulses[22].Finally termed “ripples”,these periodic surface structures appear as lines orthogonal to the direction of the electricfield vector of the incident light and show a period on the order of the wave-length of the generating light[10,23].The generally accepted explanation of these ripples is an interference between the in-cident light and a surface wave(generated by scattering).This interference leads to periodic modulation of the absorbed in-tensity and consequently to modulated ablation.Column formation in crystalline silicon as a result of multi-pulse laser irradiation has been observed in the past at different laser wavelengths(UV–NIR),for different pulse durations(fs–ns),and in different environments(vacuum, air,different gases).A certain number of laser pulses is re-quired to initiate the self-organized growth process of Si microcolumns in the irradiated region.This phenomenon is of major importance because it can limit the precision of laser ablation.For the treatment with ultrashort(fs) laser pulses,the Si-column formation was observed by sev-eral groups under different experimental conditions(λ= 248nm,τ=105fs,vacuum[24];λ=390nm,τ=250fs, vacuum[9];λ=620nm,τ=300fs,air[25];λ=780nm,τ=100fs,SF6,Cl2,N2,He,vacuum[11]).The phenomenon was also found for short-pulse(ns)excimer-laser irradiation (λ=193nm,τ=23ns,air[26];λ=248nm,τ=25ns, SF6,N2,O2,Ar[27];λ=248nm,τ=12ns,vacuum[28];λ=308nm,τ=28ns,vacuum[29]).The process strongly depends on the number of pulses ap-plied to the same spot and the laserfluence.A further key parameter for the formation process and the shape of the microstructures seems to be the ambient environment.Ox-idizing or halogen-containing atmospheres such as air,O2 or SF6support the generation of high-aspect-ratio pillars, whereas the formation of sharp spikes can be reduced in vac-uum,N2or He[11].On the other hand,column formation is rather insensitive to the laser wavelength[9,11,24,25]and the pulse duration[30,31].Influences of the doping concen-tration have not been observed[11,27].For these reasons, a chemical control of the dimensions of microcolumns seems to be possible[31].2Experimental results and analysisExperiments were carried out with two different Ti:sap-phire laser systems,a commercial CPA system(SPECTRA PHYSICS,Spitfire)at the BAM Berlin and the Vienna sys-tem,comprised of an amplifier and hollowfiber compressor,which is capable of producing5-fs pulses with a maximum energy of500µJ[32].The pulse duration of the latter one was changed between5fs and400fs by inserting dispersive material(glass blocks)in the beam path.The experimen-tal conditions were kept similar.The center wavelengths of the linearly polarized laser pulses differed by only20nm (BAM:800nm,Vienna:780nm).The different repetition rates(BAM:10Hz,Vienna:1kHz)should have no influence on the experimental results because every physical process known to be important here is terminated after1ms.An im-portant measurement–actually the one that dominates the overall error–is the energy detection.Here we used a py-roelectric detector BESTEC PM200(BAM)and the OPHIR pulse energy detector NOV A(Vienna),respectively.Different pulse numbers with varying energy were fo-cused to a diameter on the order of several10µm(BAM: f=60mm plano-convex lens,Vienna:R=100mm spher-ical silver mirror)onto the polished(111)surface of n-doped silicon samples.On these samples,a native oxide layer of about2.7nm thickness has been found from ellip-sometric measurements.For higher appliedfluences(in the single-pulse case for5-fs pulses),the sample was placed in a slightly evacuated chamber(p≈10−4mbar)in order to pre-vent ionization or non-linear effects in air and resulting pulse distortions.Inspection of the irradiated surface regions was performed using an optical microscope(Reichert–Jung,Polyvar)in No-marski mode.A more detailed characterization of morpho-logical changes of the laser-modified areas was done by means of a scanning electron microscope(SEM)equipped with a cold-field electron emission cathode(Hitachi,S-4100, accelerating voltage10kV)and an atomic force microscope (AFM,Digital Instruments,Dimension3000SPM)operated in tapping mode.Anticipating the results of our investigations,we outline the principal physical processes occurring on the Si surface after a Gaussian laser pulse was incident in Fig.1.For com-parison,a damage spot on the silicon surface generated by a single laser pulse is shown in Fig.2exhibiting different cir-cular regions of modification,annealing,and ablation.The formation of ripples cannot be seen in this picture because it only occurs after irradiation with multiple pulses onto the same sample spot.In the following section,these thresholds will be further investigated and classified quantitatively.2.1Modification thresholdsSome of the early experiments on laser-induced modification of silicon surfaces distinguish regions of amorphization and crystallization[4].We observed the same phenomena in our experiments,but the zone of amorphization showed a fur-ther substructure which we believe is related to oxidation of the surface layers of silicon.The thresholds of oxidation and amorphization are so close together that unambiguous iden-tification is hardly possible.However,in order to take this fact into account,we call the physical process in this region modification rather than amorphization.The thresholdfluences for these phenomena can be de-termined similar to the ablation thresholdfluence,namely measuring the diameter of the modified areas versus the pulse fluence and extrapolating to zero[33].In Fig.3,the square of21/xFig.1.Physical processes during the modification of silicon with femtosec-ond laser pulses and their threshold fluencesFig.2.Nomarski optical micrograph of the silicon sample surface treated with a single laser pulse in air (λ=800nm,τ=130fs,Φ0=1.5J /cm 2).The outermost ring has a diameter of 45µmthe diameter (corresponding to a modified area)is depicted versus increasing peak fluence of the laser pulses.Extending the regression of this line to zero yields the threshold values to Φmod =0.26J /cm 2and Φann =0.55J /cm 2,respectively.Forthe applied pulse duration,this is identical to the single-pulse threshold measured by Pronko et al.[20].The ablation thresholds of multi-shot experiments in air for different pulse durations are shown in Fig.4.For pulse durations below 100fs,the threshold becomes constant,a be-havior that is well known for metals [34].For higher pulse numbers,one can find no more evidence for crystallization or oxidation/amorphization.A clean edge of ablation as in Fig.9a can be recognized.From the dimen-sions of these craters,an ablation threshold is determined which cannot be distinguished from other thresholds due to morphological changes in the irradiated surface region.The values in Fig.4are significantly lower than the single-pulse thresholds evaluated from Fig.3,because the thresholds of modification and ablation depend on the number of applied laser pulses.This incubation effect rests on a non-ablating modification of the sample material by the laser pulses in such a manner that the threshold for damage decreases.This effect has been extensively studied at the surface of single-crystal metals [35].A dependence in the form of a power law was found:Φmod (N )=Φmod (1)·N ξ−1.(1)Φmod (N )denotes the modification threshold fluence for N laser pulses,and ξis a material-dependent coefficient.In-cubation is related to an accumulation of energy (i.e.non-complete dissipation of the deposited energy)into plastic stress–strain of the metal.However,this formula has also suc-cessfully been employed in the case of indium phosphide (InP)[36],where it is unclear whether intermediate storage of laser energy is mechanical or,for example,chemical (as in several glasses by F-center formation [37]).In Fig.5,the dependence of N ·Φmod (N )on the number of pulses is plot-ted for our data.The fit according to (1)(solid line)yields a coefficient ξof 0.84.From Fig.5,one can conclude that there is significant in-cubation in silicon for pulses with a duration of ≈100fs.Laser fluence Φ0[ J/cm 2]1S q u a r e d d i a m e t e r D 2[µm 2]1000200030000.50.35Fig.3.Diameter (squared)of modification and re-crystallization of the sili-con surface versus the incident peak fluence of the laser pulse (λ=800nm,τ=130fs,N =1,in air).Squares belong to the areas of modification,whereas circles belong to the re-crystallization regions.Solid lines are lin-ear regressions within the semi-logarithmic plot.The deviation of the data from the regression for high fluences is attributed to a slightly non-Gaussian beam profile (caused by apertures)22T h r e s h o l d f l u e n c e [J /c m 2]Fig.4.Ablation threshold fluence of n-Si(111)for several pulse durations,100pulses per spot,in air.Values measured at λ=780nm,except the solid circle (λ=800nm)Number of pulses N110100N ∗Φm o d (N ) [ J /c m 2]110Fig.5.Threshold fluence of laser-induced damage of silicon versus num-ber of laser pulses with a duration of τ=130fs and λ=800nm in an air environment.The solid line represents a least-squares-fit with (1),where ξ=0.84Fig.6.AFM picture of damage in silicon generated with a single Ti:sapphire laser pulse (λ=780nm,τ=5fs,Φ0=7.7J /cm 2).Dark areas indicate more ablated material.The inset at the bottom of the picture is a line-scan along the dotted white line ,the depth scale is indicated in blackThe precise nature of this effect,whether energy is stored in the form of chemical modification or by mechanical stress (as in the case of metals),cannot be deduced from these results.Interestingly,single-shot measurements with 5-fs pulses yield a damage threshold of 0.20±0.05J /cm 2,which agrees with the threshold achieved with multiple pulses within the experimental error (compare Fig.4).Obviously –for these short pulses –there is no such intermediate storage of energy below the damage threshold as it was found,for example,in fused silica [38].2.2Single-pulse experimentsSurface images taken with an atomic force microscope (AFM)and a scanning electron microscope (SEM)reveal in-teresting morphological features of the damaged areas.TheFig.7a,b.SEM picture (0◦)of damage in silicon generated with Ti:sap-phire laser pulses in air (λ=780nm,τ=5fs,Φ0=2.5J /cm 2,N =5).Three different regions of modification (ablation including ripples,re-crystallization,and amorphization)can be recognized.a Full view,b detail23formation of circular substructures (holes)within the cavities can be observed (see Fig.6).These holes vanish or are ob-scured by other morphological features when the same spot is illuminated with subsequent pulses.With increasing laser fluence,the size of these holes increases.Phenomena such as these are frequently attributed to a locally enhanced car-rier density generated either by an inhomogeneous laser beam profile or by locally enhanced absorption (scratches,crystal defects,dust).An initialization of inhomogeneous surface structures due to “hot spots”in the beam profile can be ruled out because –due to the efficient spatial filtering by guiding in a hollow fiber –the Vienna system exhibits an extremely smooth beam profile [32].External surface impurities (dust,scratches due to pol-ishing)cannot be significant,as we will see in the follow-ing argument.We consider indirect two-photon absorption with a coefficient of only 1cm /GW [39]as the domin-ant carrier-generating mechanism.Calculating the penetra-tion depth induced by this mechanism,one finds that the number of absorbing atoms in the excited volume is far smaller than the number of photons supposedly absorbed in this volume.Thus,even the indirect two-photon absorp-tion is already strongly saturated.Virtually all available electrons are excited and it is hardly conceivable that the carrier density exhibits local spikes (e.g.by absorption of defects)so distinct that locally enhanced ablation could occur.Although an enhancement of surface absorption is no appropriate explanation for the observed substructures,en-hancement of absorption at depth in the semiconductor (where the light intensity already dropped one or moreordersFig.9a–f.SEM pictures (60◦)of damage in silicon generated with Ti:sapphire laser pulses in air.a Φ0=1.0J /cm 2,b 1.3J /cm 2,c 1.8J /cm 2(λ=780nm,τ=100fs,N =100).d Φ0=2.0J /cm 2,e 2.8J /cm 2,f 4.1J /cm 2(λ=800nm,τ=130fs,N =100)of magnitude)could account for an evolving inhomogeneous energy deposition.Consequently,after the strongly saturated and overheated surface layer was removed by phase explosion,normal boil-ing including inhomogeneous nucleation of bubbles occurs in the remaining liquid layer [21].This scenario is sup-ported by the fact that larger bubbles are formed in regions of higher fluences,i.e.regions of higher temperature (and therefore slower cooling)where bubbles have more time togrow.Fig.8.SEM picture (0◦)of damage in silicon generated with Ti:sapphire laser pulses in air (λ=800nm,τ=130fs,Φ0=0.42J /cm 2,N =5)242.3Ablation with multiple pulsesThe application of a moderate number(N≈5)of laser pulses leads to characteristic laser-induced periodic surface struc-tures(ripples).In single-pulse experiments,these highly ori-ented structures were not observed,indicating that a feed-back mechanism is involved during the formation of the sur-face patterns.Fig.7shows typical surface damage in silicon(λ=800nm,τ=5fs)at afluence of2.5J/cm2.Three dif-ferent modified zones are clearly visible(compare Fig.1): ablation and ripple-formation in the central region,anneal-ing in thefirst annular structure,and modification in the outer annular border.It is interesting to note,that all these surface modifications known from longer pulses also occur at this ul-trashort pulse duration of5fs.A magnified view(Fig.7b) reveals average lateral ripple periods between650nm and 750nm which is comparable to the laser wavelength.The rip-ples were always oriented perpendicular to the electric-field vector of the incident radiation.Thus,we attribute this phe-nomenon to the well-known mechanism of interference andsubsequent localfield enhancement[10].Small globules of re-deposited material were observed on the top of the surface corrugations.The same characteristic ripple morphology was detected in the central crater region at an≈25times longer pulse duration(λ=800nm,τ=130fs,Φ0=0.42J/cm2,N=5, see Fig.8).Additionally,some outspread periodic patterned(triangular)regions are seen in the direction of the electric field.A further increased number of laser pulses(N≈100) leads to another characteristic surface morphology:the columns or pillars,already introduced in Sect.1.A certain pulse number is required to nucleate the column growth pro-cess.The evolution of silicon microcones and mirocolumns in a series of laser-generated craters,obtained with a con-stant number of100Ti:sapphire laser pulses(τ=100fs at λ=780nm,andτ=130fs atλ=800nm)at varying peak fluences in air is shown in Fig.9.At a comparatively lowfluence of1.0J/cm2(which is ≈5−6times above the ablation threshold),a uniformly ablated crater with a rough,but featureless bottom can be seen as well as highly directed nearly wavelength-sized rip-ple structures in the border region(Fig.9a).With increas-ing laserfluence,small conical structures arise from the bottom of the craters to form the initial stages of micro-columns(Fig.9b,c).The lateral and vertical extent of the columns and the spacing between them strongly depends on the localfluence.In the center of the irradiated area, the columns are wider,taller and more sparse.In the bor-der region they are packed closer together.Up to afluence of≈2J/cm2,the columns are formed in the middle of the crater(Fig.9d),while at higherfluences(Φ0=2.8J/cm2) the morphology appears crown-like.At this stage of devel-opment,the columns can protrude above the original surface plane(Fig.9e),which provides conclusive evidence for the redeposition/re-crystallization origin of these columns.At further increased laserfluences ofΦ0≈4.1J/cm2,a volcano-like structure is observed within the ablated region(Fig.9f). It is probably formed by not completely ejected mate-rial,which is redeposited at the crater walls when the crater depth exceeds a certain value.The height of the columns grows with an increasing number of laser pulses.If a critical size is reached,a destruction of the Si pillars occurs[24].Concerning the formation mechanism of the silicon columns,we suggest a similar explanation as Lowndes et al.[31].Initial surface corrugation inhomogeneously nucle-ates from local vaporization(bubble ejection from the melt layer)and/or ripple formation and subsequently re-deposited material.On the edges of these corrugations,the absorbed local laserfluence is reduced due to an altered angle of inci-dence of the laser radiation.Therefore,ablation takes place preferably at the minima and maxima of the surface topog-raphy.The silicon-rich vapor which is formed at the grooves cools during the material transport(expansion of the vapor plume)and can be re-deposited at the protruding features of the surface.During a large number of these transport cycles,a highly protruding column can be formed.Addi-tionally,the effect can be enhanced by multiple reflections of the incident laser radiation on the bodies of thecolumns, Fig.10a,b.Cross-sectional SEM picture of damage in silicon gener-ated with Ti:sapphire laser pulses in air(λ=800nm,τ=130fs,Φ0= 0.65J/cm2,N=500).a Full view,b detail25Fig.11.Scheme of the different morphological phenomena after irradiation of the silicon surface with linearly polarized femtosecond laser pulses of typically 100fs durationwhich “guides”the light into the grooves.Therefore,the re-gions between the columns again act as emitters of ablated material.A cross-section through a crater (depth ≈9µm)in silicon obtained after the application of 500subsequent laser pulses in air (λ=800nm,τ=130fs,Φ0=0.65J /cm 2)is shown in Fig.10a.A detail of the crater wall can be seen in Fig.10b.Besides an irregular surface morphology and remaining parts of small columns,only a thin thermally or chemically modi-fied layer (depth <500nm)is visible.Figure 11summarizes the different morphological fea-tures (bubbles,ripples,microcolumns)formed after irradi-ation of silicon surfaces with linearly polarized laser pulses for pulse durations of approximately 100fs.3ConclusionWe investigated laser-induced modification and ablation of silicon surfaces with laser pulse durations in the range be-tween 5fs and 400fs.The multi-pulse ablation threshold flu-ence is almost constant around 0.2J /cm 2.We found several physical processes resulting in clearly distinguishable mor-phological features.These are (from lower to higher fluences)oxidation,amorphization,re-crystallization,the formation of bubbles due to boiling below the surface,and finally ablation.Other features occur while treating the sample with multiple subsequent pulses,namely ripple formation,column growth,and crater formation due to material removal.Although these phenomena can limit the precision of 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