Cyclic crack growth tests with CRB specimens for the evaluation of the long-term performance of PE p

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红车轴草提取物对脂多糖应激断奶仔猪生长性能和炎性介质的影响

ＲＥ对脂多糖（Ｐ）Ｌ￥应激的断奶仔猪生长性能和炎性介质的影响，本试验选用３０头健康、
平均体重为（． ± ．）ｇ２ ± ）龄（ × × ）６８０１ｋ的（１１日杜长大断奶仔猪，随机区组设计设５６４按照个处理，个处理６重复，每个每个重复１头仔猪．５个处理组仔猪分别饲喂基础日粮（非应激空白对照组）基础日粮＋ＬＳＬＳ组）基础日粮＋ＬＳ＋、Ｐ（Ｐ、Ｐ
ＭｅｃｅａｄＢｏｅｎｌｙ３ＦｊｎＰｏｉｃａＥｇｅｒｇＲｓａｃｅｔｒｈｉｎｄｉｎｉｅｈｏｇ；．ｕｉｒｖｉｎｉｅｎｅｅｒｈＣｎｒｏｅｔｏａｎｌｎｉｅｆｔＰｅｅｔｎａｄＣｎｒｌｆｏｎｓ，ｏｇａ，ｕｉ６００ＣｉａｒｖｎｏｏｔｏｏｉＬｎ）Ｆｊｎ３４０，ｈｎ）ｉｎｏｏＺｓｔｎａＡｓｒｃ：ｏｓｄｅｅｅｔｏｄｃｖｒｘａｔＲＣｎｇｏｔｐｒｒａｃｄｉａｍｍｆａａｔｓｆｅｅｉｉｌｂｔｔＴｕｙｔｃｓｆｅｌｅｔｃ（Ｅ）ｏｗｈｅｏｍｅａｆｍｏｐｒｅｗａｄｐｇｗｔａｔｈｆｒｏｅｒｒｆｎｎｎｌｍｅｒｏｎｓｌ
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A literature survey on fatigue analysis approaches for rubber

* Corresponding author. Tel.: +1-419-530-8213; fax: +1-419-5308206. E-mail address: afatemi@ (A. Fatemi).
period during which cracks nucleate in regions that were initially free of observable cracks. The second phase is a period during which nucleated cracks grow to the point of failure. It will be seen that nucleation, growth, and ﬁnal failure may all be rationalized in terms of the fracture mechanical behavior of rubber. There are, however, issues unique to the crack nucleation phase, which deserve careful study. Models for predicting fatigue life in rubber follow two overall approaches. One approach focuses on predicting crack nucleation life, given the history of quantities that are deﬁned at a material point, in the sense of continuum mechanics. Stress and strain are examples of such quantities. The other approach, based on ideas from fracture mechanics, focuses on predicting the growth of a particular crack, given the initial geometry and energy release rate history of the crack. For each approach, existing theories are presented. A discussion of each approach’s strengths and limitations, and examples of how these approaches have been applied in engineering analysis are also included. Some of the information presented in this paper has been reviewed previously [5–13]. This literature survey updates these existing reviews to reﬂect recent and previously unnoticed developments. This survey also offers new interpretations of existing studies and theories, and identiﬁes areas where additional research is needed. Another paper reviews factors that affect the fatigue life of rubber [14]. These include the effects of mechanical

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TRANSIENT FATIGUE-CRACK GROWTH BEHAVIOR FOLLOWING VARIABLE-AMPLITUDE LOADING IN A MONOLITHI

TRANSIENT FATIGUE-CRACK GROWTH BEHAVIOR FOLLOWING VARIABLE-AMPLITUDE LOADING IN A MONOLITHIC SILICON NITRIDE CERAMICC.J.GILBERT and R.O.RITCHIEDepartment of Materials Science and Mineral Engineering,University of California,Berkeley,CA 94720-1760,U.S.A.Abstract ÐTransient cyclic fatigue-crack growth behavior following variable-amplitude loading sequences has been investigated in a hot-pressed,monolithic Si 3N 4ceramic (NTK EC-141),for which fracture toughness (resistance-curve)and cyclic fatigue properties are well characterized.Following rapid changes in the applied stress-intensity range,D K ,during both low-high and high-low block-load-ing sequences,transients were observed in the subsequent cyclic crack-growth rates,but only under speci®c conditions.When K min was suddenly varied with K max held constant,no transients were observed.However,when the load changes involved both D K and K max ,subtle transients were detected in the form of retardations following high-low,and accelerations following low-high block-loading sequences.The transient growth rates were typically a factor of H 2di erent from the steady-state (base-line)growth rates for load excursions of H 30±40%,and lasted for several hundred microns of crack extension before a steady-state growth rate was achieved.Such behavior is rationalized in terms of grain-bridging mechanisms active in the crack wake.#1998Elsevier Science Ltd.All rights reservedKeywords Ðsilicon nitride,fatigue-crack growth,transient crack growth,overloads.1.INTRODUCTIONI T HAS been well established that ambient-temperature crack-growth rates in a wide range of monolithic ceramics and ceramic±matrix composites are orders of magnitude faster under cyclic loading than under static loading at equivalent stress-intensity levels [1±8].Such behavior is gen-erally attributed to a crack-advance mechanism ahead of the crack tip identical to that under static loading,coupled with cyclic-induced degradation of the toughening mechanisms (crack-tip shielding)behind the crack tip.In grain-bridging ceramics,such as silicon nitride,alumina and in situ toughened silicon carbide [3±7],the latter occurs by a progressive degradation in crack bridging due to frictional wear at the sliding grain/matrix interface.Similar behavior in the form of a suppression in ®ber bridging under cyclic loading has been reported for ceramic±matrix composites[9,10].An important,but poorly understood,aspect of ceramic fatigue is the role of load history.Speci®cally,it remains unclear whether or not monolithic ceramics exhibit transient crack-growth behavior following sudden changes in the applied stress±intensity levels.Crack-growth transients are commonly observed in metals [11,12]and in transformation-toughened ceramics [13]following various variable-amplitude loading sequences.Transient retardations,involving orders of magnitude changes in growth rates and even arrest,are generally observed following single-tensile overloads and high-low block-loading sequences,whereas transient accel-erations generally follow low-high block-loading sequences.In both cases,the extent of crack growth a ected is comparable with the size of the overload plastic (or transformation)zone cre-ated ahead of the crack tip [e.g.11,12].In metals,the transients are associated with a variety of mechanisms,including crack de¯ection and the generation of residual stresses in the overload plastic zone,and the subsequent enhanced e ect of crack closure in the wake of the crack tip.Similarly,in phase-transforming ceramics such as Mg-PSZ [13],the transient growth rates are caused by changes in the degree of transformation-toughening resulting from changes in the size of the transformation zone surrounding the crack tip.In non-transforming ceramics,however,it is unclear if such transients exist.Not only are experimental data lacking,but in the few cases where transients were observed they appear to be very minor compared to both metals and transformation-toughened ceramics.This makesEngineering Fracture Mechanics Vol.60,No.3,pp.303±313,1998#1998Elsevier Science Ltd.All rights reserved Printed in Great Britain 0013-7944/98$19.00+0.00PII:S0013-7944(98)00019-8303them quite di cult to quantify,and no mechanistic understanding has been established.For example,recent work of Choi et al .[14,15]indicates subtle transients during low-high and high-low block-loading sequences in a monolithic Si 3N 4.Although this behavior was attributed to a crack-tip process zone of intense microcracking (akin to a plastic zone in metals or a transform-ation zone in phase-transforming ceramics),such microcracking has never been observed exper-imentally.Additionally,Jacobs and Chen [6]reported transient e ects in a monolithic Si 3N 4at the initial stages of fatigue-crack growth experiments.They attributed these to a competition between shielding accumulation (via grain bridging)and shielding degradation (via frictional wear);during transients it was presumed that one of these processes dominates over the other.In other studies on Al 2O 3[2],Si 3N 4[16],and an Al 2O 3/SiC w composite [17],however,no such transient behavior was detected.Accordingly,the objective of the present work is to investigate transient fatigue-crack growth behavior following overload and block-loading excursions in a hot-pressed,monolithicFig.1.Scanning electron micrographs of the microstructure of the NTK EC-141Si 3N 4obtained by etching at H 3408C in molten NaOH for H 1min,at (a)low and (b)higher magni®cation.C.G.GILBERT and R.O.RITCHI304silicon nitride,for which the long-crack,steady-state fracture toughness and cyclic fatigue prop-erties are well known[16].2.EXPERIMENTAL PROCEDURES2.1.MaterialsTests were performed on a commercial grade of silicon nitride manufactured by NTK Technical Ceramics (designation EC-141),hot pressed with H 7wt%Al 2O 3and Y 2O 3as addi-tives.The microstructure is shown in Fig.1(following an etch at 3408C in molten NaOH for H 1min),and consisted of both equiaxed a -Si 3N 4and columnar b -Si 3N 4phases.The a -Si 3N 4grains were typically H 2m m in diameter,while the b -Si 3N 4grains were generally H 1m m wide and H 5m m long.Mechanical properties are summarized in Table 1.{2.2.Fracture and cyclic fatigue-crack growth rate measurementsBoth constant-amplitude (steady-state)and variable-amplitude fatigue-crack growth rate measurements were performed in a controlled room-air environment (228C,45%relative humid-ity)on compact-tension specimens,containing microstructurally long through-thickness cracks (>3mm).Measurements were made with the conventional compact-tension geometry of width,W =25mm,and thicknesses,B ,ranging from 2to 3.5mm.Specimens were cycled using a sinu-soidal waveform at a test frequency of 25Hz under load control (to an accuracy of better than 0.2%)on a computer-controlled,servo-hydraulic mechanical test frame.A range of positive load ratios,R =K min /K max ,were used throughout the experiments and are indicated where appropriate.Constant amplitude loading tests were performed in a general accordance with ASTM standard E647.Stress-intensity factors were computed from handbook solutions [20,21],considered accurate to 20.5%over the range of a /W from 0.2to 1.0,and crack growth rates,d a /d N ,were calculated every 25m m using a simple ®nite di erence method.Crack initiation was facilitated using a half-chevron shaped starter notch,and prior to data collection samples were pre-cracked under cyclic loading for several millimeters beyond this notch.Thereafter,crack lengths were continuously monitored by means of unloading elastic compliance measurements with a 350O strain gauge attached to the back-face of the specimen.To verify such compliance measurements,crack lengths were checked periodically using a travel-ing microscope.Optical and compliance measurements were always found to be within 2%.For the variable-amplitude loading tests,crack-growth rates were measured under constant D K conditions.Once a steady-state crack velocity was achieved at a particular stress intensity range,specimens were subjected to overload and block-loading sequences in order to examine transient crack-growth behavior.All loading changes were performed within a H 2s period.Three general loading sequences were used (Fig.2):.high-low blocks,with a sudden increase in K min at constant K max [Fig.2(a)],.low-high and high-low blocks,with variations in both K max and K min [Fig.2(b)],.single tensile overload [Fig.2(c)].Transient data are presented in terms of the growth rate per cycle,d a /d N ,as a function of crack length,a .Following completion of the constant-amplitude fatigue tests,resistance curves (R-curves)were measured by loading the pre-cracked specimens to failure under load control at a rate of Table 1.Mechanical properties of EC-141Si 3N 4Purity (%)Density (g/cm 3)Modulus,E (Gpa)Bend strength s F (MPa)Initiation toughness,K 0(MPa Z m)Plateau toughness,K c (MPa Z m)Poisson's ratio,#93.0 3.22310900<4.0 5.80.27{This microstructure has not been optimized for strength and toughness,where values exceeding 1000MPa and 10MPa m p ,respectively,have been achieved[18,19].Transient fatigue-crack growth305H 0.05MPa Z m/s.Crack lengths were periodically monitored using unloading compliance,with the unloading excursions limited to less than H 10%of the current load.Data are presented in terms of crack-growth resistance,K R ,plotted as a function of crack extension,D a .All fatigue and fracture surfaces were examined using optical and scanning electron (SEM)microscopy.3.EXPERIMENTAL RESULTS3.1.Fracture and steady-state fatigue-crack growth behaviorThe hot-pressed Si 3N 4exhibited rising R-curve behavior,with the crack-extension resist-ance,K R ,increasing over the ®rst H 450m m of crack extension from an initiation value of K 0H 4MPa Z m to a plateau value of K c H 5.8MPa Z m (Fig.3).Such R-curve behavior has been shown to result from grain bridging [16].This shielding mechanism has been described extensivelyin Fig.2.Schematic illustration of the three types of loading sequences used in this study.In (a)K max was maintained constant,with sudden variations only in K min ,in (b)both K max and K min were varied,and in (c)single tensile overloads were applied within an otherwise constant loadingcondition.Fig.3.Crack-growth resistance,K R ,plotted in terms of both crack extension,D a ,and normalized crack length,a /W ,for the NTK EC-141Si 3N 4[16].C.G.GILBERT and R.O.RITCHI306the literature [e.g.18,19,22±25],and derives from closing tractions developed behind the crack tip across opposing crack faces,either via intact ligaments or contacting surfaces which interfere with one another.Steady-state,cyclic fatigue-crack growth rates (d a /d N )for the EC-141Si 3N 4are plotted as a function of the applied alternating (D K )and maximum (K max )stress intensities in Fig.4[16].Characteristic of ceramic materials,growth rates exhibit a marked dependence on stress inten-sity,with the prime dependency on K max rather than D K .This is apparent by expressing these data in terms of the modi®ed Paris power-law relationship:d a a d N C K max n D K P Y 1where C is a scaling constant (independent of K max ,D K and R ),and n and p are experimentally determined crack-growth exponents.A regression ®t to the data in Fig.4yields values of n H 29and p H 1.3,with C H 3.7Â10À28(units:m/cycle and MPa Z m)[16].3.2.Variable-amplitude fatigue-crack growth behaviorVariable-amplitude growth rates are plotted in Fig.5(a)as a function of crack length,a ,for a series of block-loading segments,where the values of D K and K max for each block are listed in Fig.5(b).Small transients were detected following many of the low-high and high-low block-loading sequences,although this was not always the case (e.g.blocks 5,7and 9);more-over,single tensile overloads appeared to have little e ect on subsequent growth rates.The con-ditions which produce these transients were investigated more carefully below using the three sequences detailed in Fig.2.Figure 6shows a series of high-low block loading sequences which involved only abrupt changes in the minimum applied stress intensity (K min ),with the value of K max heldconstantFig.4.Constant-amplitude,cyclic fatigue-crack growth rates,d a /d N ,for the NTK EC-141Si 3N 4,plotted as a function of (a)D K ,and (b)K max over a range of load ratios,R .Tests were carried out in room air at a frequency of 25Hz (sinewave)[16].Transient fatigue-crack growth 307[e.g.Figure 2(a)].For this type of sequence,there were no observable crack-growth transients.Conversely,where the load changes also involved changes in K max [e.g.Figure 2(b)],small but ®nite crack-growth rate transients were observed.An example is shown in Fig.7,where after reaching a steady-state growth rate of H 2.8Â10À10m/cycle at an applied D K of 3.8MPa Z m,D K was abruptly increased by H 30%to 5.0MPa Z m (K max increased from 4.2to 5.6MPa Z m).This resulted in a small transient acceleration;the initial growth rate of H 3Â10À7m/cycle fol-lowing the low-high sequence,however,was only H 1.5times higher than the baseline growth rate and steady-state was achieved after only H 300m m of crack extension.Similarly,for high-low loading sequences,where the applied D K was rapidly reduced H 30%from 5.0to 3.7MPa Z m (K max decreased from 5.6to 4.1MPa Z m),a transient retardation was observed over H 400m m of crack extension following the excursion;here,the inintial transient crack-growth rate was approximately half the steady-state value [the steady-state value at D K =3.7MPa Z m was clearly achieved as seen from Fig.4(a)].Similar observations have been reported by ChoietFig.6.Cyclic fatigue-crack growth rates,d a /d N ,for the NTK EC-141Si 3N 4are plotted as a function of crack extension,D a ,over a range of block-loading sequences in which K min was rapidly changed at constant K max [Fig.2(a)].Note the lack of detectable transients following changes in K min.Fig.5.Variable-amplitude,cyclic fatigue-crack growth rates,d a /d N ,at R =0.1for the NTK EC-141Si 3N 4,plotted as a function of crack length,a ,throughout a series of block-loading experiments in (a).The driving force conditions for each block are indicated in (b).C.G.GILBERT and R.O.RITCHI308al .[14,15]for a monolithic silicon nitride (hot-pressed at 17508C with 9mol%Y 2O 3and 3mol%Al 2O 3with a K c of 6.5MPa Z m).The e ect of single tensile overloads was also investigated (Fig.8).At baseline cycling at a D K of H 4MPa Z m (R =0.1),sudden 38%overloads to a K max of 5.5MPa Z m were found to result in insigni®cant variations in growth rates.4.MECHANISMS FOR TRANSIENT CRACK-GROWTH BEHAVIORThe occurrence of transient crack-growth behavior during variable-amplitude fatigue load-ing is invariably related to non steady-state conditions existing in an active process zone ahead of the crack tip or in a crack-tip shielding zone behind the crack tip (Fig.9).In simple terms,for a high-low block-loading sequence in metallic materials,the crack at the lower load level in-itially has a larger active plastic zone than it would normally experience under steady-state con-ditions at this level [Fig.9(a)].This enhanced plastic zone induces a transient retardation due to locally enhanced closure loads (which in turn reduces the e ective near-tip D K value)until steady-state conditions are re-established,generally after the crack extends over a distance com-parable with the plastic-zone size associated with the higher loads.The reverse scenario occurs for low-high loading sequences,where transient accelerations are seen.Similar arguments involving the role of crack-tip shielding have been used to explain the marked transient crack-growth behavior seen in partially-stabilized zirconia ceramics (e.g.Mg-Fig.7.Cyclic fatigue-crack growth rates,d a /d N ,for the NTK EC-141Si 3N 4are plotted as a function of crack extension,D a ,over a range of block-loading sequences in which both K min and K max were rapidly changed [Fig.2(b)].Note that small transient accelerations were detected following low-high blocks,and small transients decelerations following high-lowblocks.Fig.8.Cyclic fatigue-crack growth rates,d a /d N ,for the NTK EC-141Si 3N 4are plotted as a function of crack extension,D a .Two single overloads of K max =5.50MPa Z m were applied in an otherwise con-stant loading condition in which D K =4.00MPa Z m and K max =4.44MPa Z m (R =0.1),as illus-trated in Fig.2(C).Transient fatigue-crack growth 309PSZ [13])during fatigue under variable-amplitude loading [Fig.9(b)].Here,for a high-low block-loading sequence,the increased transformation toughening generated by the larger trans-formation zone associated with the higher loads acts to temporarily retard crack growth at the lower load level.Similarly,for a low-high sequence,the initially smaller transformation zone at the higher load level leads to a transient acceleration before the steady-state is re-established.As described above,however,transient crack-growth behavior in non-phase transforming ceramics is much more subtle and currently lacks any adequate explanation.Choi et al .[14,15]Fig.9.Schematic illustration of the mechanisms associated with crack-growth transients following sudden load changes in (a)metallic alloys,(b)transformation-toughened ceramics,and (c)grain-bridging ceramics.C.G.GILBERT and R.O.RITCHI310proposed that the observed transients in a Si3N4resulted from the presence of a crack-tip pro-cess zone ahead of the crack tip,consisting of an intense region of microcracking.The transient acceleration associated with a sudden increase in driving force is attributed to the slow develop-ment of a larger equilibrium process zone,whereas the transient deceleration associated with a sudden decrease in driving force is attributed to the development of a smaller equilibrium pro-cess zone.As in metals,the transient behavior persists until the crack grows through the process zone at the time of the load excursion.Accordingly,this process-zone size should scale directly with the crack extension associated with the crack-growth transient,i.e.several hundred microns (Fig.7).This dimension should be readily resolvable with optical or scanning electron mi-croscopy,yet such microcracking zones were not observed[14,15];indeed,zones of microcrack-ing are generally not seen in single-phase ceramics such as silicon nitride.In the present study,we propose an alternative explanation for crack-growth transients during variable-amplitude fatigue of non-transforming ceramics,speci®cally involving changes in the zone of interlocking grains behind the crack tip[Fig.9(c)].Such bridging zones are known to progressively degrade under cyclic loads due to frictional wear of the sliding grain/matrix interface[1,3±7].However,where transient accelerations are observed after low-high block-load-ing sequences,the rate of bridge degradation can be considered initially to exceed the rate of bridge creation,presumably resulting from a sudden increase in matrix/grain bridge sliding dis-tances.However,as the crack advances,growth rates decrease to the steady-state conditions as new bridging grains are formed.Similarly,following a high-low sequence,the rate of bridge cre-ation can be considered to exceed the rate of bridge degradation;a steady-state growth rate is then achieved when these competing e ects equilibrate.Jacobs and Chen[6]have also advanced similar arguments to explain observed crack-growth transients in their studies on Si3N4.Such a hypothesis can be rationalized in terms of the variation in near-tip(e ective)stress intensity during such load excursions,based on consideration of the crack bridging term,K B. Consider a crack propagating an incremental amount d a while experiencing an incremental increase in fatigue cycles,d N.The bridging term,K B,will increase in proportion to the amount of crack propagation,d a,due to the formation of additional bridging grains spanning the crack faces.Alternatively,this screening term will also decrease in proportion to the number of fatigue cycles,d N,due to frictional sliding and wear[4,5].Following previous arguments[6],this can be expressed mathematically as:d K B d K Bd ad aÀd K Bd Nd N X 2(Note that both terms can be modeled using existing bridging[24]and degradation models[4] given that all adjustable parameters are known).Under steady-state growth conditions,the rate of shielding accumulation and degradation are equal,i.e.d K B=0.Transients are therefore as-sociated with temporary departures from this equilibrium following sudden load changes. Speci®cally,whend K B d a d a bd K Bd Nd N Y 3we observe a transient deceleration,and whend K B d a d a`d K Bd Nd N Y 4we observe a transient acceleration.In other words,a crack growing at low velocity has experienced more loading cycles for a given amount of crack extension than one growing at higher velocity(closer to the K c instabil-ity).It will,therefore,have less active bridging and upon increasing the driving force,the equili-brium bridging value will be higher than the initial value.As a rseult,the crack-tip initially experiences more of the applied far®eld loading until the higher remnant bridging value is obtained.Consequently,there temporarily exists a higher growth rate which decays to the steady-state value.In the opposite case,when the driving force on a crack growing at a higher velocity is suddenly reduced,the initial amount of bridging will be higher than the equilibrium value at the lower driving force.Therefore,one observes an initially low crack growth velocityTransient fatigue-crack growth311which increases to a steady state value.In the case where only K min is altered,it is not surprising to see no apparent transients since the dependence in fatigue-crack growth is primarily on K max .In addition,single tensile overloads apparently have little e ect on the bridging zone as they cause no change in the growth rates,although this is not well understood.While the details of bridging creation and degradation are unclear,eq.(2)provides qualitative insight into transient crack growth processes in non-transforming ceramics.5.SUMMARY AND CONCLUSIONSTransient subcritical crack-growth behavior during variable-amplitude fatigue has been investigated in a hot-pressed monolithic silicon nitride (NTK EC-141),for which the microstruc-tural and mechanical properties are well characterized.pared to metals,crack-growth transients following variable amplitude load excursions in a non-transforming ceramic are relatively minor.2.Following rapid changes in the applied stress intensity (D K and K max ),in the form of low-high and high-low block-loading sequences,small transients were observed in the fatigue-crack growth rates,which lasted for several hundred microns of crack extension before steady-state growth rates were achieved.3.The transient growth rates were typically a factor of H 2di erent from the steady-state (base-line)growth rates for H 30±40%changes in the applied stress intensities.4.By contrast,single tensile overloads caused insigni®cant variations in growth rates.5.The observed transients involved accelerations immediately following low-high sequences and decelerations immediately following high-low sequences.However,no transients were observed for block-loading sequences where K max was held constant.6.Such behavior is rationalized in terms of the e ects of the zone of grain-bridging behind the crack tip,which is progressively diminished under cyclic loads.Speci®cally,the occurrence of transient accelerations or decelerations is related to the relative rates of bridge degradation and bridge creation,which in turn governs the e ective driving force at the crack tip.Acknowledgements ÐThis work was supported by the U.S.National Science Foundation under Grant no.DMR-9522134.REFERENCES1.Ritchie,R.O.and Dauskardt,R.H.,Cyclic fatigue of ceramics:a fracture mechanics approach to subcritical crack growth and life prediction.Journal of the Ceramics Society of Japan ,1991,99,1047±1062.2.Li,M.and Guiu,F.,Subcritical fatigue crack growth in aluminaÐII.Crack bridging and cyclic fatigue mechanisms.Acta Metallurgica et Materialia ,1995,43,1871±1884.thabai,S.,Ro del,J.and Lawn,B.,Cyclic fatigue from frictional degradation at bridging grains in alumina.Journal of the American Ceramic Society ,1991,74,1360±1348.4.Dauskardt,R.H.,A frictional-wear mechanism for fatigue-crack growth in grain bridging ceramics.Acta Metalurgica et Materialia ,1993,41,2765±2781.5.Gilbert,C.J.,Dauskardt,R.H.,Steinbrech,R.W.,Petrany,R.N.and Ritchie,R.O.,Cyclic fatigue in monolithic alumina:mechanisms for crack advance promoted by frictional sear of grain bridges.Journal of Materials Science ,1995,30,643±654.6.Jacobs,D.S.and Chen,I-W.,Cyclic fatigue in ceramics:a balance between crack shielding accumulation and degra-dation.Journal of the American Ceramic Society ,1995,78,513±520.7.Gilbert,C.J.,Cao,J.J.,MoberlyChan,W.J.,De Jonghe,L.C.and Ritchie,R.O.,Cyclic fatigue and resistance-curve behavior in an in situ toughened silicon carbide with Al±B±C additions.Acta Materialia ,1996,44,3199±3214.8.Kishimoto,H.,Cyclic fatigue in ceramics.JSME International Journal ,1991,34,393±403.9.Rouby,D.and Reynaud,P.,Fatigue behaviour related to interface modi®cation during load cycling in ceramic±ma-trix ®bre posite Science Technology ,1993,48,109±118.10.Kotil,T.,Holmes,J.W.and Comninou,M.,Origins of hysteresis observed during fatigue of ceramic±matrix compo-sites.Journal of the American Ceramics Society ,1990,73,1879±1883.11.Ward-Close,C.M.,Blom,A.F.and Ritchie,R.O.,Mechanisms associated with transient fatigue crack growth under variable-amplitude loading:an experimental and numerical study.Engineering Fracture Mechanics ,1989,32,613±638.12.Suresh,S.,Fatigue of Materials .Cambridge University Press,Cambridge,1991.13.Dauskardt,R.H.,Carter,W.C.,Veirs,D.K.and Ritchie,R.O.,Transient subcritical crack-growth behavior in transformation-toughened ceramics.Acta Metallurgica et Materialia ,1990,38,2327±2336.C.G.GILBERT and R.O.RITCHI312Transient fatigue-crack growth313 14.Choi,G.,Cyclic fatigue crack growth in silicon nitride:in¯uence of stress ratio and crack closure.Acta Metallurgicaet Materialia,1995,43,1489±1494.15.Choi,G.,Horibe,S.and Kawabe,Y.,Cyclic fatigue in silicon nitride ceramics.Acta Metallurgica et Materialia,1994,42,1407±1412.16.Gilbert,C.J.,Dauskardt,R.H.and Ritchie,R.O.,Behavior of cyclic fatigue cracks in monolithic silicon nitride.Journal of the American Ceramics Society,1995,78,2291±3300.17.Dauskardt,R.H.,Dalgleish,B.J.,Yao,D.,Ritchie,R.O.and Becher,P.F.,Cyclic fatigue-crack propagation in asilicon carbide whisker-reinforced alumina composite:role of load ratio.Journal of Materials Science,1993,28, 3258±3266.18.Li,C-W.,Lee,D-J.and Lui,S-C.,R-curve behavior and strength for in-situ-reinforced silicon nitrides with di erentmicrostructures.Journal of the American Ceramics Society,1992,75,1777±1785.19.Becher,P.F.,Hwang,S-L.and Hsueh,C-H.,Using microstructure to attack the brittle nature of silicon nitride cer-amics.MRS Bulletin,1995,20,23±27.20.Newman,J.C.,Stress analysis of compact specimens including the e ects of pin loading.In Fracture Analysis,105.ASTM STP560,American Society for Testing and Materials,Philadelphia,1974.21.Strawley,J.E.,Wide range stress intensity factor expressions for ASTM E399standard fracture toughness speci-mens.International Journal of Fracture Mechanics,1976,12,475±476.22.Becher,P.,Microstructural design of toughened ceramics.Journal of the American Ceramic Society,1991,74,255±269.23.Evans,A.G.,Perspective on the development of high-toughness ceramics.Journal of the American Ceramic Society,1990,73,187±206.wn,B.R.,Fracture of Brittle Solids,2nd edn.Cambridge University Press,New York,1993.25.Evans,A.G.and McMeeking,R.M.,On the toughening of ceramics by strong reinforcements.Acta Metallurgica,1986,34,2435±2441.(Received20October1997,accepted1March1998)。

番茄红素抗氧化应激效应对大鼠脊髓损伤后运动功能恢复的影响

Ｉ关键词】番茄红素；脊髓损伤；丙二醛；超氧化物歧化酶；运动功能
Ａｎｔｉ－ｏｘｉｄａｔｉｖｅＳｔｒｅｓｓＥｆｆｅｃｔｏｆＬｙｅｏｐｅｎｅｏｎＭｏｔｏｒＦｕｎｃｉｔｏｎａｔｆｅｒＳｐｉｎａｌＣｏｒｄＩｎｊｕｒｙｉｎＲａｔｓＬＩＸｉａｏ，ＬＵＯＣｈｅｎ－ｙｕ，ＹＡＮＧＺｈｅｎｇ，
ｌ３２一
中国康复理论与实践２０１３年２月第１９卷第２期ＣｈｉｎＪＲｅｈａｂｉｌＴｈｅｏｒｙＰｒａｅｔ，Ｆｅｂ．２０１３，Ｖ０１．１９，Ｎｏ．２
ＤＯＩ：１０．３９６９￣．ｉｓｓｎ．１００６－９７７１．２０１３．０２．００９
・
基础研究・
番茄红素抗氧化应激效应对大鼠脊髓损伤后运动功能恢复的影响
李晓，罗晨禹，杨拯，谢少华，张尧，杨海燕，马延仃，张艺，胡雯思，卓睿，李亚光
【摘要ｌ目损伤后运动功能恢复的影响。方法
健康成年Ｓｐｒａｇｕｅ — Ｄａｗｌｅｖ大鼠
ｗａｓｉ￣ｅｃｔｅｄｆｏｒＧｒｏｕｐＢｉｎｐｅｒｉｔｏｎｅａｌｃａｖｉｔｙ３０ｍｉｎａｆｔｅｒＳＣＩ，２０ｍｇ＆ｇＬＰｗａｓｇｉｖｅｎｏｒａｌｌｙｆｏｒＧｒｏｕｐＣ．Ｏｎｔｈｅ１ｓｔｄ，３ｒｄｄａｎｄ７ｔｈｄａｆ－ｔｅｒｏｐｅｒａｔｉｏｎ，ｔｈｅｍｏｔｏｒｆｕｎｃｔｉｏｎｗａｓａｓｓｅｓｓｅｄｂｙｓｌａｎｔｉｎｇｂｏａｒｄｔｅｓｔａｎｄＢａｓｓｏ — Ｂｅａｔｔｉｅ — Ｂｒｅｓｎａｈａｎｓｃａｌｅ（ＢＢＢ）；ｓｕｐｅｒｏｘｉｄｅｄｉｓｍｕｔａｓｅ（ＳＯＤ）ａｎｄｍａｌｏｎｄｉａｌｄｅｈｙｄｅ（ＭＤＡ）ｗｅｒｅｄｅｔｅｒｍｉｎｅｄｉｎｓｅｒｕｍ．ＲｅｓｕｌｔｓＣｏｍｐａｒｅｄｗｉｔｈＧｒｏｕｐＡ，ｔｈｅｓｃｏｒｅｓｏｆｔｈｅｓｌａｎｔｉｎｇｂｏａｒｄｔｅｓｔｓｉｇｎｉｆ－

红车轴提取物对基质金属蛋白酶的抑制活性

ＬＩＣ０．ＳｎｇＨＥＮ．ａＹｅｈｕ，ＦＡＮＧｅｘｎＸｕ．ｕ
（．研６ｒｏ１０￡ｎ
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砒
尉ｎ，￡ｎ０
０厂
ＤｒｇＤｉｏｙ。Ｃｌｇｌｓｒ＾ｃ０ｋｅ
第４卷７
第１期
吉林大学学报（理学版）ＪｕＪｆｍｎＵｉｒｔＳｉｃｄｉ）ｏｍａ０Ｊｎｖｓｙ（ｃｎｅＥｉｏｅｉｅｔｎ
Ｖｏ＿Ｎｏｌ４７．１
２００９年１月
Ｊｎ２ｏａ０９
研究简报
ｆｕｐ￡ｅｅｔａｔｃｎａｎｎ０％ｎａｏｏｄｎｂｔｍｍｅｘｒｃ０ｔｉｉｇ４ｖｎｉｓｉｈｉｉｓＭＭＰ一６ｃｉｉｙ，ｂ０ｈｎｎＡ，１ａｔｔｖｉｃａｉ
ｇｎｓｅｎａｄｄａｄｅｎｃｎ’ ｎｉｉＭＭＰ— ６ａｔｖｔｆｒｎｎｔｎｃｎｗｅｋｙｉｈｂｔＭＭＰ．６ｃｉｉｕｅｉｔｉｎｉｚｉａｔｉｈｂｔｌｃｉｉｖ，０ｍｏ０ｅｉａａＩｎｉｉ１ａｔｖｔｂｔｙ
红车轴提取物对基质金属蛋白酶的抑制活性
王书英‘ 蒋，坤杨金刚蒲利华。李，，，聪，申烨华房学迅，
（．西北大学合成与天然功能分子化学教育部重点实验室，１新药研究所，化学与材料科学学院，西安７Ｏ６；１ｏ９２．吉林大学分子酶学工程教育部重点实验室，长春１０２；３ｏｌ３．西安市儿童医院，西安７００）ｌ０３

细胞划痕实验迁移率的英文

细胞划痕实验迁移率的英文Cell Scratch Assay Migration Rate: A Critical Analysis.The cell scratch assay, commonly known as the wound healing assay, is a widely used technique to assess the migratory capacity of cells in vitro. This assay is basedon the principle of creating a denuded area or a "scratch" on a monolayer of cells and monitoring the closure of this scratch over time. The migration rate, which is calculated by quantifying the reduction in the scratch area, provides valuable insights into the migratory potential of the cells.Significance of Cell Migration.Cell migration is a fundamental biological process that plays a crucial role in various physiological and pathological conditions. It is involved in embryogenesis, wound healing, immune response, and cancer metastasis. Understanding the migratory behavior of cells is, therefore, crucial for developing therapeutic strategies againstdiseases associated with abnormal cell migration.Principles of the Cell Scratch Assay.The cell scratch assay is performed by creating a defined wound in a confluent monolayer of cells using a sterile pipette tip or a scratcher tool. This wound creates a gap in the cell monolayer, and the surrounding cells migrate into the gap to close it. The migratory capacity of the cells is then assessed by monitoring the closure of the scratch over time, typically using microscopy and image analysis software.Calculation of Migration Rate.The migration rate is a quantitative measure of the rate at which the scratch area is reduced over time. It is calculated by comparing the initial scratch area with the scratch area at subsequent time points. The migration rate can be expressed as a percentage reduction in scratch area per unit time.To calculate the migration rate, the following stepsare typically followed:1. Imaging: Acquire images of the scratch at multiple time points, such as 0 hours (immediately after scratching), 24 hours, and 48 hours. Ensure that the imaging conditions are consistent throughout the experiment.2. Quantification of Scratch Area: Use image analysis software to measure the scratch area at each time point. This software allows for accurate and reproducible measurements.3. Calculation of Migration Rate: Calculate the migration rate by subtracting the scratch area at a later time point from the initial scratch area, dividing by the initial scratch area, and then dividing by the time elapsed. The formula for migration rate (R) can be expressed as:$R = \frac{(A_0 A_t)}{A_0} \times \frac{1}{t}$。

循环加载混凝土的损伤模型---英文

DAMAGE M O D E L FOR C Y C L I C LOADING OF CONCRETE
By Wimal Suaris, 1 Member, ASCE, Chengsheng Ouyang, 2 and Viraj M. Fernando 3 ABSTRACT: A damage model for monotonic and cyclic behavior of concrete is developed. The model recognizes the tortuous nature of cracks in compression, which affects the flexibility of the material in a direction coinciding with the average plane of the cracks. An elastic potential is introduced in terms of the principal stresses and a compliance tensor dependent on the accumulated damage. Damage evolution is obtained using a loading surface and bounding surface, defined in terms of the thermodynamic force conjugates of the damage variables. The damage growth during a series of unaxial compression and cyclic tests is inferred from the amplitude attenuation of ultrasonic waveforms transmitted laterally through the specimen while the tests are in progress. The behavior of concrete under compression, tension, biaxial loading and cycling, and damage growth under both monotonic and cyclic loading are found to be predicted well by the proposed theory.

journal Cyclic Fatigue-Crack Growth and Fracture Properties in Ti 3SiC 2 Ceramics at Elevat

Cyclic Fatigue-Crack Growth and Fracture Properties inTi 3SiC 2Ceramics at Elevated TemperaturesDa Chen †Department of Materials Science and Engineering,Shanghai Jiao Tong University,ChinaKiroshi Shirato †Department of Mechanical Engineering,Nagaoka University of Technology,Niigata,JapanMichel W.Barsoum *and Tamer El-Raghy *Department of Materials Engineering,Drexel University,Philadelphia,Pennsylvania19104Robert O.RitchieMaterials Sciences Division,Lawrence Berkeley National Laboratory,and Department of Materials Science and Engineering,University of California,Berkeley,California 74720The cyclic fatigue and fracture toughness behavior of reactive hot-pressed Ti 3SiC 2ceramics was examined at temperatures from ambient to 1200°C with the objective of characterizing the high-temperature mechanisms controlling crack parisons were made of two monolithic Ti 3SiC 2materials with fine-(3–10␮m)and coarse-grained (70–300␮m)micro-structures.Results indicate that fracture toughness values,derived from rising resistance-curve behavior,were signifi-cantly higher in the coarser-grained microstructure at both low and high temperatures;comparative behavior was seen under cyclic fatigue loading.In each microstructure,⌬K th fatigue thresholds were found to be essentially unchanged between 25°and 1100°C;however,there was a sharp decrease in ⌬K th at 1200°C (above the plastic-to-brittle transition temperature),where significant high-temperature deforma-tion and damage are first apparent.The substantially higher cyclic-crack growth resistance of the coarse-grained Ti 3SiC 2microstructure was associated with extensive crack bridging behind the crack tip and a consequent tortuous crack path.The crack-tip shielding was found to result from both the bridging of entire grains and from deformation kinking and bridging of microlamellae within grains,the latter forming by delamination along the basal planes.I.IntroductionTHEternary compound Ti 3SiC 2exhibits a surprising combina-tion of properties for a ceramic;e.g.,it displays high toughness (K c Ͼ8MPa ⅐m 1/2),a ratio of hardness (ϳ4GPa)to elastic modulus (ϳ320GPa)more typical of a ductile material,relatively low density (4.5g/cm 3),and high resistance to thermal shock 1,2and oxidation.3Processed by reactive hot-pressing techniques from powders of titanium,SiC,and graphite,Ti 3SiC 2also shows a range of inelastic deformation modes not typically seen in ceramics at room temperature,4–9including grain bending,grain buckling,and significant amounts of basal slip.In general,Ti 3SiC 2appears to be one of the most damage tolerant of all nontransform-ing monolithic ceramics.4–10The essential feature of the Ti 3SiC 2structure is the relatively weak bonding between the silicon layer and the TiC octahedra,especially in shear,along the basal plane.Dislocations are mobile and multiply at room temperature.They are confined to two orthogonal directions:basal plane arrays (wherein the dislocations exist on identical slip planes)and walls or kink boundaries.Thus,in addition to regular slip,mechanisms for ambient temperature plastic deformation in Ti 3SiC 2are thought to involve the readjust-ment of local stress and strain fields from kink band (boundaries)formation,buckling,and delamination of individual grains (the delamination and associated damage being contained by the kink boundaries).9The delaminations typically occur at the intersection of the walls and arrays and result in the annihilation of the latter.It is this containment of damage that is believed to be a major source of damage tolerance in Ti 3SiC 2.Specifically,delamination along the weaker basal planes leads to the creation of microlami-nae contained within a single grain,the deformation and distortion of such laminae providing a potent contribution to toughening.In general,the plastic behavior seen in Ti 3SiC 2is unusual for carbides and is believed to be caused by this layered structure and the metallic nature of the bonding.9With such intrinsic deformation and toughness properties,Ti 3SiC 2clearly offers some potential for many structural applica-tions;however,if this is to be realized,it is important that some assessment be made of its subcritical crack-growth behavior.The fracture toughness and cyclic fatigue-crack growth behavior of monolithic Ti 3SiC 2at ambient temperatures was first characterized by Gilbert et al.10in both fine-(3–10␮m)and coarse-grained (50–200␮m)conditions.Fatigue-crack growth thresholds,⌬K th ,were found to be as high as 6and 9MPa ⅐m 1/2in the fine-and coarse-grained structures,respectively;corresponding steady-state (plateau)resistance-curve (R-curve)fracture toughnesses,K c ,were measured at 9.5and 16MPa ⅐m 1/2,respectively.The high tough-ness (the K c value for the coarse-grained structure is thought to be the highest ever reported for a monolithic,nontransforming ce-ramic)was found to be associated with a profusion of crack-bridging processes active in the crack wake.Indeed,Ti 3SiC 2J.Ro ¨del—contributing editorManuscript No.188150.Received November 14,2000;approved July 16,2001.Supported by the Office of Science,Office of Basic Energy Sciences,Materials Sciences Division,U.S.Department of Energy (Contract No.DE-AC03-76SF00098)for the Berkeley group,and by the Army Research Office (DAAD19-00-1-0435)for the Drexel University group.*Member,American Ceramic Society.†Formerly of the Department of Materials Science &Engineering,University of California,Berkeley.J.Am.Ceram.Soc.,84[12]2914–20(2001)2914displayed a far larger degree of grain bridging and sliding 6,7,10than has been observed in other ceramics,such as Al 2O 3,Si 3N 4,and SiC,11–13possibly because the deformation processes ob-served in individual grains enhanced grain bridging by increasing pullout distances and suppressing grain rupture.Fatigue-crack growth,on the other hand,was associated with the cyclic loading induced degradation of such bridging;substantial evidence was found for wear degradation at active bridging sites behind the crack tip,particularly in the coarser-grained structure.Although significant progress has been made in understanding the mechanical behavior of polycrystalline Ti 3SiC 2at ambient temperatures,far less is understood about its properties at elevated temperatures.Its mechanical response in tension and creep up to 1200°C has been documented.14–17The tensile response was found to be a strong function of strain rate and temperatures,and was characterized by a plastic-to-brittle transition ‡(PBTT)at temperatures above 1150°C.Recently,Li et al.18confirmed the ductile-to-brittle transition and showed that the Mode I fracture toughness (measured at 4.5MPa ⅐m 1/2)was not only significantly lower than that measured at ambient temperature by Gilbert et al.,10but that above 1273°C,the toughness decreased rapidly to ϳ1⁄2its room temperature value.However,as far as we are aware,nothing is known about its cyclic fatigue-crack growth behavior and the R-curve properties at these elevated temperatures.It is the objective of the present paper to examine the mechanistic aspects of crack growth in Ti 3SiC 2under both static and cyclic loading,in both fine-and coarse-grained microstructures,at temperatures from ambient to 1200°C,with the ultimate aim of defining the salient damage and crack-tip shielding mechanisms that govern the fracture toughness and cyclic fatigue-crack growth properties.II.Experimental Procedure(1)Material Processing and MicrostructuresTi 3SiC 2specimens were fabricated by a reactive hot isostatic pressing technique using TiH 2(Ϫ325mesh,99.99%),SiC (grain size ϳ20␮m,99.5%),and graphite (grain size ϳ1␮m,99%)starting powders.Powders with the desired stoichiometry were mixed in a ball mill,cold isostatically pressed at 200MPa,and annealed at 900°C for 4h in vacuo (10Ϫ3Pa)to remove the hydrogen.Specimens were then sealed in glass under vacuum and hot isostatically pressed for 4h at 1400°or 1600°C to form thefine-and coarse-grained microstructures,respectively.The coarse-grained structure consisted of large plate-like grains of 70–300␮m diameter and 5–30␮m thickness,whereas grains in the fine-grained microstructure had a diameter of ϳ7␮m and were ϳ3␮m thick (Fig.1).XRD and SEM showed that both structures contained ϳ2vol%unreacted SiC.Further details on the proce-dures for fabricating Ti 3SiC 2and its microstructural evolution have been reported elsewhere.1,2(2)Cyclic Fatigue TestingCyclic fatigue-crack growth tests were performed using 2.9mm thick compact-tension C(T)specimens (of ϳ19mm width);this geometry conforms to the ASTM Standard E-647for fatigue-crack growth rate measurements.Specimens were cycled at a constant load ratio (ratio of minimum to maximum applied loads)of R ϭ0.1and a loading frequency of ␯ϭ25Hz (sinusoidal wave form),using automated stress-intensity K control.Growth rates were monitored under decreasing ⌬K conditions with a normalized K -gradient (1/K ⅐d K /d a )set to Ϫ0.08mm Ϫ1.Before testing,the C(T)samples were polished with ϳ1␮m diamond paste and fatigue precracked at room temperature for several millimeters beyond the half-chevron-shaped starter notch.The latter notch geometry is used to facilitate crack initiation in brittle materials.At elevated temperatures (1100°and 1200°C),testing was performed on a computer-controlled servo-hydraulic mechanical testing system in flowing gaseous argon at atmospheric pressure in an environmental chamber/furnace,heated by graphite elements that maintain temperature to within Ϯ1°C.Heating and cooling rates were kept at 10°C/min to minimize any thermal shock effects.After reaching the desired temperature and before commencing the test,the furnace temperature was kept constant for 1–3h to allow the system to reach thermal equilibrium.At ambient temperatures,fatigue-crack growth rates were con-tinuously measured using the standard back-face strain elastic-unloading compliance method,using a 350⍀strain gauge affixed to the back surface of the specimen.This method could not be readily used at elevated temperatures;however,because Ti 3SiC 2has good electrical conductivity at all temperatures,crack lengths at 1100°and 1200°C could be monitored in situ using the electrical-potential method.Full details of the use of electrical-potential methods for crack length monitoring in ceramics are given elsewhere.19,20To verify such measurements at both ambi-ent and elevated temperatures,readings were checked periodically using an optical microscope.Cyclic fatigue-crack growth data are presented in terms of the growth rate per cycle,d a /d N ,as a function of the applied stress-intensity range,⌬K (ϭK max ϪK min ),the latter being computed using standard linear-elastic handbook solutions.Crack‡The exact details of the plastic-to-brittle transition temperature (PBTT)are not understood at this time.It is fairly well-established,however,that it does not involve the activation of five or more independent slip systems,4,14–16which is why we refer to the transition temperature as a PBTT instead of the more commonly used terminology of a ductile-to-brittle transitiontemperature.Fig.1.Optical micrographs of the microstructures in Ti 3SiC 2with (a)fine grains (3–10␮m)and (b)coarse grains (70–300␮m)etched in a 1:1:1by volume HF:HNO 3:H 2O solution.Note the layered nature of the grains,which is particularly evident in the coarse-grained structure.December 2001Cyclic Fatigue-Crack Growth and Fracture Properties in Ti 3SiC 2Ceramics2915profiles of selected fatigue samples,taken at midsection of the test piece perpendicular to the fracture surface,were examined with SEM.(3)R-Curve TestingFollowing completion of the fatigue tests,the R-curves were measured from the remaining uncracked portion of the specimens by loading specimens to failure at elevated temperatures.Tests were conducted under displacement control by monotonically loading the C(T)specimens and monitoring subcritical crack growth using the electrical-potential method with periodic unloads ofϽ10%.During the test,the applied load versus potential-drop curve was recorded for the further calculation of the R-curve.To minimize the effect of pre-existing crack-tip shielding on the measured initiation toughness from any grain bridging that may have developed during prior crack extension,before data collec-tion,specimens were cycled at room temperature forϳ24h at the fatigue threshold,⌬Kth,and no crack extension was detected during this period.III.Results and Discussion(1)Cyclic Fatigue-Crack Growth BehaviorThe variation in fatigue-crack growth rates,d a/d N,with applied stress-intensity range,⌬K,in Ti3SiC2at temperatures of25°, 1100°,and1200°C(Rϭ0.1,␯ϭ25Hz(sin wave)),is shown in Fig.2for the fine-and coarse-grained microstructures.Characteristic of ceramic materials at low homologous temper-atures,21,22growth rates in both Ti3SiC2microstructures at25°C display a marked sensitivity to stress intensity.In terms of a simple Paris power-law expression,d a/d Nϰ⌬K m,the exponent m was measured to be between72and85.Increasing the temperature to 1100°C results in only a small increase in growth rates,with fatigue thresholds reduced fromϳ8to7MPa⅐m1/2in the coarse-grained material,and fromϳ6.5to6MPa⅐m1/2in the fine-grained material;the slopes of the growth-rate curves remain essentially the same with mϷ79to82.However,at1200°C,which is just above the PBTT of Ti3SiC2,4,14–18the behavior is significantly different.Below ϳ10Ϫ8m/cycle,growth rates are substantially faster,the slope of the growth-rate curves are reduced,and there is a sharp decrease in the fatigue thresholds by a factor ofϳ2and3MPa⅐m1/2toϳ2.3 and4.5MPa⅐m1/2in both fine-and coarse-grained structures, respectively.For many applications of structural materials,design must be based on the concept of a fatigue threshold⌬Kthfor nocrackparison of cyclic-fatigue properties of the coarse-grained Ti3SiC2with behavior in a range of metals,intermetallics,and ceramics at both ambient and elevated temperatures.2916Journal of the American Ceramic Society—Chen et al.Vol.84,No.12pared with more traditional structural ceramics,such as Si 3N 4,23–25Al 2O 3,26and SiC,20which have ⌬K th thresholds at 1000°to 1100°C of roughly 2to 4MPa ⅐m 1/2,depending on material and testing conditions,the cyclic fatigue properties of Ti 3SiC 2(especially the coarse-grained microstructure)are clearly superior in terms of the fatigue threshold at both low and high temperatures.Indeed,at temperatures below its PBTT,Ti 3SiC 2exhibits one of the highest fatigue thresholds and cyclic crack-growth resistances observed in a monolithic nontransforming ceramic or intermetallic.This is shown in Fig.3by a comparison of the cyclic-fatigue properties of the coarse-grained Ti 3SiC 2with corresponding results in a range of metals,intermetallics,and ceramics at ambient and elevated temperatures.(2)R-Curve MeasurementsSubstantial rising R-curves were measured in both fine-and coarse-grained Ti 3SiC 2at both ambient and elevated temperatures (see Fig.4).The R-curve for the fine-grained Ti 3SiC 2rose inincrements of ϳ1.6MPa ⅐m 1/2at 25°C and ϳ1.4MPa ⅐m 1/2at 1100°C,to a steady-state toughness of roughly 10and 6MPa ⅐m 1/2,respectively.The corresponding coarse-grained Ti 3SiC 2exhibited a more steeply rising R-curve with increments of ϳ4.9to 5.3MPa ⅐m 1/2at 25°C and 2.3MPa ⅐m 1/2at 1200°C;maximum steady-state toughnesses for this structure were roughly 12–16and 7MPa ⅐m 1/2,respectively.It is clear that the crack-growth resistance of the coarse-grained microstructure is far superior to that of the fine-grained microstructure.It should be noted that the initial flaw size,a 0,appeared to affect the initial toughness,which is manifest as the starting point of the R-curve.This is considered to be associated with residual bridging in the fatigue precrack.As discussed below,there is significantly more bridging developed in the coarse-grained Ti 3SiC 2structure,which exhibits a higher plateau toughness at both ambient and elevated temperatures.Crack paths in this microstructure were significantly more tortuous,and the lengths of bridging zones in the crack wake considerably larger,which are both factors that result in the development of the enhanced crack-wake bridging.At 1200°C,i.e.,above the PBTT of Ti 3SiC 2,the coarse-grained microstructure still displays a rising R-curve,although the extent of stable crack growth is much reduced from ϳ3to 4mm at ambient temperatures to Ͻ0.5mm at 1200°C.This is presumably caused by the onset of macro (creep)deformation of the specimen and,as discussed below,extensive microcracking at these elevated temperatures.It is important to note that in contrast to most metallic and intermetallic solids for which the fracture toughness increases significantly above the PBTT,in Ti 3SiC 2it drops.This is signif-icant because it rules out the activation of additional slip systems as being the cause of the transition,which is in total accord with previous interpretations.4,14–17The fact that it drops must there-fore be related to a relaxation process ahead of the crack tip.It has recently been shown that creep and the tensile properties of Ti 3SiC 2are dominated by such stress relaxation.15–17Currently,the exact mechanism for the relaxation is not well-understood,but is conjectured to result from an annealing out of dislocation pileups by a combination of delamination,microc-racking,and/or grain boundary decohesion.The current work,while not proving these mechanisms exist,certainly supports their presence.(3)Microstructural CharacterizationsSEM images of the crack profiles during fatigue-crack growth at 25°and 1100°C in the coarse-grained microstructure are shown in Fig.5.It is apparent that at 1100°C and below,i.e.,below the PBTT,damage and shielding mechanisms are essentially un-changed from those at ambient temperatures.The fracture modeatFig.5.SEM images of the profiles of fatigue cracks propagating at (a)25°C and (b)1100°C in the coarse-grained Ti 3SiC 2microstructure.Note the crack-wake grain bridging at both temperatures,and the layered nature of the grains.Arrows indicate direction of crack growth.December 2001Cyclic Fatigue-Crack Growth and Fracture Properties in Ti 3SiC 2Ceramics2917both 25°and 1100°C is predominantly intergranular,or “interla-mellar ”,because of the delamination within the grains along the boundaries between the microlamellae (akin to a layered micro-structure);moreover,ϳ10%–20%of the fracture surface consists of transgranular/“translamellar ”cracking.Similarly,the principal shielding mechanism in the crack wake remains grain bridging at both ambient and elevated temperatures.Specifically,wake shielding can be observed both in the form of bridged grains and their frictional pullout,and by an individual or clusters of lamella (typically deformed by bending)within indi-vidual grains (Fig.6).This heavy deformation,together with the intrinsic strengths of the Ti-C bonds within the lamella,10,16can increase the proportion of grains and lamellae participating in the bridging process,which acts to enlarge the extent of the bridging zone and,hence,increases the degree of toughening.Moreover,the bridging mechanisms associated with the lamellae appear to be particularly potent.Unlike other layered structures like mica and graphite,the debonding along the basal planes is bounded within a single grain and the resulting lamellae are quite flexible;the associated bending and kink band formation during deformation provides an ideal energy-absorbing microstructural feature to increase the resistance to crack growth.As this is not seen in more traditional ceramics,it presumably accounts for the superior damage tolerance shown by Ti 3SiC 2.Typically,grain bridging and frictional pullout constitute the major source of extrinsic toughening (i.e.,crack-tip shielding)in monolithic ceramics,such as Si 3N 4,Al 2O 3,and SiC.12In Ti 3SiC 2,the additional shielding process of the bending of deformed lamella appears to further enhance crack-growth resistance and,hence,the steady-state fracture toughness,by promoting very sizable bridging-zone lengths,which can exceed ϳ5mm in the coarse-grained microstructure.Moreover,shear-faulting along the basal planes of the grain caused by sliding along the contacting surface of a bridge may also increase crack-growth resistance by reducing the severity of frictional damage in the layered micro-structures.Below the PBTT,which has been estimated to lie between 1100°and 1200°C,4,14–18no noticeable evidence of extensive creep damage,in the form of microcracking zones and cavitation ahead of the crack tip or viscous-ligament bridging by a grain-boundary glassy phase in the wake,could be seen.As such,it is clear that the primary damage (intergranular and/or interlamellar cracking and the frictional wear degradation of the bridging zone in the crack wake)and crack-tip shielding (grain and lamellae bridging)mechanisms governing high-temperature fatigue-crack growth behavior in Ti 3SiC 2up to ϳ1100°C,are essentially unchanged from those at ambient temperature.The small decrease in crack-growth resistance at 1100°C (⌬K th thresholds are ϳ8%–14%higher at 25°C)may be rationalized by considering the nature of grain bridging 27and its degradation under cyclic loading caused by frictional wear.28,29The pullout resistance from frictional tractions generated via sliding contact of opposing crack faces 11is proportional to the normal stress acting on the interface,which,in turn,is a function of the residualstressFig.6.SEM images of the profiles of fatigue cracks propagating at (a)25°C and (b)1100°C in the coarse-grained Ti 3SiC 2microstructure,showing crack wake shielding by an individual or clusters of lamella within individual grains.Arrows indicate direction of crackgrowth.Fig.7.SEM images of the profiles of fatigue cracks propagating at 1200°C,which is above the PBTT of Ti 3SiC 2in the coarse-grained Ti 3SiC 2microstructure:(a)crack-wake region and (b)crack-tip region.Note the profuse amounts of microcracks and cavities near the main crack.Arrows indicate direction of crack growth.2918Journal of the American Ceramic Society —Chen et al.Vol.84,No.12resulting from thermal expansion anisotropy during cooling from the processing temperature.As the residual stresses will “anneal out ”with increasing temperature,the normal stress will decrease,thereby reducing the pullout resistance.However,at 1200°C,above the PBTT,there is a striking change in behavior,principally in the form of significantly reduced ⌬K th thresholds that can be attributed to the onset of significant high-temperature deformation and associated microstructural dam-age.As noted above,at this temperature,macroscopic deformation of the test specimen was observed in the form of significant widening of the notch.More importantly,extensive microcracking and,to a lesser extent,cavitation,was apparent throughout the sample in both microstructures,although the intensity of damage was most severe in the vicinity of the crack tip.This is shown in Figs.7and 8for coarse-and fine-grained microstructures,respec-tively.Moreover,crack paths involved significantly more trans-granular and/or translamellar cracking,especially in the fine-grained microstructure,compared with that seen at lower temperatures,which severely diminished the propensity for grain bridging in the crack wake.In summary,Ti 3SiC 2represents an extremely damage-tolerant ceramic,with exceptional ambient temperature toughness and only a small degradation in fracture and fatigue properties at elevated temperatures below ϳ1100°C.Of the two structures examined,the coarse-grained microstructure displays a markedly higher K c fracture toughness and ⌬K th threshold at both low and high temperatures,principally resulting from substantially more tortu-ous crack paths and extensive crack-wake bridging because of much larger bridging zones,typically ϳ4–5mm in length in the coarse-grained structures compared with less than ϳ200␮m in the fine-grained structure.However,the mechanical properties of this material degrade substantially at 1200°C above the PBTT because of the onset of significant high-temperature damage mechanisms,specifically,macroscopic (creep)deformation and extensive mi-crocracking.IV.ConclusionsBased on a study of cyclic fatigue-crack propagation behavior at elevated temperatures up to 1200°C in a reactively hot-pressed monolithic Ti 3SiC 2ceramic with both fine-(3–10␮m)and coarse-grained (70–300␮m)microstructures,the following con-clusions can be drawn:(1)The cyclic fatigue-crack growth properties of Ti 3SiC 2were found to be superior to those of more traditional monolithic (nontransforming)structural ceramics,such as Si 3N 4,Al 2O 3,and SiC,both at ambient and elevated temperatures up to ϳ1100°pared with behavior at ambient temperature,fatigue thresh-olds were only reduced by between ϳ8%and 13%at 1100°C;specifically,⌬K th values were found to be ϳ6and 7MPa ⅐m 1/2in the fine-and coarse-grained microstructures,respectively.In both structures and at all temperatures below ϳ1100°C,growth rates displayed a marked sensitivity to stress intensity,with Paris power-law exponents varying from m Ϸ72to 85.(2)Mechanistically,damage and crack-tip shielding processes that were active at 1100°C were essentially unchanged from those at ambient temperature.Crack-path profiles showed a predomi-nantly intergranular and/or interlamellar mode of crack advance with consequent crack-tip shielding by both grain and lamellae bridging in the crack wake.The lamellae bridging mechanism,associated with basal plane delaminations within single grains,appears to be a particularly potent source of toughening.As this is not seen in more traditional ceramics,it presumably accounts for the superior damage tolerance shown by Ti 3SiC 2.(3)At 1200°C,above the plastic-to-brittle transition temper-ature,a striking change in the shape of the growth-rate curves was observed with ⌬K th thresholds being severely reduced by a factor of ϳ2and 3to ϳ2.3and 4.5MPa ⅐m 1/2in the fine-and coarse-grained microstructures,respectively.Such behavior was attributed to the onset of significant high-temperature damage,in the form of macroscopic deformation of the sample,the presence of widespread microcracking and cavity formation,and the onset of some degree of transgranular and translamellar cracking,which limited shielding by grain bridging in the crack wake.(4)Substantial rising R-curves were measured in both fine-and coarse-grained Ti 3SiC 2,and at both ambient and elevated temperatures with stable crack extensions between 0.5and 4mm.(5)Higher (steady-state)R-curve fracture toughness and cy-clic fatigue-crack growth resistance were achieved in the coarse-grained Ti 3SiC 2microstructure at both low and elevated temper-atures.Indeed,this microstructure exhibited ϳ15%–48%higher fatigue thresholds than the fine-grained microstructure at all temperatures between 25°and 1200°C.This is attributed to more extensive crack-wake bridging in the coarser structure in the form of larger bridging zones and a greater tortuosity in crack paths.References1M.W.Barsoum and T.El-Raghy,“Synthesis and Characterization of a Remark-able Ceramic:Ti 3SiC 2,”J.Am.Ceram.Soc.,79,1953–56(1996).2T.El-Raghy and M.W.Barsoum,“Processing and Mechanical Properties of Ti 3SiC 2,Part I,Reaction Path and Microstructure Evolution,”J.Am.Ceram.Soc.,82,2849–54(1999).3M.W.Barsoum,T.El-Raghy,and L.Ogbuji,“Oxidation of Ti 3SiC 2in Air,”J.Electrochem.Soc.,144,2508–16(1997).4T.El-Raghy,M.W.Barsoum,A.Zavaliangos,and S.Kalidindi,“Processing and Mechanical Properties of Ti 3SiC 2,Part II,Effect of Grain Size and Deformation Temperature,”J.Am.Ceram.Soc.,82,2855–60(1999).5M.W.Barsoum and T.El-Raghy,“Room-Temperature Ductile Carbides,”Metall.Mater.Trans.A ,30,363–69(1999).Fig.8.SEM images of the profiles of fatigue cracks propagating at 1200°C in the fine-grained Ti 3SiC 2microstructure:(a)crack-wake region and (b)crack-tip region.Note the profuse amounts of microcracks and cavities near the main crack.Arrows indicate direction of crack growth.December 2001Cyclic Fatigue-Crack Growth and Fracture Properties in Ti 3SiC 2Ceramics2919。

高三生物一轮复习课时作业生长素的发现及生理作用

付兑市石事阳光实验学校课时作业() [第30讲生长素的发现及生理作用]基础巩固1．2011·家研究胚芽鞘向光弯曲现象，逐步揭示了发生这种激反的一因果相关事件，下列按因果相关事件顺序排列的是( )a．胚芽鞘尖端合成生长素b．胚芽鞘尖端感受光刺激c．胚芽鞘向光弯曲生长d．生长素在背光一侧分布较多e．背光一侧细胞生长较快f．单侧光照射胚芽鞘尖端A．a→c→e→b→f→dB．a→b→d→e→c→fC．a→f→b→d→e→cD．a→b→f→d→e→c2．(双选)下列关于植物生长素的作用及其用的叙述中，不正确的是( ) A．成熟细胞比幼嫩细胞对生长素更为敏感B．顶端优势不能够说明生长素作用的两重性C．适茎生长的一浓度的生长素往往抑制根的生长D．可利用生长素类似物防止落花落果能力提升3．2011·温州八校图K301表示生长素浓度对植物根、芽和茎生长的影响，此图给你的信息是( )图K301A．生长素对3种器官的作用具有两重性，低浓度促进生长，高浓度(例如10－2 ppm)抑制生长B．A、B、C点对的生长素浓度分别是促进根、芽、茎生长的最适浓度C．D点对的生长素浓度对茎的生长具有促进作用，却抑制芽的生长D．幼嫩的细胞对生长素反灵敏，成熟的细胞对生长素反不灵敏4．2011·模拟如图K302中甲为接受单侧光照的胚芽，图乙为水平放置一段时间后的胚根，下列关于生长素作用的表述不正确的一项是( )甲乙图K302A．相同浓度的生长素对a处和c处的作用效果可能不同B．a侧生长素浓度降低抑制生长C．b侧生长素浓度升高促进生长D．c、d两点生长素浓度相同促进生长5．某同学探索向光性时，改变放置在胚芽鞘切面一侧的琼脂块中的生长素浓度，测胚芽鞘弯曲生长的角度，结果如下表。

表格数据说明( )B．胚芽鞘弯曲的程度只与生长素有关，与其他激素无关C．琼脂块中生长素浓度为0.35 mg/L时，胚芽鞘生长受抑制D．促进胚芽鞘生长的最适生长素浓度在0.25 mg/L左右6．进行研究性学习的同学对某品种番茄的花进行人工去雄后，用不同浓度的生长素类似物2,4D涂抹子房，得到的无子番茄果实平均重量见图K303：图K303据图得出的正确结论是( )A．2,4D浓度超过25 mg·L－1，对果实发育起抑制作用B．2,4D与生长素的作用效果相同C．2,4D诱导无子番茄的最适浓度大约为20 mg·L－1D．2,4D不仅有促进茎伸长的作用，还有促进果实发育的作用7．(双选)2030年代，家发现单侧光能引起某些植物体内生长素分布不均匀；2080年代家发现单侧光能引起某些植物体内抑制生长的物质分布不均匀。

Fatigue behavior of nanocrystalline metals and alloys

Fatigue behavior of nanocrystalline metals and alloysT.Hanlon a,c,E.D.Tabachnikova b,S.Suresh a,*a Department of Materials Science and Engineering,Massachusetts Institute of Technology,Cambridge,MA02139,USAb B.Verkin Institute for Low Temperature Physics and Engineering,National Academy of Sciences of the Ukraine,Kharkov-103,61103Ukrainec GE,Global Research Center,Niskayuna,NY12309,USAAvailable online2August2005AbstractIn this work,the stress-life fatigue behavior and fatigue crack growth characteristics of pure Ni were studied as a function of grain size spanning a range of tens of nanometer to tens of micrometer.The fatigue response of electrodeposited,fully dense,nanocrystalline pure Ni, with average and total range of grain sizes well below100nm,was compared and contrasted with that of electrodeposited ultraﬁne-crystalline pure Ni with an average grain size of about300nm and conventional microcrystalline Ni with an average grain size in excess of 10m m.It was found that grain reﬁnement to the nanocrystalline regime generally leads to an increase in total life under stress-controlled fatigue whereas a deleterious effect was seen on the resistance to fatigue crack growth at low and high tensile load ratio levels.To explore the generality of the above trends,systematic experiments were also performed in ultraﬁne-crystalline pure Ti produced by equal-channel angular pressing where a reduction in grain size was found to cause an increase in fatigue crack growth rates at different tensile load ratios. Grain reﬁnement from the microcrystalline to the ultraﬁne-crystalline regime by cryomilling of Al alloys also showed a similar response. Possible mechanistic origins of such trends are explored,and some general conclusions are extracted on strategies for improvements in the fatigue resistance of engineering structures by recourse to grain reﬁnement down to the nanocrystalline regime.q2005Elsevier Ltd.All rights reserved.1.IntroductionControl of the resistance of metals and alloys to fracture and fatigue through grain reﬁnement has long been a strategy for improving the structural integrity of engineering components.The vast majority of studies in this area have dealt with microcrystalline(mc)metals and alloys with an average grain size typically larger than1m m(see,for example[1]),although a limited amount of experimental information is also available for ultraﬁne-crystalline(ufc) metals[2–6].An examination of these studies leads to the following general observations[1,5].Grain reﬁnement, which leads to the strengthening of mc and ufc metals and alloys,is usually accompanied by an increase in fatigue endurance limit.Consequently,stress-life curves,which provide an indication of the dependence of total fatigue life on cyclic stress in nominally smooth fatigue specimens subjected to constant amplitude cyclic load,generally point to an improvement in fatigue resistance with decreasing grain size.Here the beneﬁcial effect on fatigue life is typically considered to arise from the inhibition of cracks that nucleate at nominally smooth surfaces in response to the higher endurance limit.In contrast to this trend,fatigue damage tolerance generally deteriorates with grain reﬁne-ment especially at low stress intensity range,D K,in the near-threshold regime.This apparent increase in resistance to fatigue fracture with increasing grain size is attributed to possible effects of the lowering of effective driving force from microstructurally induced changes in crack path and the attendant possibility of contact between crack face asperities[7].This apparently beneﬁcial effect is typically more pronounced at low D K levels where the cyclic plastic zone size and the cyclic crack tip opening displacement are smaller than the grain size in mc metals.The foregoing general trends pertaining to grain size effects of fatigue crack initiation and crack growth have thus far not been assessed fully for the broad spectrum of grain sizes spanning the microcrystalline to the nanocrystalline regime.Indeed,the recent surge in research interest in nanostructured materials has provided valuable insights into the potential beneﬁts and drawbacks of grain reﬁnement in the nanocrystalline(nc)regime,where the grain size is typically smaller than100nm,on various mechanical properties[8–11].However,the role of nanoscalegrainsInternational Journal of Fatigue27(2005)1147–1158/locate/ijfatigue0142-1123/$-see front matter q2005Elsevier Ltd.All rights reserved.doi:10.1016/j.ijfatigue.2005.06.035*Corresponding author.Tel.:C16172533320;fax:C16172530868.E-mail address:ssuresh@(S.Suresh).in inﬂuencing the resistance to fatigue of engineering metals and alloys has not been explored in sufﬁcient detail,despite the fact that the potential use of nanocrystalline materials in load-bearing engineering structures critically depends on their tolerance to the onset and progression of damage from cyclic loading.One factor contributing to this relative lack of information on the fatigue response of nc materials is the difﬁculty in producing truly nc metals and alloys(whose average as well as extreme grain dimensions are all smaller than about100nm)in sufﬁciently large quantities to facilitate‘valid’tests that comply with the American Society for Testing and Materials(ASTM)standards for minimum specimen dimensions and small-scale yielding for characterization using linear elastic fracture mechanics (LEFM).In an attempt to circumvent these difﬁculties, several studies[2–4]have employed ufc metals produced by equal-channel angular pressing(ECAP)to explore grain size effects on cyclic deformation.However,the high initial defect density inherent in these materials signiﬁcantly biases the fatigue response of the material.Furthermore, such processing methods are not presently amenable to produce truly nc metals in that the microstructure comprises a signiﬁcantly large fraction of grains whose dimensions exceed100nm;these grains typically dominate the overall mechanical response.The only study available to date on the fatigue response of fully nc metals is a preliminary report from the present work[5].This investigation was initiated with the objective of probing the effects of cyclic loading on the fatigue resistance of fully dense nc metals.The stress-life(S–N) fatigue response and the fatigue crack growth resistance of truly nc electrodeposited pure Ni was assessed,with the crack growth response further examined over a wide range of tensile load ratios.In order to assess the effects of grain size control and grain-boundary engineering on fatigue response,the results obtained for nc Ni are compared with those for ufc and mc metals and alloys,wherever feasible. In order to further assess the broad generality of the conclusions of this study,additional crack growth experiments were also conducted in a cryomilled ufc Al–Mg alloy and an equal channel angular pressed(ECAP) Grade2pure Ti,for which sizeable quantities of bulk specimens were available such that conventional fatigue testing could be employed.From these observations, mechanistic interpretations of the effects of grain size on fatigue response are developed,and possible strategies for grain size engineering in damage-critical applications are suggested.2.Materials and experimental methods2.1.Model material systemThe choice of a model material system,and correspond-ing fabrication technique,was predicated upon the following requirements:(1)the attainability of a true nc grain structure where the average and the largest grain size are both below100nm,(2)the ability to achieve a high level of material purity,with sufﬁcient reproducibility,and (3)the capability to produce a fully dense structure.Such requirements ensure that the material produced can accurately be described as nc.Particular attention was paid to maintaining a range of grain sizes well below 100nm to avoid having any larger‘outlier’grains dominate the fatigue response.While there are many methods currently available to fabricate nc materials[12–22],electrodeposited Ni was chosen as the model system for this investigation.It can be produced over a broad range of grain sizes,from the nc to the mc regime,and is capable of satisfying the above-mentioned requirements.Two notable advantages of the electrodeposition process are the ability to produce relatively large(in-plane)quantities of uniform,fully dense material(e.g.80!80mm),and the capacity to conﬁne the grain size to a narrow distribution(Fig.1). Although the attainable thickness resides in the millimeter range,samples in this investigation were purposely limited to a thickness of approximately100–150m m,to ensure through-thickness grain size uniformity and to avoid processing induced residual stresses.Electrodeposited Ni foils with two different grain sizes(nc Ni with an average grain size of20–40nm and a ufc equi-axed structure with an average grain size of approximately300nm)were procured from Integran Corporation,Toronto,Canada.2.2.Experimental detailsFor the purpose of minimizing processing-induced residual stresses,all fatigue specimens were extracted from the electrodeposited Ni foils by way of electro-discharge machining.Both low-and high-cyclefatigue Fig.1.Grain size distribution of electrodeposited pure nc Ni,illustrating the relatively narrow range of grain sizes,as well as the conﬁnement of all grains below the100nm range.T.Hanlon et al./International Journal of Fatigue27(2005)1147–1158 1148experiments were carried out in a laboratory air environ-ment(approximately258C and50%relative humidity).Full details of the stress-life set-up and experimental methods are available elsewhere[5,20].High cycle fatigue crack growth experiments for the nc,ufc,and mc Ni foils were conducted using single edge-notched specimens,where fatigue cracks were initiated in cyclic tension at load ratios ranging from 0.1to0.7at a cyclic frequency of10Hz(sinusoidal waveform).The specimens were39mm long,9.9mm wide,and100m m in thickness.Changes in crack length as a function of the number of fatigue cycles were monitored optically with a traveling microscope.The crack growth rate,d a/d N,was monitored as the length of the crack increased under a constant range of imposed cyclic loads.To ensure that small-scale yielding conditions prevailed,all data were collected such that the remaining uncracked ligament length was always at least twenty times greater than the maximum plastic zone size at the crack tip.In order to assess the overall generality of the trends observed in electrodeposited Ni,we studied the fatigue properties of two additional material systems,for which larger bulk specimens could be produced.First,a cryomilled Al-7.5Mg alloy was fabricated in billet form,50mm in diameter,and several inches in length(and was procured from the University of California,Irvine).This choice was motivated by the fact that grain size effects could be assessed in the ufc regime using specimens whose dimensions are sufﬁciently large(35!35!5mm)to meet the requirements of conventional standards for fatigue testing of bulk materials.A complete review of the Al-7.5Mg powder production and consolidation techniques is given in[23,24].In addition,ufc,ECAP pure Ti was also investigated,with direct comparisons made to its mc counterpart.Compact specimens,5mm in thickness,were extracted from extruded Al–Mg billets in the circumferential–radial (C–R)conﬁguration.The notch tip was machined to a radius of0.09mm and the specimen faces were polished to a mirrorﬁnish,with aﬁnal0.25m m polishing step.The through-holes,located on either side of the notch to pin-load the specimens,were machined after the fatigue pre-crack was introduced via cyclic compression loading[1,25, 26].Full details of the experimental technique are given in[5].For the study of fatigue of ECAP materials,billets of Grade2commercially pure mc Ti,40mm in diameter and 150mm in length,were subjected to eight ECAP cycles at a temperature of4258C,using a molybdenum disulphide lubricant and a die angle pact specimens,3mm in thickness,were extracted in the radial–longitudinal(R–L) conﬁguration from billets from the same batch exposed to either zero(mc Ti)or eight(ufc Ti)such pressing cycles.In-plane dimensions of the specimens measured33!31.75mm.The notch tip was machined to an initial radius of0.12mm,and subsequently sharpened to a radius of 0.03mm with a razor blade sprayed with a0.25m m diamond polishing suspension.A fatigue pre-crack was introduced in cyclic tension at load ratios of R Z0.1and0.3, at a cyclic frequency of10Hz(sinusoidal waveform)at room temperature.Crack growth was monitored in-situ with a telescopic video camera module,and ex-situ with an optical microscope.2.3.Shape factor considerationsIn order to characterize the variation of the fatigue crack growth rate as a function of D K,a proper evaluation of the shape factor,f(a/W),speciﬁc to the present specimen geometry(single-edge-notched tension specimen)and loading conﬁguration was required.For a rectangular plate of width,W,containing a through-thickness edge crack of length,a,it is commonly known that the stress intensity factor can be expressed by:K I Z sﬃﬃﬃapfaW;f aWZ1:99K0:41aWC18:7aW2K38:48aW3C53:85aW4ð1ÞHowever,the above expression is limited to loading conﬁgurations capable of producing a uniform stress throughout the specimen.Typically,this can be accom-plished by utilizing proper sample geometry(i.e.height, H O2W).Eq.(1)must be modiﬁed when a non-uniform stress is applied.In order to conduct experiments at a sufﬁciently high frequency,it was necessary to load the edge-notch speci-mens using rigid,as opposed to pin-loaded grips.In the latter,there exists a rotational degree of freedom,which produces an opening moment at the crack tip during tensile loading.Upon loading with such grips,a signiﬁcant degradation of the imposed sinusoidal waveform occurred at approximately3–5Hz.This was accompanied by machine resonance,which rendered the system unstable. Conversely,when the specimens were loaded with rigid grips(i.e.‘ﬁxed end displacement’loading),where the ends of the specimens were displaced by a constant amount,a clean sinusoidal waveform at the desired frequency of 10Hz was easily achieved.Several authors[27–30]have previously recognized that the application of Eq.(1)to aﬁxed-end displacement loading conﬁguration can substantially overestimate the mode I stress intensity factor.The boundary conditions associated with this type of loading are such that a closing bending moment is imposed on the crack,even during tensile loading,due to the lack of rotational freedom in the grips.This imposed closing moment effectively reduces the stress intensity factor,and therefore the driving force for crack propagation.To account for this relative reduction in D K,the shape factor,f(a/W),was evaluated viaﬁnite element analysis forT.Hanlon et al./International Journal of Fatigue27(2005)1147–11581149the speciﬁc loading geometry used in this investigation.In the analysis,the J -integral was calculated for a /W values ranging from 0.1to 0.7,fully encompassing the range investigated experimentally.Assuming that the stress intensity factor,K I ,has the form given in Eq.(1),the shape factor can be expressed as follows:f a W Z ﬃﬃﬃﬃﬃﬃﬃﬃJE a s r (2)where E is Young’s modulus and the relationship in Eq.(2)pertains to plane stress conditions.By applying a prescribed load,J can be systematically solved for over a range of crack lengths.In this manner,the shape factor can be fully characterized for the loading conﬁguration and specimen geometry of interest.To conﬁrm the accuracy of the analysis,a constant stress loading conﬁguration was modeled,and the resulting shape factor was found to compare well with the standard formula given in Eq.(1).Eq.(3)speciﬁes the shape factor determined from the analysis (detailed descriptions of which can be found in Ref.[20])for the ﬁxed end displacement case,which was employed in all calculations of the stress intensity factor used to interpret the fatigue crack growth results.f a W Z 2:03K 0:43a W C 3:96a W2K 6:04a W 3C 5:68a W4(3)3.Results3.1.Structure and tensile propertiesSpecimens of nc and ufc Ni were fully characterized in their as-received state by recourse to electron microscopy (Fig.2)and/or x-ray diffraction.The electrodeposited nc Ni had a columnar grain structure with an aspect ratio of 7–10,whereas the electrodeposited ufc Ni was nearly equiaxed.The micrographs in Fig.2a,b depict relatively defect-free initial structures,a critical factor in assessing the role ofgrain size on the overall fatigue response.Full details of theelectrodeposition process and resulting Ni structure are reported elsewhere [14,31,32].Tensile testing of nc and ufc Ni specimens revealed 0.2%offset yield strength values of 930and 525MPa,respectively,and strain to failure values of 3and 10%,respectively.The mc Ni had a yield strength of 180MPa,tensile strength of 450MPa and tensile strain to failure of 35%.Transmission electron microscopy (TEM)images show that the Al–Mg alloy investigated has a relatively equiaxed grain structure,with an average grain size of w 300nm [24].Its yield and tensile strengths were measured experimentally to be 540and 551MPa,respectively.The structure of the ufc ECAP Ti was also relatively equiaxed,with an average grain size of approximately 250nm (Fig.3)and a yield strength of 635MPa.Image analysis of the mc Ti revealed an average grain size of w 22m m,while its yield strength was measured at 430MPa.Full processing and property details are reported in [33].3.2.Fatigue response3.2.1.Stress-controlled fatigueThe effect of grain size on the fatigue resistance of initially smooth-surfaced pure Ni is shown in Fig.4in the form of stress-life fatigue curves,from our earlier report [5].It is evident that nc Ni has a slightly (and reproducibly)higher resistance to stress-controlled fatigue loading than ufc Ni.Additionally,the endurance limit of the mc Ni is signiﬁcantly below that of both the nc and ufcmaterial,Fig.2.Electron and optical micrographs of the (a)nc Ni,(b)ufc Ni,and (c)mc Ni investigated [5].Fig.3.(a)Optical micrograph showing the grain size distribution of the Grade 2pure mc Ti investigated (zero pressing cycles).(b)TEM image of the ufc ECAP Ti,subjected to 8pressing cycles at a temperature of 4258C.T.Hanlon et al./International Journal of Fatigue 27(2005)1147–11581150clearly illustrating the beneﬁcial effects of grain size reduction on the resistance to S–N fatigue.3.2.2.Fatigue crack growthThe variation in fatigue crack growth rate with respect to D K for pure Ni at load ratios of R Z0.1,0.3,and0.7is plotted in Fig.5.In order to enforce the assumptions inherent to LEFM,all data collection was truncated to incorporate only those data points corresponding to an uncracked ligament length of at least20times the plastic zone size at the tip of the crack during fatigue crack growth experiments.Due to the relatively limited strength of mc Ni, valid fatigue crack growth experiments using D K as the characterizing parameter were not possible under the abovementioned requirements.It is evident from Fig.5 that the resistance to fatigue crack growth is substantially lower in nc Ni at all levels of applied loading,over a wide range of load ratios.To circumvent the foregoing issues associated with the extent of crack-tip plasticity,a set of data was also collected relating the change in crack length to the number of fatigue cycles in mc,ufc,and nc Ni(Fig.6).Each material was subjected to identical initial loading conditions of D K Z 9.5MPa m1/2,R Z0.3,and a cyclic frequency of10Hz at room temperature.Fig.6clearly illustrates that the crack growth rate in the nc Ni is signiﬁcantly higher than that in the ufc and mc Ni.The effects of load ratio R on the fatigue crack growth response in nc and ufc Ni are plotted in Fig.7, where an increase in R leads to faster crack growth in both materials over the entire range of D K examined.These results are replotted in Fig.8where the stress intensity factor range D K required for a growth rate of10K6mmper Fig.5.Variation of the fatigue crack growth rate,d a/d N,as a function of D K for pure electrodeposited ufc,and nc Ni at load ratios(a)R Z0.1,(b)R Z0.3,and (c)R Z0.7,at a fatigue frequency of10Hz at room temperature.101010107StressRange(MPa)No. of Cycles to FailureFig.4.The effects of grain size from the micro to the nano-regime on thecyclic stress vs.total number of cycles to failure plot in pure Ni.(From[5]).T.Hanlon et al./International Journal of Fatigue27(2005)1147–11581151cycle in nc and ufc Ni is plotted as a function of R and maximum stress intensity factor,K max .The deleterious effect of grain reﬁnement on crack growth is evident in this ﬁgure.Further,discussion of the trends seen in this ﬁgure is taken up in a later section.To further explore the validity and generality of the above fatigue crack growth trends to ﬁne-grained metals and alloys produced by other processing methods,additional ufc materials fabricated via cryomilling (ufc Al-7.5Mg)and equal channel angular pressing (ufc pure Ti),for which larger bulk specimens could be procured,were examined.Since the solid solution Al-7.5Mg alloy234567800.20.40.60.81Load Ratio234567805101520255K max (MPa m 1/2)K (M P a m 1/2)K (M P a m 1/2)Fig.8.Stress intensity factor range D K required to induce a growth rate of 10K 6mm per cycle in nc and ufc Ni is replotted from the information in Fig.7as a function of R (top ﬁgure)and maximum stress intensity factorrange,K Ãmax (bottom ﬁgure).The solid lines in both ﬁgures show actual trend lines whereas the dashed lines denote assumed trends.Here it is assumed that there exists a critical R above which crack growth is unaffected by R .The ﬁlled symbols in the latter ﬁgure are experimental data and the open symbols are interpolated points extracted from the informationgiven in the upper ﬁgure.D K Ãand K Ãmax ,respectively,denote the limiting or threshold values of alternating and maximum values of stress intensity factor required for the particular growth rate of 10K 6mm per cycle.These values increase with increasing crack growth rate.The apparently detrimental effect of grain reﬁnement on crack growth is evident in bothﬁgures.V a r i a t i o n I n C r a c k L e n g t h (m m )Fatigue Cycles (In Thousands)Fig.6.Variation in crack length as a function of the number of imposed fatigue cycles for mc,ufc,and nc Ni subjected to an initial D K of 9.5MPa m 1/2,load ratio R Z 0.3,and cyclic frequency of 10Hz at roomtemperature.10–710–610–510–4234567891020d a /d N (m m /c y c le )d a /d N (m m /c y c le )K (MPa m 1/2)K (MPa m 1/2)10–710–610–510–4234567891020Fig.7.Fatigue crack growth rate as a function of D K at different load ratios for nc Ni (upper ﬁgure)and ufc Ni (lower ﬁgure).T.Hanlon et al./International Journal of Fatigue 27(2005)1147–11581152could only be fabricated via cryomilling,which ultimately results in a very ﬁne grain structure,a direct comparison to mc Al-7.5Mg could not be made.However,Al-5083is a close mc counterpart,and is often used for comparison purposes [24].Consistent with the results obtained for electrodeposited Ni,it was found that the ufc Al-7.5Mg fatigue crack growth rates over the entire d a /d N range,from threshold to ﬁnal failure,were substantially higher than those in the mc Al-5083(Fig.9).The threshold stress intensity factor range was also considerably lower in the ufc material.In addition,the critical value of D K at which catastrophic failure occurred was several times smaller in the cryomilled Al–Mg.Examination of the fracture surface of the ufc material revealed a signiﬁcant amount of cracking at inclusion particles [5],which were likely introduced during the cryomilling process.Such particles are believed to play a signiﬁcant role in the lowering of the critical value of D K .Fatigue crack growth response of pure mc and ECAP-processed ufc Ti was also fully characterized from threshold to ﬁnal failure.Fig.10shows effects of grain size on the variation of d a /d N as a function of D K at load ratios of R Z 0.1and 0.3.Here,grain reﬁnement from the mc to the ufc regime leads to a reduction in D K th by a factor of 2.5.The rate of fatigue crack propagation is more than an order of magnitude higher in the ufc Ti,over a wide range of applied loading,further reﬂecting the same trends captured in the electrodeposited Ni.4.DiscussionStudies of fatigue crack initiation in microcrystalline metals and alloys have long considered the critical role ofpurity,surface preparation and material strength in inﬂuencing the stress-life response.In nominally smooth-surfaced and defect-free fatigue specimens without notches and other stress concentrators,the initiation of cracks is known to occur in high-cycle fatigue at surfaces.An increase in the strength and hardness of the material,especially near the surface regions which serve as crack nucleation sites,is therefore considered to impart a greater resistance to fatigue crack initiation and hence to S–N fatigue.In the present experiments involving pure Ni and smooth laboratory specimens,this trend is seen to extend down to the nanocrystalline regime,where the higher strength and hardness of the nc material serves to provide a greater resistance to high cycle fatigue than the ufc and mc Ni.1Prior work on the fatigue crack growth resistance of conventional mc metals and alloys has shown (see,for example,[1]for a review of the literature)that predominantly crystallographic and stage I crack growth mechanisms arising at low D K levels of fatigue crack growth lead to microstructurally tortuous crack paths in the coarser grained materials.Grain reﬁnement in mc alloys then serves to reduce the extent of such crack path tortuosity,especially in the near-threshold regime where the cyclic plastic zone size is typically smaller than the average grain size and where the maximum crack-tip opening displacement can be markedly smaller than the surface asperity dimension.These seemingly small periodic deﬂections in the path of the crack can lead to changes in effective D K as a result of the net reduction in the local stress intensity factor range compared to the case of a perfectly straight crack front [7].In addition,when the10–810–710–610–510–410–310–2110d a /d N (m m /c y c le )∆K (MPa m 1/2)Fig.9.Variation of the fatigue crack growth rate,d a /d N ,as a function of the stress intensity factor range,D K ,for ufc cryomilled Al-7.5Mg at R Z 0.1–0.5at a fatigue frequency of 10Hz at room temperature.Also shown are the corresponding crack growth data for a commercial mc 5083aluminum alloy at R Z 0.33[14].10–810–710–610–510–410–311d a /d N (m m /c y c le )K (MPa m 1/2)Fig.10.Variation in fatigue crack growth rate as a function of D K for commercially pure mc Ti and ECAP ufc Ti at a fatigue frequency of 10Hz at room temperature.1Nanocrystalline metals produced by electrodeposition are known to contain hydrogen which is introduced during processing.However,in the present work,the effects of such impurities are not found to play a major role in introducing damage in nc Ni,as corroborated through high-resolution transmission electron microscopy [31].T.Hanlon et al./International Journal of Fatigue 27(2005)1147–11581153。

血管内皮生长因子及其受体与Cyclin E-CDK2在肝癌发生发展过程中的表达变化

血管内皮生长因子及其受体与Cyclin E-CDK2在肝癌发生发展过程中的表达变化王东;李笑岩;李红星;白咸勇【摘要】Objective To study the expression of vascular endothelial growth factor (VEGF) and its receptor, Cyclin E and CDK2 in the liver of the rat liver cancer model, and to explore the function and significance of VEGF, VEGFR1, Cyclin E and CDK2 in the liver cancer progression. Methods Based on the DEN-induced hepatocarcinogenesis model, the expressions of VEGFR1, Cyclin E and CDK2 were detected with immunohistochemical staining, and the quantity of VEGF in the serum was detected with enzyme linked immunosorbent assay. Results The quantity of VEGF in the serum of rats is lowest in the control group, and increased gradually in the experimental groups. The average optic density values of VEGFR1,Cyclin E and CDK2 protein expression had increasing trends during the development of liver cancer. The quantity of VEGF in the rat serum and the average optic density values of VEGFR1, Cyclin E and CDK2 protein were positively correlated (r = 0. 834, F = 42. 1274 , P＜0. 05). Conclusion Abnormal expression of VEGF and its receptor, promotes hepatocarcinogenesis and tumor development, which might be related with the abnormal expression of Cyclin E and CDK2 protein.%目的通过观察血管内皮生长因子(VEGF)及其受体和细胞周期蛋白E (Cyclin E)在肝癌模型大鼠肝脏中表达情况,探讨VEGF与细胞周期相关蛋白在肝癌发生发展过程中的作用.方法建立诱发性肝癌模型,采用酶联免疫吸附试验检测血清中VEGF量的变化,免疫组化技术检测VEGFR1、Cyclin E和细胞周期蛋白依赖性激酶(CDK2)的表达情况.结果血清中的VEGF含量在对照组中最低,在实验组中逐渐增多,以癌变期含量最高.VEGFR1、Cyclin E和CDK2蛋白表达的平均光密度值均随着肝癌的发生发展有增高的趋势,大鼠血清中的VEGF量与肝脏组织中VEGFR1、Cyclin E和CDK2蛋白表达的平均光密度值随着肝癌的发生发展呈正相关(r=0.834,F=42.1274,P<0.05).结论 VEGF及其受体VEGFR1在肝癌发生发展中异常表达,促进肝癌的发生发展,可能与Cyclin E、CDK2细胞周期蛋白异常表达有关.【期刊名称】《中国组织化学与细胞化学杂志》【年(卷),期】2011(020)004【总页数】6页(P310-315)【关键词】诱发性肝癌;VEGF;VEGFR1;Cyclin E;CDK2【作者】王东;李笑岩;李红星;白咸勇【作者单位】滨州医学院组织胚胎学教研室,山东烟台,264000;滨州医学院组织胚胎学教研室,山东烟台,264000;滨州医学院组织胚胎学教研室,山东烟台,264000;滨州医学院组织胚胎学教研室,山东烟台,264000【正文语种】中文【中图分类】R329肝癌是死亡率仅次于胃癌、食道癌的第三大常见恶性肿瘤,它的发生发展是一个多因素、多阶段的过程[1]。

MEKC血样中药物,外文翻译

MEKC血样中药物.第1页原始文件沃尔夫冈·BuchbergerÆ马蒂亚斯Ferdig鲁道夫·索默Æ翠THANH VO 雷帕霉素在人体血液中的微量分析胶束电动色谱收稿日期：2004年3月16日/日期：11 2004/5月接受日期：2004年5月18号/发表时间：2004年7月20日Ó施普林格出版社2004年摘要毛细管电泳法与UV检测波长为278 nm的分析已经开发在人类的免疫抑制剂雷帕霉素（西罗莫司）血微克每升的水平低。

分离有在酸性载体电解质含有已取得十二烷基硫酸钠和30％（体积/体积）腈。

为样品净化和富集，离线固固相萃取的步骤中使用的基于二氧化硅的反相材料和毛细管聚焦技术采用。

后者允许注射增加扩大样本量，而不会过度带。

虽然这种新方法是不敏感，比现有的质谱联用液相色谱程序法，它是完全适合常规分析的RA-pamycin血液中这种药物治疗的患者。

最后，但并非最不重要的，低的成本使一个有吸引力的建立的方法替代。

关键词雷帕霉素Æ胶束电动血色谱Æ分析介绍雷帕霉素（西罗莫司）是三烯大环内酯类抗生素具有抗真菌，消炎，抗肿瘤和这是目前使用的免疫抑制特性用于预防器官移植排斥反应的反式种植园。

雷帕霉素的结构图给出。

1。

雷怕霉素已被证明是阻止T 细胞活化和扩散以及激活P70 S6文FK-506结合激酶，表现出很强的结合蛋白质。

它也抑制活性的蛋白质哺乳动物雷帕霉素靶蛋白（mTOR）的，该功能蒸发散在促进肿瘤生长的信号转导通路。

雷帕霉素结合到受体蛋白质（FKBP12），和的rapamycin/FKBP12复杂的结合mTOR的并阻止mTOR的互动与目标蛋白质在这一信号通路。

对于栽种的器官的患者是必不可少的监测雷帕霉素在他们的血液中的水平。

决定雷帕霉素在血液中的中断可以确定迄今已完成高效液相色谱紫外检测器[1-4]，最近由高效液相色谱法，连字符和质谱（MS）[5-8]。

nature.

Cancer CellArticleEnhancing Tumor-Speciﬁc Uptake of the Anticancer Drug Cisplatin with a Copper ChelatorSeiko Ishida,1,2Frank McCormick,1Karen Smith-McCune,1,3and Douglas Hanahan1,2,*1Helen Diller Family Comprehensive Cancer Center2Diabetes Center and Department of Biochemistry&Biophysics3Department of Obstetrics,Gynecology,and Reproductive SciencesUniversity of California,San Francisco,San Francisco,CA94143,USA*Correspondence:dh@DOI10.1016/r.2010.04.011SUMMARYUptake of the anticancer drug cisplatin is mediated by the copper transporter CTR1in cultured cells.Here we show in human ovarian tumors that low levels of Ctr1mRNA are associated with poor clinical response to platinum-based ing a mouse model of human cervical cancer,we demonstrate that combined treatment with a copper chelator and cisplatin increases cisplatin-DNA adduct levels in cancerous but not in normal tissues,impairs angiogenesis,and improves therapeutic efﬁcacy.The copper chelator also enhances the killing of cultured human cervical and ovarian cancer cells with cisplatin.Our results identify the copper transporter as a therapeutic target,which can be manipulated with copper chelating drugs to selectively enhance the beneﬁts of platinum-containing chemotherapeutic agents.INTRODUCTIONCisplatin is widely used for the treatment of solid tumors.It is a component ofﬁrst-line treatment for testicular,ovarian, cervical,endometrial,bladder,head and neck,lung,and gastro-esophageal cancers(Hussain et al.,2008;Omura,2008;Gal-lagher et al.,2008;Yang et al.,2009;Lustberg and Edelman, 2007;Fischer and Arcaro,2008;Cen and Ajani,2007).It is also used as a second-or third-line treatment for prostate and pancreatic cancers,metastatic cancer of the breast,and for melanoma and gliomas(Carrick et al.,2004;Oh et al.,2007; Saif and Kim,2007;Atallah and Flaherty,2005;Glioma Meta-Analysis Trialists Group,2002).The dosage and efﬁcacy of cisplatin,however,are limited by its side effects,the most prom-inent being nephrotoxicity(Pabla and Dong,2008).Cisplatin exerts its cytotoxic effect by forming an intrastrand crosslink on DNA(Jamieson and Lippard,1999).Studies of cisplatin-resistant cell lines have revealed drug uptake as a crit-ical step that governs cisplatin sensitivity in vitro(Hall et al., 2008).A yeast genetic screen for cisplatin-resistant mutants identiﬁed the copper transporter CTR1as a major mediator of cisplatin uptake in yeast and mouse cells(Ishida et al.,2002). Upon exposure to excess copper,Ctr1protein undergoes endo-cytosis and degradation(Ooi et al.,1996;Petris et al.,2003). Pretreatment of cells with high copper results in decreased cisplatin uptake and increased resistance to this drug in a CTR1-dependent manner(Ishida et al.,2002),whereas chelating copper with bathocuprione disulphonate has opposite effects(S.I.and I.Herskowitz,unpublished data).These results demonstrated that cisplatin uptake can be modulated by copper levels in vitro through Ctr1p.In this study,we examined the in vivo roles of Ctr1and copper in cisplatin uptake and response using a mouse model of human cervical cancer.Cervical cancer is the second most common cause of cancer-related death among women worldwide. Cisplatin has been one of the most active single agents and a component of combined-agent chemotherapy regimens for patients with advanced cervical cancer.The response rates for cisplatin alone range from13%–19%,with an overall survival time of6–9months;the current combined-agent therapy is only modestly better,with response rates of27%–36%,with overall survival of8–10months(Tao et al.,2008).Cervical cancer574Cancer Cell17,574–583,June15,2010ª2010Elsevier Inc.is caused by integration and persistent expression of genes of human papilloma viruses(HPV),of which HPV type16is detected in a majority of cases of invasive carcinoma(Castell-sague,2008).In the mouse model of cervical carcinoma, K14-HPV16/E2,the HPV16-encoded oncogenes are under the control of the keratin14promoter,resulting in targeted onco-gene expression in keratin14-positive squamous epithelial cells, a natural host-cell target for HPV infection in humans(Arbeit et al.,1996).When the estrogen levels of K14-HPV16females are maintained at a level near their natural peak during estrus with time-released implants of17b-estradiol(estrogen,E2),the female mice(referred to as HPV16/E2mice hereafter)develop progressive cervical intraepithelial neoplasia(CIN)lesions at the transformation zone,the site implicated in the genesis of human cervical cancer(Elson et al.,2000).By6months of age, these CIN lesions progress to invasive squamous cell carcinoma in90%of the animals.Both male and female K14-HPV16mice also develop dysplasia of the skin,and skin tumors arise later, with50%penetrance by12months of age(Coussens et al., 1996).Using this mouse model of human cervical carcinoma,HPV16/ E2,we ask whether a copper chelator selectively increases cisplatin uptake and killing of tumor cells without affecting normal organs,thereby increasing the antitumoral efﬁcacy of cisplatin.We also analyze human ovarian tumors for expression of the copper transporter Ctr1to evaluate its association with clinical response to platinum drugs,and assess the capability of the copper chelator to enhance killing of human ovarian and cervical cancer cell lines with cisplatin.RESULTSHuman Ctr1mRNA Levels in Tumors Are Associatedwith Clinical Response to Platinum-Based Chemotherapy in Ovarian CancerPrevious studies using mouse embryonicﬁbroblasts(MEFs) deleted for the mouse Ctr1gene demonstrated a correlation between Ctr1gene dosage and cisplatin sensitivity in vitro:cells that are heterozygous for a knockout allele of mouse Ctr1were more resistant than wild-type cells to cisplatin,and homozygous mutant cells were more resistant than heterozygous cells (Ishida et al.,2002).Theseﬁndings raised the hypothesis that Ctr1mRNA levels correlate with and thus explain differential clin-ical responses to platinum drugs.To address this possibility,we analyzed tumor specimens and the clinical course of15ovarian cancer patients.The standard treatment for advanced ovarian cancer is primary cytoreductive surgery followed by chemo-therapy that includes a platinum drug.Resistance to platinum-based chemotherapy remains a primary factor in disease recurrence.Figure1A lists the baseline characteristics of the 15patients,all with advanced-stage(stage III or IV),serous epithelial ovarian carcinoma.The tumors were of mixed grade (Figure1A).Following optimal cytoreductive surgery resulting in residual tumors of1cm or less in diameter,the patients received regimens containing either cisplatin or carboplatin, and paclitaxel.Prior clinical studies have revealed equal efﬁcacy for the two platinum compounds(Ozols et al.,2003;du Bois et al.,2003).The mechanisms of resistance are thought to be similar for the two drugs,which share the same putative trans-porter CTR1(Holzer et al.,2006).Of the15patients,9proved to be sensitive to platinum-based therapy(no evidence of recur-rence within6months after completion of primary therapy), whereas the remaining6were refractory or resistant(progres-sion or recurrence within6months of completion of therapy, as indicated by an increase in CA-125levels,or by CT,ultraso-nography,or second-look laparoscopy)(see Figure S1A available online).RNA was prepared from these patients’tumor samples,and Ctr1mRNA levels were measured by quantitative reverse-transcriptase polymerase chain reaction(RT-PCR). The mean relative Ctr1mRNA level was substantially higher in platinum-sensitive patients than in refractory/resistant patients (43%versus18%of internal control gene expression,p= 0.016)(Figure1B).Expression levels of the internal control gene human GUS,which encodes b-glucuronidase,were equiv-alent between the two patient groups(average±standard deviation[SD]Ct in sensitive group was28.2±0.8versus 28.1±1.3in resistant/refractory group,p=0.69)(Figure S1B). Theseﬁndings suggest that Ctr1expression in ovarian tumors is associated with platinum drug response.In order to validate theseﬁndings in an independent data set, we used clinical and array-based expression data from The Cancer Genome Atlas(TCGA)that are deposited at the Data Coordinating Center for public access(http://cancergenome. /).We deﬁned a subset of91patients with stage III or IV serous epithelial ovarian cancer who had undergone a cytore-ductive surgery followed by adjuvant chemotherapy consisting of a platinum drug and a taxane(/ tcga/homepage.htm)(Figure1C).Results for Ctr1expression were extracted from expression data that were produced using Affymetrix HT Human Genome U133A Array Plate Set (HT_HG-U133A)(/tcga/dataAccess Matrix.htm).The patients were divided into two groups by the median Ctr1expression value,resulting in45‘‘high Ctr1’’(above median)and46‘‘low Ctr1’’(below median)patients(Figure S1C). High expression of Ctr1was associated with increased disease-free survival(Figure1D;log-rank statistic8.67,hazard ratio0.53, p=0.003),supporting our hypothesis that Ctr1levels are prognostic of treatment outcome.Theseﬁndings suggest that mRNA levels of the human copper transporter Ctr1in tumors are clinically relevant to the efﬁcacy of platinum therapy.Ctr1mRNA Levels Correlate with Cisplatin Adduct Levels in Mouse TissuesTo address the in vivo role of CTR1in cisplatin uptake,we exam-ined the levels of Ctr1mRNA and cisplatin-induced DNA adducts in various tissues of HPV16/E2,the mouse model of human cervical cancer.Overall,the levels of cisplatin adducts corre-lated with Ctr1mRNA in most organs tested—skin,lung,liver, pancreas,and uterus(Figure2).Kidney contained the highest amount of cisplatin adducts and Ctr1mRNA,which is interesting in light of the prevalence of nephrotoxicity among cisplatin-treated patients.The colon contained fewer adducts relative to Ctr1mRNA levels.In the duodenum and intestine,CTR1protein is localized to the apical surface facing the intestinal lumen and in intracellular compartments(Nose et al.,2006;Kuo et al.,2006). It is possible that such localized CTR1in the digestive tract is not involved in the uptake of cisplatin,which is administered intravenously to patients,and in this experiment was suppliedCancer CellModulating Copper Improves Cisplatin EfﬁcacyCancer Cell17,574–583,June15,2010ª2010Elsevier Inc.575intraperitoneally.Notably,the cancerous HPV16/E 2cervix contained more adducts relative to Ctr1mRNA levels.Expres-sion in the cervix of the internal control gene mouse L19,to which Ctr1mRNA was normalized,was equivalent to that of othertissues (average ±SD tissue L19Ct 23.5±1.1versus 23.3in the cervix)(Figure S2A),indicating that the observed low levels of Ctr1mRNA in the cervix is not a result of higher internal control gene expression in the tumors.Despite the low Ctr1mRNA levels,CTR1protein was substantially expressed in the can-cerous cervix (Figure 3D),suggesting increased protein syn-thesis and/or reduced protein turnover in this neoplastic tissue.We also measured mRNA levels of the copper exporters,ATP7A and ATP7B ,which are involved in exporting cisplatin in vitro (Komatsu et al.,2000;Samimi et al.,2004).Consistent with prior reports (Bull et al.,1993;Tanzi et al.,1993),ATP7B was primarily expressed in the liver (Figure S2B).We were able to detect ATP7A mRNA in all the tissues tested,but did not observe an inverse correlation between ATP7A mRNA levels and tissue platinum levels,further supporting our proposition that the copper transporter Ctr1is the major determinant for cisplatin accumulation in vivo as well.The Copper Chelator Tetrathiomolybdate Increases Cisplatin Adduct Levels in Tumors but Not Normal OrgansHaving observed an in vivo correlation between levels of the copper transporter Ctr1and cisplatin adducts in the majority of tissues examined,we next asked whether reducing systemic copper could result in increased cisplatin adducts in tumors.We treated HPV16/E 2mice with the copper chelator tetrathio-molybdate (TM),which is currently granted orphan designation in the United States and Europe for the treatment of Wilson’s disease,an autosomal recessive disorder characterized by excess copper deposition in various organs.TM reduces bio-available copper levels,primarily by forming a tripartite complex with ingested copper and protein,preventing the absorptionofFigure 1.Low Levels of Ctr1mRNA in Ovarian Tumors Are Associated with Poor Clinical Outcome(A)Baseline characteristics of the 15patients with advanced-stage epithelial ovarian carcinoma,whose tumors were analyzed for relative expres-sion of the Ctr1copper transporter.(B)Tumor Ctr1mRNA levels in platinum sensitive and refractory/resistant ovarian cancer patients.Bars represent mean values.Ctr1mRNA levels are expressed as a percentage relative to the internal control of human GUS mRNA.(C)Characteristics of 91ovarian cancer patients from The Cancer Genome Atlas study.(D)Kaplan-Meier estimate of disease-free survival for 91advanced-stage ovarian cancer patients according to Ctr1expression levels.Patients whose tumors expressed Ctr1mRNA at levels higher than the median value were classiﬁed as ‘‘High Ctr1’’(red;45patients);the remaining patients were labeled ‘‘Low Ctr1’’(blue;46patients).The vertical hash marks (at 7and 10months,High Ctr1)repre-sent censored patients who were disease-free at the last follow-up time indicated.See also Figure S1.Figure 2.Levels of Ctr1mRNA and Cisplatin Adducts in Various Organs of the HPV16/E 2Female MiceOrgans were harvested from three 6-month-old HPV16/E 2females and Ctr1mRNA levels were determined by quantitative RT-PCR (hatched bars).Mouse L19mRNA levels were used as internal controls.For measurement of cisplatin adducts,genomic DNA was puriﬁed from each organ (n =8)2hr after 6mg/kg cisplatin was injected.Platinum was measured by ICP-MS and normalized to the amount of DNA (solid bars).Error bars represent the standard error of the mean.See also Figure S2.Cancer CellModulating Copper Improves Cisplatin Efﬁcacy576Cancer Cell 17,574–583,June 15,2010ª2010Elsevier Inc.copper in the intestine (Brewer and Merajver,2002).We deter-mined the maximally tolerated daily dose of TM for the HPV16/E 2female mice to be 1mg (data not shown).When the mice are subjected to this daily dose of TM,plasma copper levels start decreasing after 3days,and reach a plateau after 2weeks.We did not observe any side effects during 3weeks of daily TM treat-ment,at which point the plasma copper levels had decreased by 23%(Figure 3A).This TM dosage resulted in a 2.7-fold increase in cisplatin adduct levels in the cancerous cervix of the HPV16/E 2mice (7.2m g Pt/g DNA in control versus 19.8m g Pt/g DNA in TM-treated mice,p =0.02),but not in the wild-type cervix (Figure 3B).Such a difference in the number of adducts is considered signiﬁcant because it can result in a 1000-fold differ-ential in cell viability in vitro (Hall et al.,2008).Notably,this TM regimen did not increase adducts in the kidney (Figure 3B),which is the major organ affected by cisplatin toxicity,norinFigure 3.The Copper Chelator Tetrathiomolybdate Enhances Cisplatin Uptake into the Cervix of HPV16/E 2Mice(A)Plasma copper levels in six-month-old HPV16/E 2females after a 3week daily treatment with 1mg tetrathiomolybdate (TM).Bars represent mean values.(B)Cisplatin adduct levels in the cervix,skin,and kidney of six-month-old HPV16/E 2females (n =12)and wild-type females (n =5)that received either the three-week TM treatment (solid bars)or no treatment (hatched bars).6mg/kg cisplatin was injected two hours before organs were harvested.Adduct levels were determined as in Figure 2.Error bars represent the standard error of the mean.(C)CTR1immunohistochemistry reveals positive cells (green)in tumors (T)of the 6-month-old HPV16/E 2cervix (right panel)but not in the epithelium (E)of the wild-type cervix (left panel).DAPI staining is shown in blue.The panels shown are representative images of 12ﬁelds in three tissue sections per mouse collected from three mice.Scale bars represent 100m m.(D)Tissue CTR1protein levels are elevated in the neoplastic cervix,but unchanged in the context of TM treatment.Wild-type and HPV16/E 2mice were either treated with 1mg TM daily for 3weeks prior to being sacriﬁced at 6months of age,or did not receive such treatment.A total of 5m g membrane extracts from tissues was pooled from three mice and applied on each well.(E)CTR1protein (green)is predominantly localized in intracellular compartments in cervical carcinoma cells,and is unchanged by TM treatment.HPV16/E 2mice were either treated with 1mg TM daily for 3weeks prior to being sacriﬁced at 6months of age,or did not receive such treatment.Tissue sections were immunostained with anti-CTR1(green).Nuclei are stained with DAPI (blue).The images were collected using a 20x objective lens on a ﬂuorescent microscope,and are representative of >500cells in 12sections from six tumors.Scale bars represent 5m m.Cancer CellModulating Copper Improves Cisplatin EfﬁcacyCancer Cell 17,574–583,June 15,2010ª2010Elsevier Inc.577Figure 4.Preclinical Trials Assessing Therapeutic Effects of the Copper Chelator Tetrathiomolybdate in Combination with Cisplatin on De Novo Cervical Carcinomas(A)Schematic diagrams of each therapeutic arm.Cohorts of 14female HPV16/E 2mice were subjected either to a 3week oral administration of 1mg tetrathio-molybdate (TM)daily starting at 5months of age,or to a single intraperitoneal injection of 6mg/kg cisplatin at 5.75months of age,or to both.All the mice were sacriﬁced at 6months of age.Cancer CellModulating Copper Improves Cisplatin Efﬁcacy578Cancer Cell 17,574–583,June 15,2010ª2010Elsevier Inc.other noncancerous organs of the mice.Interestingly,in the absence of TM treatment,the kidneys of HPV16/E 2mice accu-mulated 50%more platinum adducts than those of wild-type mice (Figure 3B).This increase in the adduct level in the kidneys of HPV16/E 2mice may be due to renal failure from urinary tract obstruction,which is not uncommon both in cervical cancer patients and in this mouse model.TM treatment also produced a marginal 1.4-fold increase in the adduct level in the dysplastic skin of the HPV16/E 2mice (2.8m g Pt/g DNA in control versus 3.9m g Pt/g DNA in TM-treated mice,p =0.03),but not in the skin of wild-type mice,suggesting that neoplastic tissues may become more sensitive to changes in copper levels.Notably,we found that CTR1protein is highly expressed in tumors of the HPV16/E 2cervix but not in the epithelium of the wild-type cervix,the origin of cervical carcinoma (Figure 3C).We were,however,unable to detect any change in Ctr1mRNA (data not shown)or protein levels (Figure 3D)in response to TM treatment in the neoplastic HPV16/E 2cervix or in normal tissues.In tumor cells of the HPV16/E 2cervix,CTR1protein was predominantly localized in intracellular compartments;CTR1remained intracel-lular in TM-treated HPV16/E 2cervix (Figure 3E).It is possible that the modest change in systemic copper (23%reduction)by TM (Figure 3A)is not sufﬁcient to cause measurable changes in the levels of CTR1protein,which is already high in these tumors,or in its localization.However,TM treatment resulted in in-creased cisplatin uptake in the HPV16/E 2cervix (Figure 3B),demonstrating that even a small reduction in systemic copper levels can increase the transport activity of CTR1in tumors.These data indicate that expression of the copper transporter CTR1is upregulated in neoplastic tissues compared to cognate normal tissues,and that its transport activity could be further elevated selectively in tumors by reducing systemic copper.The Copper Chelator Tetrathiomolybdate Enhances Cisplatin EfﬁcacyFor assessment of possible therapeutic beneﬁts of using TM to enhance the uptake and consequent cytotoxicity of cisplatin in tumors,cohorts of 14female HPV16/E 2mice were subjected either to a 3week regimen of oral 1mg TM daily starting at 22weeks of age,or to a single intraperitoneal injection of 6mg/kg cisplatin at 25weeks of age,or to both (Figure 4A).At 22weeks,all HPV16/E 2females have high-grade dysplasias (CIN 2/3),and about half have invasive cervical carcinomas.Animals from all treatment arms were sacriﬁced at 26weeks of age,when 90%have cervical carcinomas.Tumor volumes were calculated by measuring tumor areas on serial sections.Cisplatin treatment reduced tumor volumes to 55%of untreated cervix (2.7mm 3in control versus 1.5mm 3with cisplatin,p =0.01)(Fig-ure 4B).In the cisplatin-treated cervix,cells had more abundant cytoplasm that appeared keratinized (Figure 4C).Nuclei were enlarged and bizarre in shape,characteristic features of human cervical cancer treated with radiation (Robboy et al.,2002),which causes DNA damage.TM alone also modestly decreased tumor volumes in the cervix,to 65%of controls (2.7mm 3in control versus 1.7mm 3with TM,p =0.04)(Figure 4B).In the TM-treated cervix,the center of a tumor mass was necrotic and ﬁlled with inﬂammatory cell inﬁltrates (Figure 4C).The percentage of area covered with meca-32-positive endothelial cells in the TM-treated tumors was 2.6-fold less compared with those from untreated mice (Figures 4D and 4E),consistent with prior obser-vations that TM is antiangiogenic (Brewer and Merajver,2002).When cisplatin was combined with the TM treatment,tumor volume decreased to 19%of untreated cervix (2.7mm 3in control versus 0.5mm 3with TM and cisplatin,p <0.0001)(Figure 4B).Microscopically,the cervical carcinoma from the combination arm showed histologic features of the effects seen with both cisplatin and TM monotherapies:enlarged nuclei,necrotic centers,inﬂammatory cell inﬁltrates,and reduced vascularity (Figures 4C and 4D).These results demonstrate that TM enhances efﬁcacy of cisplatin synergistically.The Copper Chelator Tetrathiomolybdate Increases Cisplatin Sensitivity of Cancer CellsSeeking to separate the antiangiogenic effect of TM on tumor vasculature from its effect on tumor cell uptake of cisplatin,we asked whether TM treatment increases cisplatin sensitivity and uptake in cultures of SiHa cells,which are derived from an HPV16-positive human cervical carcinoma (Friedl et al.,1970).Plasma TM levels in the HPV16/E 2females range from 150m M to 4m M,over the period from 3to 24hr after oral administration of 1mg TM,respectively (data not shown).We determined that 24hr incubation with up to 10m M TM does not inhibit prolifera-tion of SiHa cells (Figure 5A).When SiHa cells were incubated with 10m M TM for 24hr and then subjected to a 2hr cisplatin treatment,we observed a 2-fold increase in cisplatin sensitivity and a 45%increase in cisplatin accumulation inside TM-treated cells compared to cells that were not preincubated with TM (Figures 5B and 5C).We also assessed effectiveness of the(B)Tumor volumes in the cervix of 6-month-old HPV16/E 2mice from each treatment arm.Approximately 30sections per mouse from cohorts of 14female HPV16/E 2mice were analyzed as described in Experimental Procedures .Error bars represent the standard error of the mean.(C)Representative H&E-stained cervices from each treatment arm,with the treatment indicated below each panel.Cells are shown at a higher magniﬁcation on the right.The magniﬁcation is scaled by a bar below each panel indicating 50m m.In the nontreated cancerous cervix (top left),invasive squamous cell carcinomas contain large keratinized cells with hyperchromatic nuclei.In the cisplatin-treated cervix (top right),the cancer cells have more abundant cytoplasm that appears keratinized,with enlarged,bizarre-shaped nuclei.In the TM-treated cervix (bottom left),the center of a tumor mass is necrotic and ﬁlled with inﬂammatory cell inﬁltrates.The cervix treated with both TM and cisplatin (bottom right)shows histological features of the effects seen with both cisplatin and TM monotherapies:enlarged nuclei,necrotic centers,and inﬂammatory cell inﬁltrates.The panels shown are representative of four ﬁelds in 30tissue sections collected from cohorts of 14HPV16/E 2mice.(D and E)Vascular density of tumors after treatment.15tumor sections from ﬁve mice per treatment arm were stained with a pan-endothelial cell antibody Meca-32,followed by FITC-conjugated secondary antibody to visualize tumor endothelial cells.Images were collected from ﬁve to seven ﬁelds per section,and vascular density was determined using the Metamorph Angiogenesis Tube Formation program.Data are expressed as the percentage of area covered with Meca-32-positive tube structures (E).Error bars represent the standard error of the mean.Representative images of Meca-32(green)staining in the cervix of control and TM-treated HPV16/E 2mice are shown in (D).Nuclei are stained with DAPI (blue).Scale bars represent 50m m.Cancer CellModulating Copper Improves Cisplatin EfﬁcacyCancer Cell 17,574–583,June 15,2010ª2010Elsevier Inc.579Figure 5.Effects of the Copper Chelator Tetrathiomolybdate (TM)on Cisplatin Sensitivity in Cultured Mammalian Cells(A)Effect of TM on cell proliferation in human cervical carcinoma cells.SiHa cells were plated in triplicates and treated with 0,1,10,or 100m M TM for 24hr.Numbers of cells before and after the treatment are shown.Error bars represent the standard error of the mean.(B and C)Effect of TM on cisplatin sensitivity and accumulation in human cervical carcinoma cells.SiHa cells were plated in triplicates and either pretreated with 10m M TM or mock treated for 24hr before a 2hr incubation with cisplatin.For cell survival (B),data are expressed as percentage of viable cells compared with control cultures not exposed to cisplatin.For determining cisplatin accumulation (C),cells were lysed and,after a centrifugation,the supernatant was used to determine the platinum content and protein concentration.Cellular platinum readings were normalized to protein concentration.Mean values of the triplicates are shown.Error bars represent the standard error of the mean.(D)TM increases cisplatin sensitivity in human ovarian cancer cells.Cells were plated in triplicates and either pretreated with 10m M TM or mock treated for 6hr (A2780),12hr (SKOV3),or 24hr (OVCA4)before cisplatin treatment.Data are expressed as percentage of viable cells compared with control cultures not exposed to cisplatin.(E)TM increases cisplatin sensitivity in a Ctr1-dependent manner.Isogenic mouse embryonic ﬁbroblasts that are either wild-type or homozygous for a knockout allele of Ctr1were either pretreated with 10m M TM or mock treated for 24hr before a 2hr incubation with cisplatin.The percentage of viable cells compared with control cultures not exposed to cisplatin is shown.Cancer CellModulating Copper Improves Cisplatin Efﬁcacy580Cancer Cell 17,574–583,June 15,2010ª2010Elsevier Inc.combination therapy for ovarian cancer,using three human ovarian cancer cell lines derived from patients’tumors,A2780, SKOV3,and OVCA4(Eva et al.,1982;Fogh and Trempe, 1975).These cells had different growth properties:the doubling time was16hr for A2780,24hr for SKOV3,and36hr for OVCA4 (data not shown).We determined that incubation of these cells with10m M TM did not have an antiproliferative effect for up to 6hr in A2780,12hr in SKOV3,and24hr for OVCA4(data not shown).When these cells were subsequently treated with cisplatin,they exhibited1.8-to2.2-fold increased sensitivity compared with cells that were not preincubated with TM (Figure5D),demonstrating that TM can increase cisplatin sensi-tivity of ovarian cancer cells as well.To establish that the effects of TM on enhancing cisplatin-mediated killing were dependent on Ctr1,we compared isogenic MEFs that were either wild-type or deleted for the Ctr1gene(Lee et al.,2002).The ability of TM to increase cisplatin killing proved to be dependent on Ctr1:in wild-type MEFs,the percentage of cells that survived a2hr treatment with1mM cisplatin was reduced to0.16%in cultures pretreated with TM,as compared with1%in control cultures(Figure5E).In contrast,the percentage of cells surviving cisplatin treatment was not affected by TM treatment in Ctr1null MEFs(Figure5D).These results further support the role of the copper chelator TM and the copper transporter CTR1in enhancing cisplatin uptake and killing of cancer cells.DISCUSSIONOur studies in a mouse model of HPV16-induced cervical carci-noma demonstrate that the copper chelator tetrathiomolybdate has a synergistic antitumor effect when combined with cisplatin treatment,by increasing uptake of cisplatin into tumors but not normal tissues while concomitantly inhibiting tumor angiogen-esis.We show that cisplatin adduct levels in various mouse organs correlate with mRNA levels of Ctr1,supporting its in vivo involvement in cisplatin uptake.Pretreatment of cultured human cervical and ovarian cancer cells with TM resulted in increased cisplatin sensitivity,encouraging applicability to human cancers.TM-enhanced cisplatin killing was dependent on Ctr1because mouse cells deleted for the Ctr1gene were not sensitized to cisplatin by TM,in comparison to isogenic cells with wild-type Ctr1.Notably,low levels of Ctr1mRNA in human ovarian tumors are associated with decreased disease-free survival after platinum-based therapy in patients,suggesting that Ctr1may serve as a clinical predictor of response to plat-inum agents.Previous genetic experiments with yeast and mouse cells demonstrated that the copper transporter Ctr1is a major determinant of cisplatin uptake and sensitivity,which could be modulated by copper levels.These in vitro results suggested a strategy to increase cisplatin uptake,but a challenge was to ﬁnd a clinically feasible method to do so selectively in tumors, without affecting normal organs to avoid untoward toxicities of cisplatin.Surprisingly,we were able to achieve this goal by reducing systemic copper levels with a copper chelator in a mouse model of cervical cancer.We do not fully understand the mechanism by which reduced systemic copper leads to increased cisplatin uptake only in tumors.Copper is essential for a variety of key cellular processes such as respiration,free radical detoxiﬁcation,and iron uptake(Kim et al.,2008).Cancer cells may have a greater demand for copper than normal cells for proliferation and survival.Indeed,the copper transporter CTR1was more highly expressed in cervical carcinoma of the HPV16/E2mice than in the wild-type cervix.Concordant with our observation is a recent demonstration using microbeam synchrotron X-rayﬂuorescence,showing that copper is concentrated in tumor regions of tissue specimens obtained from invasive ductal carcinoma of the breast(Farquharson et al.,2008).Although pretreatment of mice with the copper chelator TM resulted in increased cisplatin sensitivity and adducts in tumors, we were unable to detect any increase in the levels of Ctr1mRNA or protein.In vitro,elevated extracellular copper causes endocy-tosis and degradation of human CTR1protein(Petris et al., 2003).We did not,however,observe any change in CTR1protein localization in the TM-treated cervix of HPV16/E2mice.It is possible that copper starvation enhances cisplatin transport activity by changing the conformation of CTR1,allowing more cisplatin to be transported into cells.Exogenous copper has been shown to induce structural rearrangements in yeast Ctr1p (Sinani et al.,2007).Further biochemical studies will be required to elucidate the mechanism by which copper chelation causes increased cisplatin uptake.Drug-induced toxicity is a common cause for discontinuation or dose reduction of chemotherapeutic drugs during cancer treatment.Cancers that are deemed‘‘resistant’’to a chemother-apeutic drug fail to respond to the dosage tolerated by patients without untoward side effects.These‘‘resistant’’cancers might in principle still respond to a higher dosage if toxicities were not manifest.By reducing bioavailable copper with the copper chelator TM,we were able to increase cisplatin activity in highly metabolic tumors while comparatively sparing normal organs. TM was developed for the treatment of patients with Wilson’s disease,an autosomal recessive disorder of copper transport that results in excessive accumulation of copper and toxicity. Phase II and III trials with Wilson’s disease patients have demon-strated that TM is a fast-acting and well-tolerated drug(Brewer, 2009).The major side effect is anemia and leukopenia due to copper depletion,which can be reversed with a drug holiday or dose reduction.Additionally,the independent anti-angiogenic effects of the copper chelator TM present it as a drug that targets both tumor parenchyma and stroma,by enhancing cisplatin efﬁcacy against the cancer cells while inhibiting angiogenesis in the tumor bination regimens involving copper chelating and platinum-containing drugs may improve the treatment of cervical,ovarian,and other cancers for which cisplatin is currently in use,and for cancers that are treated with carboplatin, whose uptake is also mediated by CTR1(Holzer et al.,2006). Such a therapeutic strategy may even prove effective in treating cancers that are inherently resistant to cisplatin or have devel-oped resistance.With the development of a second-generation TM analog that depletes copper more quickly and is more stable(Lowndes et al.,2008),it may be possible to manipulate the activity of CTR1even more effectively.The mechanistic prin-ciples and results elaborated in this report should motivate discussion of analogous clinical trials in patients with cervical, ovarian,and other platinum-responsive cancers.Cancer CellModulating Copper Improves Cisplatin EfﬁcacyCancer Cell17,574–583,June15,2010ª2010Elsevier Inc.581。

高强箍筋混凝土柱的耗能性能

第４２卷第６期２０２０年１１月沈　阳　工　业　大　学　学　报ＪｏｕｒｎａｌｏｆＳｈｅｎｙａｎｇＵｎｉｖｅｒｓｉｔｙｏｆＴｅｃｈｎｏｌｏｇｙＶｏｌ４２Ｎｏ６Ｎｏｖ２０２０收稿日期：２０１９－０９－３０．基金项目：河南省科技厅项目（１７２１０２３１０７４９）；河南省教育厅项目（１６Ａ５６００３０）．作者简介：万海涛（１９７９－），男，江西抚州人，副教授，博士，主要从事结构抗震等方面的研究．本文已于２０２０－０９－２５０９∶４２在中国知网优先数字出版．网络出版地址：ｈｔｔｐ：∥ｋｎｓ．ｃｎｋｉ．ｎｅｔ／ｋｃｍｓ／ｄｅｔａｉｌ／２１．１１８９．Ｔ．２０２０１１０２．１３０３．０１６．ｈｔｍｌｄｏｉ：１０．７６８８／ｊ．ｉｓｓｎ．１０００－１６４６．２０２０．０６．１４高强箍筋混凝土柱的耗能性能万海涛，侯文彬，禹钿龙（河南大学土木建筑学院，河南开封４７５００４）摘　要：为了研究使用ＣＲＢ５５０级高强钢筋替代普通钢筋作为钢筋混凝土柱构件箍筋使用的可行性，对５根不同等级箍筋的ＲＣ柱构件进行低周往复加载试验，再使用有限元软件对其进行模拟，将试验结果与模拟结果进行对比，验证有限元模拟的可靠性．通过调整纵筋及箍筋直径、箍筋间距、柱截面尺寸等参数设计模拟了８根柱构件，从构件的延性和耗能角度研究不同等级箍筋柱构件的抗震性能．结果表明，使用ＣＲＢ５５０级钢筋替代ＨＲＢ４００级钢筋作为ＲＣ柱的箍筋后，ＲＣ柱的延性性能会适当提高，耗能性能变化不大，说明ＣＲＢ５５０级钢筋可以替代ＨＲＢ４００级钢筋作为ＲＣ柱构件的箍筋使用．关　键　词：高强钢筋；箍筋；钢筋混凝土；柱构件；低周往复加载试验；有限元模拟；延性；耗能中图分类号：ＴＵ３７５３文献标志码：Ａ文章编号：１０００－１６４６（２０２０）０６－０６９４－０７ＥｎｅｒｇｙｄｉｓｓｉｐａｔｉｏｎｐｅｒｆｏｒｍａｎｃｅｏｆＲＣｃｏｌｕｍｎＷＡＮＨａｉｔａｏ，ＨＯＵＷｅｎｂｉｎ，ＹＵＤｉａｎｌｏｎｇ（ＳｃｈｏｏｌｏｆＣｉｖｉｌＥｎｇｉｎｅｅｒｉｎｇａｎｄＡｒｃｈｉｔｅｃｔｕｒｅ，ＨｅｎａｎＵｎｉｖｅｒｓｉｔｙ，Ｋａｉｆｅｎｇ４７５００４，Ｃｈｉｎａ）Ａｂｓｔｒａｃｔ：ＩｎｏｒｄｅｒｔｏｓｔｕｄｙｔｈｅｆｅａｓｉｂｉｌｉｔｙｏｆｕｓｉｎｇｔｈｅＣＲＢ５５０ｈｉｇｈｓｔｒｅｎｇｔｈｒｅｉｎｆｏｒｃｅｍｅｎｔｉｎｓｔｅａｄｏｆｏｒｄｉｎａｒｙｒｅｉｎｆｏｒｃｅｍｅｎｔａｓｒｅｉｎｆｏｒｃｅｄｃｏｎｃｒｅｔｅ（ＲＣ）ｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓｔｉｒｒｕｐ，ｌｏｗｃｙｃｌｅｃｙｃｌｉｃｌｏａｄｉｎｇｔｅｓｔｓｗｅｒｅｐｅｒｆｏｒｍｅｄｕｓｉｎｇｆｉｖｅＲＣｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓｗｉｔｈｄｉｆｆｅｒｅｎｔｓｔｉｒｒｕｐｇｒａｄｅｓ，ａｎｄｔｈｅＲＣｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓｗｅｒｅｓｉｍｕｌａｔｅｄｗｉｔｈｆｉｎｉｔｅｅｌｅｍｅｎｔｓｏｆｔｗａｒｅ．Ｔｈｅｅｘｐｅｒｉｍｅｎｔａｌｍｅａｓｕｒｅｍｅｎｔｓｗｅｒｅｃｏｍｐａｒｅｄｗｉｔｈｔｈｅｓｉｍｕｌａｔｉｏｎｒｅｓｕｌｔｓｔｏｖｅｒｉｆｙｔｈｅｒｅｌｉａｂｉｌｉｔｙｏｆｆｉｎｉｔｅｅｌｅｍｅｎｔｓｉｍｕｌａｔｉｏｎ．Ｅｉｇｈｔｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓｗｅｒｅｓｉｍｕｌａｔｅｄｔｈｒｏｕｇｈａｄｊｕｓｔｉｎｇｓｕｃｈｐａｒａｍｅｔｅｒｓａｓｌｏｎｇｉｔｕｄｉｎａｌｒｅｉｎｆｏｒｃｅｍｅｎｔａｎｄｓｔｉｒｒｕｐｄｉａｍｅｔｅｒｓ，ｓｔｉｒｒｕｐｓｐａｃｉｎｇａｎｄｃｏｌｕｍｎｓｅｃｔｉｏｎｓｉｚｅ．Ｔｈｅｓｅｉｓｍｉｃｐｅｒｆｏｒｍａｎｃｅｏｆｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓｗｉｔｈｄｉｆｆｅｒｅｎｔｓｔｉｒｒｕｐｇｒａｄｅｓｗａｓｓｔｕｄｉｅｄｆｒｏｍｔｈｅｐｅｒｓｐｅｃｔｉｖｅｓｏｆｃｏｍｐｏｎｅｎｔｄｕｃｔｉｌｉｔｙａｎｄｅｎｅｒｇｙｄｉｓｓｉｐａｔｉｏｎ．ＴｈｅｒｅｓｕｌｔｓｓｈｏｗｔｈａｔｗｈｅｎｔｈｅＣＲＢ５５０ｇｒａｄｅｒｅｉｎｆｏｒｃｅｍｅｎｔｉｓｕｓｅｄａｓｔｈｅｓｔｉｒｒｕｐｏｆＲＣｃｏｌｕｍｎｉｎｓｔｅａｄｏｆｔｈｅＨＲＢ４００ｇｒａｄｅｒｅｉｎｆｏｒｃｅｍｅｎｔ，ｔｈｅｄｕｃｔｉｌｉｔｙｏｆＲＣｃｏｌｕｍｎｃａｎｂｅｐｒｏｐｅｒｌｙｉｍｐｒｏｖｅｄ，ａｎｄｔｈｅｅｎｅｒｇｙｄｉｓｓｉｐａｔｉｏｎｐｅｒｆｏｒｍａｎｃｅｗｉｌｌｃｈａｎｇｅｉｎｃｏｎｓｉｄｅｒａｂｌｙ，ｓｈｏｗｉｎｇｔｈａｔｔｈｅＣＲＢ５５０ｇｒａｄｅｒｅｉｎｆｏｒｃｅｍｅｎｔｃａｎｒｅｐｌａｃｅｔｈｅＨＲＢ４００ｇｒａｄｅｒｅｉｎｆｏｒｃｅｍｅｎｔａｓｔｈｅｓｔｉｒｒｕｐｏｆＲＣｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓ．Ｋｅｙｗｏｒｄｓ：ｈｉｇｈｓｔｒｅｎｇｔｈｒｅｉｎｆｏｒｃｅｍｅｎｔ；ｓｔｉｒｒｕｐ；ｒｅｉｎｆｏｒｃｅｄｃｏｎｃｒｅｔｅ；ｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔ；ｌｏｗｃｙｃｌｅｃｙｃｌｉｃｌｏａｄｉｎｇｔｅｓｔ；ｆｉｎｉｔｅｅｌｅｍｅｎｔｓｉｍｕｌａｔｉｏｎ；ｄｕｃｔｉｌｉｔｙ；ｅｎｅｒｇｙｄｉｓｓｉｐａｔｉｏｎ随着现代社会进步和经济发展，建筑结构正朝着超高层、超大层的方向发展，钢筋混凝土框架结构目前在高层建筑中应用较广，柱构件是钢筋混凝土框架结构的主要受力构件，柱构件的性能对结构整体的抗震性能有着重要影响．在国家当前可持续发展和推广绿色建筑的背景下，在结构沈　阳　工　业　大　学　学　报第４２卷中推广使用高强度、高性能钢筋有利于降低资源消耗，保护生态环境．目前，国家正在大力推广高强钢筋，但ＣＲＢ５５０级钢筋在抗震结构构件中还很少使用，因此，有必要对ＣＲＢ５５０级高强箍筋混凝土柱构件的抗震性能进行研究，为高强钢筋的应用推广工作提供参考．Ｗａｎｇ等［１］配制了５个配置ＨＲＢ６００级增强钢筋的活性粉末混凝土节点试件，结果表明，ＨＲＢ６００级钢筋的配置减轻了破坏，降低了强度退化和刚度退化，减少了残余变形，提高了变形能力和耗能能力；孙双喜［２］试验了８根纵筋和箍筋均为１４２０级ＰＣ钢棒的高强混凝土短柱，结果表明，配置高强钢棒后构件的承载力和变形能力均有所提高；刘伦等［３］通过对８根配置ＣＲＢ６００Ｈ级高强箍筋柱和１根配置ＨＲＢ４００级普通箍筋柱进行低周往复加载试验，分析了高强箍筋混凝土柱的抗震抗剪性能，结果表明，ＣＲＢ６００Ｈ级高强箍筋柱与ＨＲＢ４００级普通箍筋柱相比，抗剪承载力差异不大，但前者具有更好的延性和极限变形能力；刘佳妮等［４］分析了含箍特征值和箍筋形式对ＣＲＢ６００Ｈ级高强箍筋框架梁抗震抗剪性能的影响，结果表明，含箍特征值较高的构件，其承载力和变形能力较高，耗能能力较低，箍筋形式对抗剪承载力和变形能力影响不大，箍筋肢数多的构件耗能能力较高．对５根轴压比为０９的配置不同箍筋的混凝土柱构件进行低周往复加载试验，并使用有限元软件ＡＢＡＱＵＳ对柱构件进行数值模拟，将５根钢筋混凝土柱构件的模拟结果与试验的破坏形态、骨架曲线进行对比，验证了有限元模拟结果的可靠性．通过调整纵筋及箍筋直径、配箍间距、柱截面尺寸等参数设计并模拟了８根柱构件，从延性性能和耗能性能两个方面分析评价了ＲＣ柱构件的抗震性能，为ＣＲＢ５５０级高强钢筋在中国的推广及应用提供一定的参考．１　试验及有限元模拟概况１１　试件设计为了研究柱构件受力的最不利情况，柱构件的轴压比全部选用《建筑抗震设计规范》（ＧＢ５００１１２０１０）［５］中规定的０９０．柱构件截面尺寸分别为３５０ｍｍ×３５０ｍｍ与３００ｍｍ×３００ｍｍ两种，纵筋选用直径为１６、２２ｍｍ的ＨＲＢ４００级钢筋，箍筋选用直径为１０、１２ｍｍ的ＨＲＢ４００级钢筋与直径为７５、９５ｍｍ的ＣＲＢ５５０级钢筋，采用对称配筋，柱构件的详细参数如表１所示．混凝土强度等级为Ｃ３０，混凝土保护层厚度为４０ｍｍ，按照《混凝土结构试验方法标准》（ＧＢ／Ｔ５０１５２２０１２）［６］对混凝土进行取样并测试其强度，测得混凝土强度值如表２所示．按照《金属材料拉伸试验第１部分：室温试验方法》（ＧＢ／Ｔ２２８１２０１０）［７］对钢筋进行拉伸试验，测得钢筋强度参数如表３所示．柱构件通过基础梁固定，基础梁尺寸为１２００ｍｍ×６００ｍｍ×５００ｍｍ．表１　柱构件参数Ｔａｂ１　Ｐａｒａｍｅｔｅｒｓｏｆｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓ编号轴压比截面尺寸／ｍｍ纵筋纵筋配筋率／％箍筋箍筋配筋率／％Ａ１１０９３５０×３５０８１６１３１１０＠１００１７９Ａ１２０９３５０×３５０８１６１３１Ｒ７５＠１００１０７Ａ２２０９３５０×３５０８２２２４８Ｒ７５＠１００１０７Ｂ１１０９３００×３００８１６２０６１２＠１５０２０６Ｂ１２０９３００×３００８１６２０６Ｒ９５＠１５０１２２注：ＨＲＢ４００级钢筋；Ｒ为ＣＲＢ５５０级钢筋．表２　混凝土强度参数Ｔａｂ２　Ｐａｒａｍｅｔｅｒｓｏｆｃｏｎｃｒｅｔｅｓｔｒｅｎｇｔｈ（Ｎ·ｍｍ－２）立方体抗压强度标准值轴心抗压强度标准值轴心抗压强度设计值３１５０２１０７１５０５１２　有限元模型建立使用有限元软件ＡＢＡＱＵＳ对上述５根柱构件进行模拟，通过试验结果与有限元软件ＡＢＡＱＵＳ表３　钢筋强度参数　Ｔａｂ３　ＰａｒａｍｅｔｅｒｓｏｆｒｅｉｎｆｏｒｃｅｍｅｎｔｓｔｒｅｎｇｔｈＭＰａ钢筋类别弹性模量屈服强度极限强度１０２０１×１０５４８６６０５１２２０４×１０５４８９５９７１６２０２×１０５４７９６００２２２００×１０５４７７５９６Ｒ７５１９０×１０５５５９６９５Ｒ９５１９０×１０５５５６６９０５９６第６期万海涛，等：高强箍筋混凝土柱的耗能性能模拟结果之间的对比，来验证有限元模拟的准确度及可靠性，再通过调整纵筋及箍筋直径、配箍间距、柱截面尺寸等参数设计模拟了８根柱构件．所有模拟的柱构件参数如表４所示．表４　柱构件模型参数Ｔａｂ４　Ｐａｒａｍｅｔｅｒｓｏｆｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｍｏｄｅｌ编号轴压比截面尺寸／ｍｍ纵筋纵筋配筋率／％箍筋箍筋配筋率／％Ａ１１Ｍ０９３５０×３５０８１６１３１１０＠１００１７９Ａ１２Ｍ０９３５０×３５０８１６１３１Ｒ７５＠１００１０７Ａ１３Ｍ０９３５０×３５０８１６１３１Ｒ６５＠７０１００Ａ２１Ｍ０９３５０×３５０８２２２４８１０＠１００１７９Ａ２２Ｍ０９３５０×３５０８２２２４８Ｒ７５＠１００１０７Ａ２３Ｍ０９３５０×３５０８２２２４８Ｒ６５＠７０１００Ｂ１１Ｍ０９３００×３００８１６２０６１２＠１５０２０６Ｂ１２Ｍ０９３００×３００８１６２０６Ｒ９５＠１５０１２２Ｂ１３Ｍ０９３００×３００８１６２０６Ｒ７５＠１５００８６Ｂ２１Ｍ０９３００×３００８２２３３８Ｒ９５＠１５０１２２Ｂ２２Ｍ０９３００×３００８２２３３８Ｒ７５＠１５００８６Ｂ３１Ｍ０９３００×３００８１６２０６Ｒ９５＠１００１８０Ｂ３２Ｍ０９３００×３００８１６２０６Ｒ７５＠１００１２７钢筋混凝土柱构件建模时使用分离式模型，混凝土为三维实体单元Ｃ３Ｄ８Ｒ，钢筋为三维桁架单元Ｔ３Ｄ２，如图１所示．混凝土本构采用混凝土损伤塑性模型，泊松比为０２．由于本试验研究不同等级箍筋的性能差异，故ＣＲＢ５５０级高强钢筋采用双斜线本构模型，ＨＲＢ４００级钢筋采用理想弹塑性模型，泊松比均为０３．结合计算精度与运算时间要求，钢筋混凝土柱构件的网格尺寸设置为５０ｍｍ．级循环加载的最大水平承载力的连线，反映了构件在不同阶段所受荷载与其弹塑性形变的关系．通过有限元软件ＡＢＡＱＵＳ模拟获得的骨架曲线与试验获得的骨架曲线之间的对比分析，可以验证有限元模拟的可靠性，如图３所示．以构件Ｂ１２为代表对有限元模拟的骨架曲线与试验的骨架曲线进行对比．图３　Ｂ１２与Ｂ１２Ｍ骨架曲线对比Ｆｉｇ３　ＳｋｅｌｅｔｏｎｃｕｒｖｅｃｏｍｐａｒｉｓｏｎｏｆＢ１２ａｎｄＢ１２Ｍ由有限元软件ＡＢＡＱＵＳ模拟得到的骨架曲线与试验得到的骨架曲线吻合较好，模拟骨架曲线的屈服、峰值、极限状态的荷载、位移值与试验结果相差不大，说明有限元模拟能够较为准确地反映出实际试验得到的荷载位移之间的变化关系．３　柱的耗能性能分析３１　延性性能分析在荷载位移骨架曲线中，在最大承载力的前后段有明显的平台，此阶段构件有很大的变形但承载力没有显著降低，这种性质称为延性，延性大的构件破坏过程缓慢［８］．《建筑抗震试验规程》（ＪＧＪ／Ｔ１０１２０１５）［９］规定，结构的位移延性性能采用延性系数μ表示，其计算公式为μ＝ΔｕΔｙ（１）式中：Δｕ为试件的极限位移；Δｙ为试件的屈服位移．μ越大则构件能够产生的塑性变形就越大，地震对结构的影响就越小．为了保证高层建筑具有足够的抗震性能，一般要求结构的延性系数μ＞３．构件的极限位移取荷载位移骨架曲线下降段０８５倍的极限承载力处对应的位移，屈服位移通过Ｐａｒｋ法确定［１０－１２］．各构件在正、反两种加载方向下模拟骨架曲线上的特征数据和延性系数，如表５所示．表５　柱构件特征数据及延性系数Ｔａｂ５　Ｃｈａｒａｃｔｅｒｉｓｔｉｃｄａｔａａｎｄｄｕｃｔｉｌｉｔｙｃｏｅｆｆｉｃｉｅｎｔｏｆｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓ构件编号加载方向荷载／ｋＮ屈服峰值极限位移／ｍｍ屈服峰值极限延性系数Ａ１１ＭＡ１２ＭＡ１３ＭＡ２１ＭＡ２２ＭＡ２３ＭＢ１１ＭＢ１２ＭＢ１３ＭＢ２１ＭＢ２２ＭＢ３１ＭＢ３２Ｍ正向　２３７　２７１　２３０　１４４　２６０　５６２负向－２４０－３０６－２６０－１３８－２８７－５７５正向２３６２７５２３４１０６２２９５４０负向－２４９－２８４－２４１－１２７－２６１－６０８正向２４１２８５２４２１２３２９０５６４负向－２３２－２９３－２４９－１３１－３１２－５７６正向３００３１７２６９１６２３０７５６２负向－３１８－３２４－２７５－１３７－２６７－５５０正向２８６３００２５５１４６２２５５２７负向－２７９－２９８－２５３－１５２－１８２－５３２正向３２４３４３２９２１５９２８５５８３负向－３１３－３３８－２８７－１４８－２７３－５７７正向１７９１９６１６７１０６２４２５２０负向－１８５－２０４－１７３－１４７－２３０－４７２正向１８２１９７１６７１３７２４３５３８负向－１８４－２１０－１７９－１２３－２６０－５１９正向１７５１８８１６０１００１９０３９７负向－１７８－１８７－１５９－９２－１２１－４３９正向２３７２６５２２５１３３２２９４７６负向－２３１－２６１－２２２－１２９－２３６－４４１正向２３５２５７２１８１２６２２８４５３负向－２３１－２４８－２１１－１２８－２２４－４９７正向２０１２２１１８８１０１２４９４９４负向－２１４－２２８－１９４－１０８－２１３－４５５正向２０５２２９１９５１２２２０３５１４负向－１９７－２３０－１９６－１１８－１９３－５０４４０４９４５３７３６３８４０４１４４３５３７４６４２７９６第６期万海涛，等：高强箍筋混凝土柱的耗能性能由表５可以看出：１）所有柱构件的延性系数均大于３，符合抗震规范要求，表明高强箍筋混凝土柱构件和普通箍筋混凝土柱构件都具有良好的延性性能．２）按照等强度原则，构件Ａ１１Ｍ布置了直径为１０ｍｍ的ＨＲＢ４００级箍筋，构件Ａ１２Ｍ布置了直径为７５ｍｍ的ＣＲＢ５５０级高强箍筋，后者的箍筋配筋率比前者降低了０７２％，但荷载值和位移值相近，高强箍筋混凝土柱构件的延性系数比普通箍筋混凝土柱构件的提高了０９；构件Ｂ１１Ｍ与构件Ｂ１２Ｍ的荷载值和位移值相近，后者延性系数稍大，表明使用较小直径的ＣＲＢ５５０级高强箍筋替代较大直径的ＨＲＢ４００级普通箍筋后，钢筋混凝土柱构件的延性性能会有一定程度的提高．３）对于箍筋种类和直径相同的柱构件，如构件Ａ１１Ｍ和构件Ａ２１Ｍ，前者的延性系数为４０，后者为３７，说明提高柱构件的纵筋配筋率后延性性能会降低，原因是构件的纵筋配筋率提高后刚度增强，延缓了构件屈服，而极限状态基本没有发生变化，最终导致延性系数降低，其余构件如Ａ２２Ｍ的延性系数比Ａ１２Ｍ降低了２６５％，构件Ａ２３Ｍ的延性系数比Ａ１３Ｍ降低了１５６％，构件Ｂ２１Ｍ的延性系数比Ｂ１２Ｍ降低了１２５％，构件Ｂ２２Ｍ的延性系数比Ｂ１３Ｍ降低了１５９％．４）Ｂ组纵向钢筋直径为１６ｍｍ的ＨＲＢ４００级钢筋柱构件中，配置同一种箍筋的钢筋混凝土柱构件延性系数会随着箍筋间距的减小和强度的增强而提高，但效果不明显．３２　耗能能力分析结构的耗能能力是评价其抗震性能的重要指标，钢筋混凝土结构作为弹塑性结构，在地震作用时能够吸收地震传递的能量．结构在反复荷载作用下，加载时吸收的能量与卸载时释放能量的差值就是构件在一个循环加载过程中能量耗散能力，在滞回曲线中表现为滞回环包围的面积，如图４所示．吸收或者耗散的能量越大，结构的抗震性能越好．图４　滞回环面积示意图Ｆｉｇ４　Ｓｃｈｅｍａｔｉｃｄｉａｇｒａｍｏｆｈｙｓｔｅｒｅｔｉｃｌｏｏｐａｒｅａ评价结构的能量耗散能力主要有能量耗能系数Ｅ和等效黏滞阻尼系数ξｅｑ两种指标，计算公式为Ｅ＝ＳＡＢＣ＋ＳＡＣＤＳΔＯＢＥ＋ＳΔＯＤＦ　（２）ξｅｑ＝ＳＡＢＣ＋ＳＡＣＤ２π（ＳΔＯＢＥ＋ＳΔＯＤＦ）　（３）使用数据处理软件Ｏｒｉｇｉｎ对滞回曲线上各滞回环进行数值积分运算，得到构件滞回环包围的面积，总滞回耗能是各滞回环包围面积的总和，如表６所示．由于各构件的屈服、峰值、极限三个状态点的荷载值与位移值各不相同，加载级数不同，为了能够采用同一标准对各柱构件的滞回耗能性能进行对比分析，将总滞回耗能除以构件的屈服荷载和屈服位移之积，进行归一化无量纲处理．表６　柱构件的滞回耗能Ｔａｂ６　Ｈｙｓｔｅｒｅｔｉｃｅｎｅｒｇｙｄｉｓｓｉｐａｔｉｏｎｏｆｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓ加载方式Ａ１１ＭＡ１２ＭＡ１３ＭＡ２１ＭＡ２２ＭＡ２３ＭＢ１１ＭＢ１２ＭＢ１３ＭＢ２１ＭＢ２２ＭＢ３１ＭＢ３２Ｍ第一级加载　１０６２０　１２３８０　１１４７０　１７８００　１５４９０　１５８１０　６８５０　６６９０　５２６０　９２１０１０５７０　６７２０　７３７０第二级加载８７１７０１２５２００１０９３１０１０４８００１０５７４０９７７４０２８６６０３０５７０２９４２０５３２９０５０６７０３２６９０３６０１０第三级加载１３７８３０１９６１３０１８８７６０１８０８７０１７６５３０１７９６３０６９３９０８２８００６７８４０９４８１０９２６４０８８２３０８５６１０第四级加载１８９７７０２５３６３０２６０３１０２６９５２０２６５５００２５６１７０１２０９３０１２７８９０９８３８０１５４８７０１４４１６０１４０８５０１５４９５０第五级加载２４４８００３４５１７０３３６２４０３５９１１０３４２２６０３３８６１０１９３７４０１８２３００１３９３１０２３０８９０２２８６８０２０７０８０１９７３２０第六级加载３０１９５０３９１７３０３７９６１０４２７３８０４０２６６０３９６８４０２５１４２０２３７４４０１８１８３０２７４３８０２６２６１０２５５０８０２６４９７０第七级加载３４７６８０４２５３１０３９４２７０４７７２００４３０９２０４２１４８０－－－－－－－第八级加载３７２３１０－－－－－－－－－－－－总计１６９２１３０１７４９５５０１６７９９７０１８３６６８０１７３９１００１７０６２８０６７０９９０６６７６９０５２２０４０８１７４５０７８９３３０７３０６５０７４６２３０ＰｙΔｙ３３６２４２８３２０３００１８４６０８３４２０８２４８９２１２３１７７２４４２０１６９３８３０６６０２９５８９２１７０７２４１２８滞回耗能归一化５０３６１８５６０３９９４１３３４９２９０２７３３０８２６７２６７３３７３０９注：Ｐｙ为试件的屈服荷载．８９６沈　阳　工　业　大　学　学　报第４２卷从表６中可以看出：１）随着加载级数的增加，每级滞回耗能增长率在临近破坏时有下降趋势．截面尺寸较大的Ａ组构件滞回耗能明显高于截面尺寸小的Ｂ组构件．２）对比三组柱构件Ａ１１Ｍ和Ａ１２Ｍ、Ａ２１Ｍ和Ａ２２Ｍ、Ｂ１１Ｍ和Ｂ１２Ｍ发现，按照等强度原则分别配置了ＣＲＢ５５０级箍筋和ＨＲＢ４００级箍筋的混凝土柱构件归一化总滞回耗能值相近，都具有良好的能量耗散能力．在低周往复加载试验过程中，钢筋混凝土柱构件会出现强度和刚度上的退化，这种现象会对柱构件滞回环面积的大小产生影响．为了表达构件的这个特征，在现代工程抗震中采用等效黏滞阻尼系数ξｅｑ来评价构件在地震作用中耗能性能的好坏．构件的等效黏滞阻尼系数ξｅｑ越大，其耗能性能就越好．所有模拟柱构件各级加载循环的等效黏滞阻尼系数变化情况如图５所示．图５　柱构件的等效黏滞阻尼系数Ｆｉｇ５　Ｅｑｕｉｖａｌｅｎｔｖｉｓｃｏｕｓｄａｍｐｉｎｇｃｏｅｆｆｉｃｉｅｎｔｏｆｃｏｌｕｍｎｃｏｍｐｏｎｅｎｔｓ对比Ａ、Ｂ两组不同截面的柱构件，等效黏滞阻尼系数随着截面尺寸的降低而减小；在纵向配筋率相同的Ｂ组构件中，采用同一种类箍筋混凝土柱构件的等效黏滞阻尼系数随着箍筋间距的减小而增大．较大箍筋直径的ＨＲＢ４００级箍筋柱构件Ａ２１Ｍ和较小箍筋直径的ＣＲＢ５５０级箍筋柱构件Ａ２２Ｍ的等效黏滞阻尼系数曲线贴近，说明两者的耗能能力基本相同；较大箍筋直径的ＨＲＢ４００级箍筋柱构件Ａ１１Ｍ和较小箍筋直径的ＣＲＢ５５０级箍筋柱构件Ａ１２Ｍ的等效黏滞阻尼系数曲线出现了上下交错的情况，说明两者的耗能能力不相上下．总体上说，按照等强度原则配置了ＣＲＢ５５０级高强箍筋的柱构件和配置了ＨＲＢ４００级箍筋的柱构件的等效黏滞阻尼系数相近，耗能能力相当．４　结　论本文通过分析得出以下结论：１）通过有限元模拟得到的应变云图与构件的试验破坏形态图基本相似，且二者的骨架曲线拟合较好，说明有限元模拟可以较为准确地对试验进行模拟．２）配置较小直径的ＣＲＢ５５０级高强箍筋混凝土柱构件的骨架曲线与较大直径的ＨＲＢ４００级箍筋混凝土柱构件的骨架曲线吻合较好，承载力相差在１５％以内，说明ＣＲＢ５５０级高强钢筋可以替代ＨＲＢ４００级普通钢筋作为钢筋混凝土柱构件的箍筋．３）在延性性能方面，所有柱构件的延性系数均大于３０，具有良好的变形能力；在ＣＲＢ５５０级高强箍筋等强度替代ＨＲＢ４００级箍筋后，钢筋混凝土柱构件的延性性能会适当提高．４）在能量耗散方面，配置了较小直径的ＣＲＢ５５０级高强箍筋混凝土柱构件与对应的ＨＲＢ４００级普通箍筋混凝土柱构件的归一化总滞回耗能值、等效黏滞阻尼系数接近，都具有很好的耗能能力．参考文献（Ｒｅｆｅｒｅｎｃｅｓ）：［１］ＷａｎｇＤＨ，ＪｕＹＺ，ＺｈｅｎｇＷＺ．Ｓｔｒｅｎｇｔｈｏｆｒｅａｃｔｉｖｅｐｏｗｄｅｒｃｏｎｃｒｅｔｅｂｅａｍｃｏｌｕｍｎｊｏｉｎｔｓｒｅｉｎｆｏｒｃｅｄｗｉｔｈｈｉｇｈｓｔｒｅｎｇｔｈ（ＨＲＢ６００）ｂａｒｓｕｎｄｅｒｓｅｉｓｍｉｃｌｏａｄｉｎｇ［Ｊ］．ＳｔｒｅｎｇｔｈｏｆＭａｔｅｒｉａｌｓ，２０１７，４９（１）：１３９－１５１．［２］孙双喜．高强钢筋高强混凝土短柱抗震性能研究［Ｄ］．沈阳：沈阳建筑大学，２０１５．（ＳＵＮＳｈｕａｎｇｘｉ．Ｒｅｓｅａｒｃｈｏｎｓｅｉｓｍｉｃｐｅｒｆｏｒｍａｎｃｅｏｆｈｉｇｈｓｔｒｅｎｇｔｈｃｏｎｃｒｅｔｅｓｈｏｒｔｃｏｌｕｍｎｒｅｉｎｆｏｒｃｅｄｗｉｔｈｈｉｇｈｓｔｒｅｎｇｔｈｒｅｂａｒ［Ｄ］．Ｓｈｅｎｙａｎｇ：ＳｈｅｎｙａｎｇＪｉａｎｚｈｕＵｎｉｖｅｒｓｉｔｙ，２０１５．）［３］刘伦，朱爱萍，翟文，等．配置ＣＲＢ６００Ｈ高强箍筋混凝土柱抗震抗剪性能试验研究［Ｊ］．结构工程师，２０１９，３５（３）：１９２－２０１．（ＬＩＵＬｕｎ，ＺＨＵＡｉｐｉｎｇ，ＺＨＡＩＷｅｎ，ｅｔａｌ．ＥｘｐｅｒｉｍｅｎｔａｌｒｅｓｅａｒｃｈｏｎｓｅｉｓｍｉｃｓｈｅａｒｂｅｈａｖｉｏｒｏｆＣＲＢ６００Ｈｈｉｇｈｓｔｒｅｎｇｔｈｓｔｉｒｒｕｐｓｃｏｎｃｒｅｔｅｃｏｌｕｍｎｓ［Ｊ］．ＳｔｒｕｃｔｕｒａｌＥｎｇｉｎｅｅｒｓ，２０１９，３５（３）：１９２－２０１．）［４］刘佳妮，朱爱萍，翟文，等．配置ＣＲＢ６００Ｈ箍筋的混凝土框架梁抗震抗剪试验研究［Ｊ］．工业建筑，９９６第６期万海涛，等：高强箍筋混凝土柱的耗能性能２０１８，４８（１１）：６５－７２．（ＬＩＵＪｉａｎｉ，ＺＨＵＡｉｐｉｎｇ，ＺＨＡＩＷｅｎ，ｅｔａｌ．ＥｘｐｅｒｉｍｅｎｔａｌｒｅｓｅａｒｃｈｏｎｓｅｉｓｍｉｃｂｅｈａｖｉｏｒａｎｄｓｈｅａｒｂｅｈａｖｉｏｒｏｆｃｏｎｃｒｅｔｅｆｒａｍｅｂｅａｍｗｉｔｈＣＲＢ６００Ｈｓｔｉｒｒｕｐｓ［Ｊ］．ＩｎｄｕｓｔｒｉａｌＣｏｎｓｔｒｕｃｔｉｏｎ，２０１８，４８（１１）：６５－７２．）［５］中华人民共和国住房和城乡建设部．ＧＢ５００１１２０１０建筑抗震设计规范［Ｓ］．北京：中国建筑工业出版社，２０１６．（ＭｉｎｉｓｔｒｙｏｆＨｏｕｓｉｎｇａｎｄＵｒｂａｎＲｕｒａｌＤｅｖｅｌｏｐｍｅｎｔｏｆｔｈｅＰｅｏｐｌｅ’ｓＲｅｐｕｂｌｉｃｏｆＣｈｉｎａ．ＧＢ５００１１２０１０Ｃｏｄｅｆｏｒｓｅｉｓｍｉｃｄｅｓｉｇｎｏｆｂｕｉｌｄｉｎｇｓ［Ｓ］．Ｂｅｉｊｉｎｇ：ＣｈｉｎａＡｒｃｈｉｔｅｃｔｕｒｅ＆ＢｕｉｌｄｉｎｇＰｒｅｓｓ，２０１６．）［６］中华人民共和国住房和城乡建设部．ＧＢ／Ｔ５０１５２２０１２混凝土结构试验方法标准［Ｓ］．北京：中国建筑工业出版社，２０１２．（ＭｉｎｉｓｔｒｙｏｆＨｏｕｓｉｎｇａｎｄＵｒｂａｎＲｕｒａｌＤｅｖｅｌｏｐｍｅｎｔｏｆｔｈｅＰｅｏｐｌｅ’ｓＲｅｐｕｂｌｉｃｏｆＣｈｉｎａ．ＧＢ／Ｔ５０１５２２０１２Ｓｔａｎｄａｒｄｆｏｒｔｅｓｔｍｅｔｈｏｄｏｆｃｏｎｃｒｅｔｅｓｔｒｕｃｔｕｒｅｓ［Ｓ］．Ｂｅｉｊｉｎｇ：ＣｈｉｎａＡｒｃｈｉｔｅｃｔｕｒｅ＆ＢｕｉｌｄｉｎｇＰｒｅｓｓ，２０１２．）［７］中华人民共和国国家质量监督检验检疫总局．ＧＢ／Ｔ２２８１２０１０金属材料拉伸试验第１部分：室温试验方法［Ｓ］．北京：中国标准出版社，２０１０．（ＧｅｎｅｒａｌＡｄｍｉｎｉｓｔｒａｔｉｏｎｏｆＱｕａｌｉｔｙＳｕｐｅｒｖｉｓｉｏｎ，ＩｎｓｐｅｃｔｉｏｎａｎｄＱｕａｒａｎｔｉｎｅｏｆｔｈｅＰｅｏｐｌｅ’ｓＲｅｐｕｂｌｉｃｏｆＣｈｉｎａ．ＧＢ／Ｔ２２８１２０１０Ｍｅｔａｌｌｉｃｍａｔｅｒｉａｌｓｔｅｎｓｉｌｅｔｅｓｔｉｎｇｐａｒｔ１：ｍｅｔｈｏｄｏｆｔｅｓｔａｔｒｏｏｍｔｅｍｐｅｒａｔｕｒｅ［Ｓ］．Ｂｅｉｊｉｎｇ：ＳｔａｎｄａｒｄｓＰｒｅｓｓｏｆＣｈｉｎａ，２０１０．）［８］沈聚敏，周锡元，高小旺，等．抗震工程学［Ｍ］．北京：中国建筑工业出版社，２０００．（ＳＨＥＮＪｕｍｉｎ，ＺＨＯＵＸｉｙｕａｎ，ＧＡＯＸｉａｏｗａｎｇ，ｅｔａｌ．Ａｓｅｉｓｍｉｃｅｎｇｉｎｅｅｒｉｎｇ［Ｍ］．Ｂｅｉｊｉｎｇ：ＣｈｉｎａＡｒｃｈｉｔｅｃｔｕｒｅ＆ＢｕｉｌｄｉｎｇＰｒｅｓｓ，２０００．）［９］中华人民共和国住房和城乡建设部．ＪＧＪ／Ｔ１０１２０１５建筑抗震试验规程［Ｓ］．北京：中国建筑工业出版社，２０１５．（ＭｉｎｉｓｔｒｙｏｆＨｏｕｓｉｎｇａｎｄＵｒｂａｎＲｕｒａｌＤｅｖｅｌｏｐｍｅｎｔｏｆｔｈｅＰｅｏｐｌｅ’ｓＲｅｐｕｂｌｉｃｏｆＣｈｉｎａ．ＪＧＪ／Ｔ１０１２０１５Ｓｐｅｃｉｆｉｃａｔｉｏｎｆｏｒｓｅｉｓｍｉｃｔｅｓｔｏｆｂｕｉｌｄｉｎｇｓ［Ｓ］．Ｂｅｉｊｉｎｇ：ＣｈｉｎａＡｒｃｈｉｔｅｃｔｕｒｅ＆ＢｕｉｌｄｉｎｇＰｒｅｓｓ，２０１５．）［１０］万海涛．钢筋混凝土梁、柱构件抗震性能试验及其基于变形性能的参数研究［Ｄ］．广州：华南理工大学，２０１０．（ＷＡＮＨａｉｔａｏ．ＳｅｉｓｍｉｃｐｅｒｆｏｒｍａｎｃｅｔｅｓｔｏｆＲＣｂｅａｍｓａｎｄｃｏｌｕｍｎｓａｎｄｒｅｓｅａｒｃｈｏｎｐａｒａｍｅｔｅｒｏｆｄｅｆｏｒｍａｔｉｏｎｐｅｒｆｏｒｍａｎｃｅ［Ｄ］．Ｇｕａｎｇｚｈｏｕ：ＳｏｕｔｈＣｈｉｎａＵｎｉｖｅｒｓｉｔｙｏｆＴｅｃｈｎｏｌｏｇｙ，２０１０．）［１１］禹钿龙．ＣＲＢ５５０级高强箍筋混凝土柱的滞回耗能性能研究［Ｄ］．开封：河南大学，２０１９．（ＹＵＤｉａｎｌｏｎｇ．ＳｔｕｄｙｏｎｈｙｓｔｅｒｅｔｉｃｅｎｅｒｇｙｄｉｓｓｉｐａｔｉｏｎｐｅｒｆｏｒｍａｎｃｅｏｆＲＣｃｏｌｕｍｎｗｉｔｈＣＲＢ５５０ｓｔｉｒｒｕｐｓ［Ｄ］．Ｋａｉｆｅｎｇ：ＨｅｎａｎＵｎｉｖｅｒｓｉｔｙ，２０１９．）［１２］孙艳丽，董文天，刘娟，等．钢骨钢管混凝土柱与钢梁空间节点抗震性能的非线性有限元分析［Ｊ］．沈阳建筑大学学报（自然科学版），２０１７，３３（５）：８８１－８８９．（ＳＵＮＹａｎｌｉ，ＤＯＮＧＷｅｎｔｉａｎ，ＬＩＵＪｕａｎ，ｅｔａｌ．ＮｏｎｌｉｎｅａｒＦＥＭａｎａｌｙｓｉｓｏｎｓｅｉｓｍｉｃｐｅｖｆｏｒｍａｎｃｅｏｆｓｐａｃｅｊｏｉｎｔｏｆｓｔｅｅｌｔｕｂｕｌａｒｃｏｌｕｍｎｓｆｉｌｌｅｄｗｉｔｈｓｔｅｅｌｒｅｉｎｆｏｒｃｅｄｃｏｎｃｒｅｔｅａｎｄｓｔｅｅｌｂｅａｍ［Ｊ］．ＪｏｕｒｎａｌｏｆＳｈｅｎｙａｎｇＪｉａｎｚｈｕＵｎｉｖｅｒｓｉｔｙ（ＮａｔｕｒａｌＳｃｉｅｎｃｅ），２０１７，３３（５）：８８１－８８９．）（责任编辑：钟　媛　英文审校：尹淑英）００７沈　阳　工　业　大　学　学　报第４２卷。

钙离子和氯化钴促进成纤维细胞色素上皮衍生因子的表达

相对表达量，别为２３和１４倍，者差异具有显著性（００）分．倍．二Ｐ＜．５。结论
氯化钴的调节。
真皮成纤维细胞表达ＰＤ而且受钙离子和ＥＦ，
（键词］成纤维细胞；色素上皮衍生因子；钙离子；氯化钴关
［ｂｔａｔ０ｂｅｔｅＴｏｉｖｓｉａｅｈｘｒｓｉｎｏＥＦａｄｉｅｕａｉｎｂａｃｕａｄｃｂｌＡｓｒｃ］ｊｃｉｅｔｔｔｅｅｐｅｓｆＤｎｓｇｌｔｙｃｌｉｍｎｏａｔｖｎｇｏＰｔｒｏ
ｃｌｒｄ（Ｃｌ）ｉｅｍａｉｒｂａｔ．ＭｅｈｄｎｅｎｎｒｌｈｍａｏｅｋｎｓｅｉｅｓｗｅｅｇｏｐｄｈｏｉｅＣｏ，ｎｄｒｌｂｏｌｓｓｆｔｏｓＮｉｔｅｏｍａｕｎｆｒｓｉｐｃｍｎｒｒｕｅｉｔｒＩｃｏｄｎｏｔｅｒａｅｅｏ（＝１）ｏｂｖ０（ｎｏＩｏＩａｃｒｉｇｔｈｉｇｓｂｌｗｎ３ｒａｏｅ４ｎ一６．Ｄｅｍａｉｒｂａｔｒｅａａｅ）ｒｌｆｏｌｓｓｗｅｅｓｐｒｔｄｂ
ｂｙｃｌａｎａｅｔｐｅＩａｄｃｌｕｒｄａ。ｉｏｌｇｅｓｙｎｕｔｅｔ３７Ｃｎ５ＣＯ．Ｃａｃｕｍｔ１０ｍＭｎｄＣｏ２ａＯ２ｌｉａ．ａＣ１ｔＩ０￣Ｍｒｎｃ — ｗｅｅｉｕｂａｅｔｉｏａｔｏ４ｔｄｗｉｈｆｂｒｂｌｓｓｆｒ２ｈ．ＴｈｅｍＲＮＡｎｒｅｎｌｖｅｆＰＥＤＦｗａｔｒｉｄｂｙＲＴ— ａｄｐｏｔｉｅ１ｏｓｄｅｅｍｎｅＰＣＲｎｍ— ａｄｉ

哈氏合金碳化物析出英语

哈氏合金碳化物析出英语Hastelloy Alloy and Carbide Precipitation.Hastelloy, a renowned brand of nickel-based superalloys, is widely used in harsh environments due to its exceptional corrosion resistance and high-temperature stability. However, one of the key considerations in the applicationof Hastelloy alloys is the precipitation of carbides, which can significantly affect the mechanical properties and corrosion resistance of the material.Understanding Carbide Precipitation.Carbide precipitation refers to the formation ofcarbide particles within the alloy matrix during specific heat treatments or exposure to elevated temperatures. These carbide particles can vary in size, shape, and distribution, depending on the alloy composition, heat treatment parameters, and exposure conditions.In Hastelloy alloys, carbide precipitation can occur due to the interaction of carbon with other alloying elements, such as chromium, molybdenum, and tungsten. These elements have a strong affinity for carbon, leading to the formation of stable carbide compounds.Effects of Carbide Precipitation.The precipitation of carbides in Hastelloy alloys can have both beneficial and detrimental effects. On the one hand, the formation of fine, uniformly distributed carbide particles can enhance the strength and hardness of the alloy, improving its wear resistance and fatigue life.However, the presence of larger or irregularly shaped carbide particles can have negative impacts. Theseparticles can act as sites for corrosion initiation, leading to localized corrosion attacks, such as pitting or crevice corrosion. Additionally, carbide precipitation can affect the alloy's toughness and ductility, making it more susceptible to cracking or fracture under certain conditions.Controlling Carbide Precipitation.To minimize the negative effects of carbide precipitation, several strategies can be employed. Firstly, alloy composition can be optimized to reduce the amount of carbon and carbide-forming elements, thus limiting the potential for carbide formation. Secondly, heat treatment protocols can be carefully designed to control the kinetics of carbide precipitation, ensuring that the particles are fine and uniformly distributed.Moreover, the use of stabilizers, such as titanium or aluminum, can help to maintain the stability of the carbide particles, preventing their growth and coarsening during exposure to elevated temperatures.Applications Considerations.When selecting Hastelloy alloys for specific applications, it is crucial to consider the potential for carbide precipitation. Applications that involve exposureto high temperatures or corrosive environments may require special attention to ensure that carbide precipitation does not compromise the alloy's performance.For instance, in chemical processing equipment or oil and gas exploration equipment, where Hastelloy alloys are commonly used, regular inspections and maintenance are essential to monitor for carbide precipitation and take corrective measures if necessary.In conclusion, carbide precipitation is a crucial aspect to consider when working with Hastelloy alloys. By understanding its mechanisms, effects, and controlling strategies, engineers and material scientists can ensure that these alloys perform reliably in even the most demanding applications.。

β环糊精包埋碳量子点

β环糊精包埋碳量子点英文回答：β-cyclodextrin-encapsulated carbon quantum dots are a fascinating topic in the field of nanomaterials. These carbon quantum dots, also known as CQDs, are tiny nanoparticles with unique optical and electronic properties. They have a wide range of potential applications in fields such as bioimaging, sensing, and drug delivery.The encapsulation of carbon quantum dots within β-cyclodextrin, a cyclic oligosaccharide, offers several advantages. Firstly, β-cyclodextrin provides a protective shell around the carbon quantum dots, preventing their degradation and enhancing their stability. This is particularly important for applications that require long-term stability, such as in vivo imaging or drug delivery systems.Furthermore, the β-cyclodextrin shell can also serveas a carrier for hydrophobic molecules. This means that the carbon quantum dots can be loaded with drugs or other therapeutic agents, enhancing their potential as drug delivery vehicles. The hydrophobic molecules can be released from the β-cyclodextrin shell in response to specific stimuli, such as changes in pH or temperature, allowing for controlled drug release.In addition to their protective and drug delivery capabilities, β-cyclodextrin-encapsulated carbon quantum dots also exhibit enhanced optical properties. The encapsulation process can lead to improved photoluminescence efficiency and stability of the carbon quantum dots, making them ideal for applications in bioimaging and sensing. For example, the enhanced photoluminescence properties can be utilized for fluorescent labeling of biological samples, enabling the visualization and tracking of specific cells or molecules in complex biological systems.Overall, the encapsulation of carbon quantum dotswithin β-cyclodextrin offers a versatile platform for thedevelopment of advanced nanomaterials with enhanced stability, drug delivery capabilities, and optical properties. The combination of these features makes them highly promising for a wide range of applications in fields such as medicine, environmental monitoring, and electronics.中文回答：β-环糊精包埋碳量子点是纳米材料领域的一个引人注目的研究课题。

4G蛋白偶联信号通路

第四十六页，共四十八页。
C
蛋白激酶的作用
第四十七页，共四十八页。
内容总结
G蛋白偶联信号通路。不同组织的不同细胞虽有很大差异，但总体可以分为三类：。 cAMP信号的终止： cAMP磷酸二酯酶(PDE)。膜受体与信号分子结合后，激活膜上的Gq蛋白(一种G蛋白)。信号的终止是通过去磷酸化形成自由的肌醇。对代谢的调节作用：PKC被激活后可引起一系列靶蛋白的丝氨酸和〔或〕苏氨酸残基发生磷酸化反响。PKC对基因的活化过程可分为早期反响和晚期反响两个阶段。PKC 对基因的活化分为早期反响和晚期反响
内质网
磷脂酶在细胞内信使DG、 IP3和 Ca2+的形成中起着关键作用。
钙信号系统
磷脂酰肌醇信号通路的最大特点是胞外信号被膜受体接受后，同时产生两个胞内信使，分别启动两个信号传递途径即IP3—Ca 2 +和DG—PKC途径，实现细胞对外界的应答，因此把这一信号系统称
之为“双信使系统〞。
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细胞核
PKC 对基因的早期活化和晚期活化
第四十五页，共四十八页。
目录
在许多细胞中，PKC的活化可增强特殊基因转录。有两条途径： ①PKC激活一条蛋白激酶的级联反响，导致基因调控蛋白的磷酸化和激活； ②PKC的活化，导致一种抑制蛋白的磷酸化，使基因调控蛋白摆脱抑制状态释放出来，进入细胞核，刺激特殊基因的转录。
酶
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利钠肽
第八页，共四十八页。
腺苷酸环化酶(adenylate
cyclase， AC) 是膜整合蛋白，能够将ATP转变成cAMP，引起细胞的信号应答。
第九页，共四十八页。

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StamberShow preview | Related articles | Related reference work articlesPurchase$ 31.50340The long term financial impacts of CVD: Livingstandards in retirementOriginal Research ArticleInternational Journal of Cardiology , In Press, CorrectedProof , Available online 20 November 2010Deborah Schofield, Simon Kelly, Rupendra Shrestha,Megan Passey, Emily Callander, Richard Percival Show preview | Related articles | Related reference work articlesPurchase $ 31.50341 An integrated safety prognosis model for complexsystem based on dynamic Bayesian network and antcolony algorithm Original Research ArticleExpert Systems with Applications, Volume 38, Issue 3,March 2011, Pages 1431-1446Jinqiu Hu, Laibin Zhang, Lin Ma, Wei LiangShow preview | Related articles | Related reference work articlesPurchase$ 41.95342 Deploying effective service strategy in the operationsstage of high-speed rail Original Research ArticleTransportation Research Part E: Logistics andTransportation Review, Volume 47, Issue 4, July 2011,Pages 507-519Jui-Sheng Chou, Changwan Kim, Yao-Chen Kuo, Nai-ChiOuShow preview | Related articles | Related reference work articlesPurchase$ 41.95Research highlights► Appraisal of high-speed rail service quality and performance. ► Evaluation of overall and social demographic attribute-level satisfaction. ► Passenger satisfaction index (PSI) model for high-speed rail corporations. ► PSI with the modified importance-performance analysis to implement an effective service strategy.343Methodology for modelling plug-in electric vehicles inthe power system and cost estimates for a systemwith either smart or dumb electric vehicles OriginalResearch ArticleEnergy, Volume 36, Issue 3, March 2011, Pages1758-1767Juha Kiviluoma, Peter MeibomShow preview | Related articles | Related reference work articlesPurchase$ 35.95Research highlights► The costs of …smart‟ and …dumb‟ electric vehicles were estimated. ► The power system benefits of smart electric vehicles were 227 €/vehicle/year. ► Two models were used: a generation planning model and a unit commitment model. ► Impact of electric vehicles in the long-term power plant portfolio are important. ► A model for electric vehicles was made for the stochastic unit commitment model.344Simulation of a supercapacitor/Li-ion battery hybrid for pulsed applications Original Research ArticleJournal of Power Sources , Volume 195, Issue 9, 1 May 2010, Pages 2731-2736D. Cericola, P.W. Ruch, R. Kötz, P. Novák, A. WokaunShow preview | Related articles | Related reference work articlesPurchase$ 41.95345Inter-regional comparisons of production technology in Canada's timber harvesting industries OriginalResearch ArticleForest Policy and Economics , Volume 8, Issue 8, 1November 2006, Pages 797-810David M. Nanang, Asghedom Ghebremichael Show preview | Related articles | Related reference work articlesPurchase $ 19.95346Exploring passenger preferences in airline service attributes: A noteJournal of Air Transport Management , Volume 15, Issue 1,January 2009, Pages 52-53Ching-Fu Chen, Tsai-Fang WuShow preview | Related articles | Related reference work articles Purchase $ 35.95347Networking beyond the software code? an explorative examination of the development of an open source carproject Original Research ArticleTechnovation , Volume 30, Issues 11-12,November-December 2010, Pages 627-634Gordon Müller-Seitz, Guido Reger Show preview | Related articles | Related reference work articlesPurchase $ 31.50348Cathode material influence on the power capability and utilizable capacity of next generation lithium-ionbatteriesJournal of Power Sources , Volume 195, Issue 12, 15 June 2010, Pages 3922-3927Michael A. Roscher, Jens Vetter, Dirk Uwe Sauer Show preview | Related articles | Related reference work articlesPurchase$ 41.95349Indicators for energy security Original Research Article Energy Policy , Volume 37, Issue 6, June 2009, Pages2166-2181Bert Kruyt, D.P. van Vuuren, H.J.M. de Vries, H.GroenenbergPurchase $ 19.95Show preview | Related articles | Related reference work articles35Mil operacionesbariátricas Original Research ArticleCirugía Española , Volume 79,Issue 6, June 2006, Pages349-355Aniceto Baltasar, Rafael Bou,Marcelo Bengochea, Carlos Serra, Nieves PérezShow preview | Related articles | Related reference work articlesPurchase $ 31.50results 326 - 35011,426 articles found for: pub-date > 2005 and tak((Targetlink or CATIA or UG or Simulink or SIL or MIL or model or develope or HIL or test) and ((Pure electric) or passenger or car or battery or equipment or material or "long-term") and ("long-term" or development or dynamic or market or application or promotion or access or customer) and (performance or security or function or design or company or industry or economic or cost or afford or popular or welcome or poor or rural or town or village or easy))动力软件工程师)|公司行业：电子技术/半导体/集成电路公司性质：国企公司规模： 500人以上比比你的竞争力发布日期： 2011-03-25 工作地点：上海-浦东新区招聘人数： 1 工作年限：三年以上语言要求：英语熟练学历：本科职位描述岗位职责:1)上位机软件和单片机软件的设计开发.2)单片机方案的应用3)生产工艺及流程文件的制订.4)产品的测试分析和试产.5)系统架构的搭建.6)与客户的项目沟通.岗位要求1)本科以上学历,汽车电子软件硬件相关专业.2)三年以上经验,较强的沟通和现场应变能力;熟练运用C语言,完成2款以上单片机程序开发,熟悉汽车CAN通信技术;熟悉ROHS、UL、ISO9001等,掌握至少一种视窗开发工具(如VC或LABVIEW等),有电源管理系统实际经验的为佳.对锂电池、铅酸电池等动力电池技术有一定的了解;对纯电动或混合动力汽车的电池管理系统有一定了解;有汽车特别是电动汽车/混合动力方面工作经历者优先.3)有团队合作意识,有良好的心态.联系方式电子邮箱：hr@比比你的竞争力『收藏』『推荐给朋友』『该公司其他职位』『打印』粉丝团(23)|上海长园维安电子线路保护股份有限公司（原名上海维安热电材料股份有限公司）创建于1996年。