nature上关于复合材料的英文文献

2020年在nature catalysis上发表重要成果
2020年，在《Nature Catalysis》杂志上发表了一项重要成果，该成果由某科研团队经过多年努力终于成功研发出一种新型的催化剂，能够有效地将废弃塑料转化为高附加值的产品。

这项成果的研发背景是，随着人类对塑料的依赖程度不断加深，废弃塑料的污染问题日益严重，给生态环境带来了巨大的压力。

因此，科研团队一直在寻找一种能够有效处理废弃塑料的方法。

该科研团队通过多年的研究，成功研发出这种新型催化剂。

该催化剂能够在常温常压下将废弃塑料中的聚乙烯和聚丙烯等塑料成分转化为燃料和化学品等高附加值的产品。

这种转化过程不仅能够有效处理废弃塑料，而且能够产生经济效益，具有很高的应用价值。

该成果的发表引起了广泛关注。

在《Nature Catalysis》杂志上，该论文被选为封面文章，并得到了编辑部的特别推荐。

该论文的发表不仅证明了该科研团队在催化剂研究方面的实力，也标志着人类在解决废弃塑料污染问题方面取得了重要进展。

未来，该科研团队将继续优化这种新型催化剂的制备工艺和应用范围，希望能够为解决全球废弃塑料污染问题做出更大的贡献。

同时，他们也希望通过与产业界的合作，将这种技术应用于实际生产中，为人类创造更加美好的生态环境和可持续发展未来。

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nature 环氧树脂和复合材料中c-o键的催化断开环氧树脂和复合材料中的C-O键是一种非常重要的化学键，它们在材料的性能中起着关键的作用。

然而，在某些情况下，需要催化剂来断开这些键，以便进行一些特定的化学反应。

本文将介绍环氧树脂和复合材料中C-O键的催化断开。

首先，让我们来了解一下环氧树脂和复合材料中的C-O键。

环氧树脂是一种重要的高分子材料，具有优异的物理性能和化学性能。

它通常由环氧树脂单体和固化剂组成，其中环氧树脂单体中含有若干个环氧基(-CH2-CHO-)，固化剂则含有若干个活性氢原子。

在反应中，环氧基和活性氢原子发生加成反应，形成C-O键，从而使环氧树脂单体和固化剂发生交联反应，形成三维网络结构。

这种网络结构赋予了环氧树脂材料优异的物理性能和化学性能。

复合材料是由多种不同的材料组成的材料，其中包括基质材料和增强材料。

基质材料通常是一种聚合物，如环氧树脂，而增强材料则可以是玻璃纤维、碳纤维、芳纶纤维等。

在复合材料中，C-O键通常是由基质材料中的环氧树脂单体和增强材料中的表面官能团发生反应形成的。

这种反应可以增强基质材料和增强材料之间的结合力，从而提高复合材料的强度、刚度和耐久性。

然而，在某些情况下，需要催化剂来断开环氧树脂和复合材料中的C-O键，以便进行一些特定的化学反应。

例如，在环氧树脂中加入一定量的酸类催化剂可以促进环氧基与活性氢原子之间的加成反应，从而加速固化过程。

在复合材料中，如果需要将增强材料与基质材料分离开来，可以使用一些特殊的催化剂来断开它们之间的C-O键。

常用的催化剂包括酸类催化剂、碱类催化剂、金属催化剂等。

酸类催化剂可以促进环氧基与活性氢原子之间的加成反应，从而加速固化过程。

碱类催化剂可以使环氧基发生开环反应，从而断开C-O键。

金属催化剂可以通过与环氧基形成配位键来促进开环反应。

这些催化剂在实际应用中都有广泛的应用。

总之，C-O键在环氧树脂和复合材料中起着非常重要的作用。

/Journal of Reinforced Plastics and Composites/content/30/19/1621The online version of this article can be found at:DOI: 10.1177/07316844114268102011 30: 1621 originally published online 7 November 2011Journal of Reinforced Plastics and Composites N. Venkateshwaran, A. ElayaPerumal and M. S. JagatheeshwaranEffect of fiber length and fiber content on mechanical properties of banana fiber/epoxy compositePublished by: can be found at:Journal of Reinforced Plastics and Composites Additional services and information for/cgi/alerts Email Alerts:/subscriptions Subscriptions: /journalsReprints.nav Reprints:/journalsPermissions.nav Permissions:/content/30/19/1621.refs.html Citations:What is This?- Nov 7, 2011OnlineFirst Version of Record- Dec 16, 2011Version of Record >>ArticleEffect of fiber length and fiber contenton mechanical properties of banana fiber/epoxy compositeN.Venkateshwaran,A.ElayaPerumal and M.S.JagatheeshwaranAbstractThe main factors that influence the properties of composite are fiber length and content.Hence the prediction of optimum fiber length and content becomes important,so that composite can be prepared with best mechanical prop-erties.Experiments are carried out as per ASTM standards to find the mechanical properties namely,tensile strength and modulus,flexural strength and modulus,and impact strength.In addition to mechanical properties,water absorption capacity of the composite is also studied.Further,fractured surface of the specimen are subjected to morphological study using scanning electron microscope.The investigation revealed the suitability of banana fiber as an effective reinforce-ment in epoxy matrix.Keywordspolymer composites,banana fiber,mechanical properties,scanning electron microscopeIntroductionNowadays,polymers are used everywhere in the day-to-day life.Plastics found its way when the need for low weight high strength material became important for various applications.The research in thefield of poly-mer and polymer-based components has gained wide-spread recognition owing to its property;however,its bio-degradability is still a matter of concern.Further, glassfiber reinforced polymers(GFRP)have become appealing substitutes for aluminum,concrete,and steel due to its high strength-to-weight ratio,ease of handling,and for being corrosion-free.Moreover, they can also be engineered to get the desired proper-ties.1Since large-scale production and fabrication of glassfiber causes environmental problems and also health hazards,a suitable alternate which is environ-mental friendly is the need of the hour.Naturalfibers that are low cost,lightweight and environmental friendly provide an excellent alternative to glassfiber. Joshi et al.2reviewed the life cycle assessment of natural fiber and glassfiber composite and found that natural fibers are environmentally superior to glassfiber,and also reduces the polymer content as reinforcement. Schmidt and Beyer,Wotzel et al.,and Corbiere et al.carried out some important works using the natural fibers as reinforcement in polymer matrix for use in automobile parts.Schmit and Beyer3have replaced the glassfiber polypropylene(PP)with hemp-PP com-posite for auto-insulation application.Wotzel et al.4 have used hemp-epoxy to replace glassfiber acryloni-trile butadien–styrene(ABS)for usage in auto-side panel.Similarly,Corbiere et al.5replaced glassfiber PP with Curaua PP for transporting pallet.All these studies revealed that the naturalfiber based polymer composite has successfully replaced the glassfiber. Pothan et al.6studied the effect offiber length and con-tent on the mechanical properties of the short banana/ polyester composite.Study shows that30–40mmfiber length and40%fiber loading provides better mechan-ical properties.Idicula et al.7investigated the mechan-ical performance of banana/sisal hybrid composite and Department of Mechanical Engineering,Anna University,Chennai,India.Corresponding author:N.Venkateshwaran,Department of Mechanical Engineering,Anna University,Chennai,IndiaEmail:venkatcad@Journal of Reinforced Plasticsand Composites30(19)1621–1627!The Author(s)2011Reprints and permissions:/journalsPermissions.navDOI:10.1177/0731684411426810the positive hybrid effect for tensile strength was found to be in the ratio of4:1(banana:sisal). Further,the tensile strength of the composite is better when bananafiber is used as skin and sisal as core material.Visco-elastic property of the banana/ sisal(1:1ratio)hybrid composite was studied by Idicula et al.8The study shows that sisal/polyester composite has maximum damping behavior and high-est impact strength as compared to banana/polyester and hybrid composite.Sapuan et al.9prepared the composite by reinforcing woven bananafibers with epoxy matrix.Tensile test result showed that the woven kind of reinforcement has better strength and the same was confirmed using Anova technique also. Venkateshwaran and ElayaPerumal10reviewed the various work in thefield of bananafiber reinforced with polymer matrix composite with reference to phys-ical properties,structure,and application. Venkateshwaran et al.11studied the effect of hybridi-zation on mechanical and water absorption properties. Investigation revealed that the addition of sisal in bananafiber composite upto50%increases the mechanical properties.Sapuan et al.12designed and fabricated the household telephone stand using woven banana fabric and epoxy as resin.Zainudin et al.13studied the thermal stability of banana pseudo-stem(BPS)filled unplastisized polyvinyl chlo-ride(UPVC)composites using thermo-gravimetric analysis.The study revealed that the incorporation of bananafiller decreases the thermal stability of the composite.Zainudin et al.14investigated the effect of bananafiller content in the UPVC matrix.The inser-tion offiller increases the modulus of the composite and not the tensile andflexural strength.Zainudin et al.15studied the effect of temperature on storage modulus and damping behavior of bananafiber rein-forced with UPVC.Uma Devi et al.16studied the mechanical properties of pineapple leaffiber rein-forced with polyester composite.Study found that optimum mechanical properties are achieved at 30mmfiber length and30%fiber content.Dabade et al.17investigated the effect offiber length and weight ratio on tensile properties of sun hemp and palmyra/polyester composite.The optimumfiber length and weight ratio were30mm and around 55%,respectively.From the above literatures,it is evident that the fiber length and content are the important factors that affect properties of the composite.Hence in this work,the effect offiber length and weight percentage on the mechanical and water absorption properties of the bananafiber epoxy composite is investigated. Further,the fractured surface of the composite are subjected to fractography study to evaluate the frac-ture mechanism.ExperimentalFabrication of compositeA molding box made of well-seasoned teak wood of dimensions300Â300Â3mm3is used to make a com-posite specimen.The top,bottom surfaces of the mold and the walls are coated with remover and kept for drying.Fibers of different length(5,10,15,and 20mm)and weight percentage(8,12,16,and20)are used along with Epoxy(LY556)and Hardener (HY951)for the preparation of composite.Testing standardsThe tensile strength of the composite was determined using Tinnus Olsen Universal Testing Machine (UTM)as per ASTM D638standard.The test speed was maintained at5mm/min.In this case,five specimens were tested with variedfiber length andfiber weight ratio.The average value of tensile load at breaking point was calculated.Theflexural strength was determined using the above-mentioned UTM as per ASTM D790procedure.The test speed was maintained between1.3and1.5mm/min. In this case,five samples were tested and the average flexural strength was reported.The impact strength of the composite specimen was determined using an Izod impact tester according to ASTM D256 Standards.In this case,five specimens were tested to obtain the average value.Figures1to5show the effect offiber length and weight content on ten-sile,flexural,and impact properties.Water absorp-tion behavior of banana/epoxy composites in water at room temperature was studied as per ASTM D570to study the kinetics of water absorption. The samples were taken out periodically andFigure1.Effect of fiber length and weight percentage on tensile strength.1622Journal of Reinforced Plastics and Composites30(19)weighed immediately,after wiping out the water from the surface of the sample and using a precise 4-digit balance to find out the content of water absorbed.All the samples were dried in an oven until constant weight was reached before immersing again in the water.The percentage of moisture absorption was plotted against time (hours)and are shown in Figures 6–13.Scanning electron microscopeThe fractured surfaces of the specimens were exam-ined directly by scanning electron microscope Hitachi-S3400N.The fractured portions of the sam-ples were cut and gold coated over the surface uni-formly for examination.The accelerating voltage used in this work was 10kV.Figures 14to 17show the fractured surface characteristics of the compositespecimen.Figure 6.Effect of moisture on fiber content;Fiber length –5mm.Figure 3.Effect of fiber length and weight percentage on flexural strength.Figure 2.Effect of fiber length and weight percentage on tensilemodulus.Figure 4.Effect of fiber length and weight percentage on flexuralmodulus.Figure 5.Effect of fiber length and weight percentage on impact strength.Venkateshwaran et al.1623Figure 12.Effect of moisture on fiber length;Fiber wt%–16.Figure 7.Effect of moisture on fiber content;Fiber length –10mm.Figure 11.Effect of moisture on fiber length;Fiber wt%–12.Figure 10.Effect of moisture on fiber length;Fiber wt%–8.Figure 8.Effect of moisture on fiber content;Fiber length –15mm.Figure 9.Effect of moisture on fiber content;Fiber length –20mm.1624Journal of Reinforced Plastics and Composites 30(19)Results and discussion Mechanical propertiesFor the tensile test,composite specimens are made of fibers of different length (5,10,15,and 20mm)and weight ratio (8,12,16,and 20)were used to calculate the tensile strength.Figures 1and 2show the effect of fiber length and weight ratios on tensile strength and modulus of the composite,respectively.Figure 1shows that the increase in fiber length and weight ratio increases the tensile strength and modulus upto 15mm fiber length and 12%weight ratio.Further increases cause the properties to decrease because of lower fiber–matrix adhesion and the quantity of fiber content being more than matrix.From Figures 1and 2,the maximum tensile strength and modulus oftheFigure 14.SEM micrograph of tensile fracturedspecimen.Figure 15.SEM micrograph of fractured specimen under flexuralload.Figure 16.SEM micrograph of fractured specimen under impactload.Figure 17.Micrograph of poorinterface.Figure 13.Effect of moisture on fiber length;Fiber wt%–20.Venkateshwaran et al.1625composite are16.39MPa and0.652GPa,respectively for thefiber length of5mm and12%weight ratio. Flexural strength and modulus for differentfiber lengths(5,10,15,and20mm)and weight ratios(8, 12,16,and20)are shown in Figures3and4,respec-tively.It was found that the maximumflexural strength and modulus are57.53MPa and8.92GPa,respectively for thefiber length of15mm andfiber weight of16%.The results of the pendulum impact test are shown in Figure 5.As thefiber weight and length increases impact strength also increases upto16%fiber weight ratio and then begin to decrease.The maximum impact strength of 2.25J/m was found for thefiber length 20mm and16%fiber weight.Although the variousfiber lengths and weight per-centage provides the maximum mechanical properties, from Figures10,12,and14it can be concluded that the optimumfiber length andfiber weight percentage is 15mm and16%respectively as the properties variation with15mm and16%are negligible when compared to the maximum mechanical properties provided by differ-entfiber lengths and weight percentage indicated as above.The mechanical properties provided above are better than coir18and palmyra.19Water absorption studyThe effects offiber length and content on the water absorption study are shown in Figures6–13.Figures 6to9show the effect offiber content on the water absorption property of the banana/epoxy composite. It shows that as thefiber content increases the moisture uptake of the composite also increases.This is due to the affinity of the bananafiber towards the moisture. The maximum moisture absorption for the composite is around5%for all length and weight percentage of composite.Figures10to13show the effect offiber length on the water uptake capability of composite.It indicates that the variation of length(5,10,15,and 20mm)does not have much impact as compared with thefiber content.The moisture absorption percentage of bananafiber/epoxy composite seems to be lesser than hempfiber20andflaxfiber21composite. Fractography studyMicrographs of fractured tensile,flexural,and impact specimens are shown in Figures14–17.Figure14shows the micrograph of fractured surface of specimen under tensile load.It clearly indicates that the failure is due to fiber pull out phenomenon.Figure15shows the frac-tured surface of the specimen under bending load. Micrograph also shows the bending offibers due to the application of load.Figure16shows the failure of the composite under impact load.Further,it also shows the striation occurring on the matrix surface and the presence of hole due tofiber pull out.Figure17shows the micrograph of20mmfiber length and20%fiber weight composite specimen.It clearly indicates that the clustering offibers result in poor interface with matrix,and in turn decreases the mechanical properties of the composite.ConclusionBased on thefindings of this investigation the following conclusions can be drawn:.The optimumfiber length and weight ratio are 15mm and16%,respectively for bananafiber/ epoxy composite..Moisture absorption percentage of banana/epoxy composite for all length and weight percentage is around5..Also,the moisture uptake capability of the compos-ite is greatly influenced byfiber content than length. .SEM image shows that increasing thefiber content above16%results in poor interface betweenfiber and matrix.References1.Houston N and Acosta F.Environmental effect of glassfiber reinforced polymers.In:Proceedings of2007Earth Quake Engineering Symposium for Young Researcher, Seattle,Washington,2007.2.Joshi SV,Drzal LT,Mohanty AK and Arora S.Are nat-ural fiber composites environmentally superior to glass fiber reinforced posite Part A2004;35: 371–376.3.Schmidt WP and Beyer HM.Life cycle study on a naturalfiber reinforced component.In:SAE Technical Paper 982195.SAE Total Life-Cycle Conference,1–3 December,1998,Graz,Austria.4.Wotzel K,Wirth R and Flake R.Life cycle studies onhemp fiber reinforced components and ABS for automo-tive parts.Die Angewandte Makromolekulare Chemie1999;272:121–127.5.Corbiere-Nicollier T,Laban BG and Lundquist.Lifecycleassessment of bio-fibers replacing glass fibers as reinforce-ment in plastics.Resour Conserv Recycl2001;33:267–287.6.Pothan LA,Thomas S and Neelakantan NR.Shortbanana fiber reinforced polyester composites:mechanical, failure and aging characteristics.J Reinf Plast Compos 1997;16:744–765.7.Idicula M,Neelakantan NR and Oommen Z.A study ofthe mechanical properties of randomly oriented short banana and sisal hybrid fibre reinforced polyester compos-ites.J Appl Polym Sci2005;96:1699–1709.1626Journal of Reinforced Plastics and Composites30(19)8.Idicula M,Malhotra SK,Joseph K and Thomas S.Dynamic mechanical analysis of randomly oriented short banana/sisal hybrid fibre reinforced polyester pos Sci Technol2005;65:1077–1085.9.Sapuan SM,Leenie A,Harimi M and Beng YK.Mechanical property analysis of woven banana/epoxy composite.Mater Design2006;27:689–693.10.Venkateshwaran N and ElayaPerumal A.Banana fiberreinforced polymer composite-a review.J Reinf Plast Compos2010;29:2387–2396.11.Venkateshwaran N,ElayaPerumal A,Alavudeen A andThiruchitrambalam M.Mechanical and water absorption behavior of banana/sisal reinforced hybrid composites.Mater Design2011;32:4017–4021.12.Sapuan SM and Maleque MA.Design and fabrication ofnatural woven fabric reinforced epoxy composite for household telephone stand.Mater Design2005;26: 65–71.13.Zainudin ES,Sapuan SM,Abdan K and MohamadMTM.Thermal degradation of banana pseudo-stem fibre reinforced unplastisized polyvinyl chloride compos-ites.Mater Design2009;30:557–562.14.Zainudin ES,Sapuan SM,Abdan K and MohamadMTM.The mechanical performance of banana pseudo-stem reinforced unplastisized polyvinyl chloride compos-ites.Polym Plast Technol Eng2009;48:97–101.15.Zainudin ES,Sapuan SM,Abdan K and MohamadMTM.Dynamic mechanical behaviour of bananapseudo-stem filled unplasticized polyvinyl chloride com-posites.Polym Polym Compos2009;17:55–62.16.Uma Devi L,Bhagawan SS and Sabu Thomas.Mechanical properties of pineapple leaf fiber-reinforced polyester composites.J Appl Polym Sci1997;64: 1739–1748.17.Dabade BM,Ramachandra Reddy G,Rajesham S andUdaya kiran C.Effect of fiber length and fiber weight ratio on tensile properties of sun hemp and palmyra fiber reinforced polyester composites.J Reinf Plast Compos 2006;25:1733–1738.18.Harish S,Peter Michael D,Bensely A,Mohan Lal D andRajadurai A.Mechanical property evaluation of natural fiber coir composite.Mater Characterisation2009;60: 44–49.19.Velmurugan R and Manikandan V.Mechanical proper-ties of palmyra/glass fiber hybrid posite Part-A2009;38:2216–2226.20.Dhakal HN,Zhang ZY and Richardson MOW.Effect ofwater absorption on the mechanical properties of hemp fibre reinforced unsaturated polyester composites.Compos Sci Technol2007;67:1674–1683.21.Alix S,Philippe E,Bessadok A,Lebrun V,Morvan V andMarais S.Effect of chemical treatments on water sorption and mechanical properties of flax fibres.Bioresour Technol2009;100:4742–4749.Venkateshwaran et al.1627。

j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y207(2008)1–12j o u r n a l h o m e p a g e:w w w.e l s e v i e r.c o m/l o c a t e/j m a t p r o t ecReviewReview of warm forming of aluminum–magnesium alloysSerkan Toros,Fahrettin Ozturk∗,Ilyas KacarDepartment of Mechanical Engineering,Nigde University,51245Nigde,Turkeya r t i c l e i n f oArticle history:Received31October2007 Received in revised form11March2008Accepted31March2008Keywords:Warm formingAluminum–magnesium(Al–Mg) alloys5XXX series a b s t r a c tAluminum–magnesium(Al–Mg)alloys(5000series)are desirable for the automotive industry due to their excellent high-strength to weight ratio,corrosion resistance,and weldability. However,the formability and the surface quality of thefinal product of these alloys are not good if processing is performed at room temperature.Numerous studies have been conducted on these alloys to make their use possible as automotive body materials.Recent results show that the formability of these alloys is increased at temperature range from200 to300◦C and better surface quality of thefinal product has been achieved.The purpose of this paper is to review and discuss recent developments on warm forming of Al–Mg alloys.©2008Elsevier B.V.All rights reserved.Contents1.Introduction (1)2.Aluminum for passenger vehicles (2)3.Formability of aluminum–magnesium sheets (5)3.1.The effects of blankholder force and drawbead geometry (5)3.2.The effects of temperatures and strain rates (6)3.3.The effects of lubrication (10)4.Conclusion (10)Acknowledgements (10)References (10)1.IntroductionAluminum alloys are produced and used in many forms such as casting,sheet,plate,bar,rod,channels and forgings in various areas of industry and especially in the aerospace∗Corresponding author.Tel.:+903882252254.E-mail address:fahrettin@.tr(F.Ozturk).industry.The advantages of these alloys are lightweight, corrosion resistance,and very good thermal and electrical conductivity.The aforementioned factors plus the fact that some of these alloys can be formed in a soft condition and heat treated to a temper comparable to structural steel make0924-0136/$–see front matter©2008Elsevier B.V.All rights reserved. doi:10.1016/j.jmatprotec.2008.03.0572j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y207(2008)1–12it very attractive for fabricating various aircraft and missile parts.The present system utilized to identify aluminum alloys is the four digit designation system.The major alloy element for each type is indicated by thefirst digit,i.e.,1XXX indicates aluminum of99.00%minimum;2XXX indicates that copper is the main alloying element.Manganese for 3XXX,silicon for4XXX,magnesium for5XXX,magnesium and silicon for6XXX,zinc for7XXX,lithium for8XXX,and unused series for9XXX are main alloying elements.In industry,low carbon steels have been commonly used for a long time due to their excellent formability at room temperature,strength,good surfacefinish,and low cost.How-ever application of the aluminum and its alloys in thisfield were ranked far behind steels because of cost and formability issues,despite their high-strength-to-weight ratio and excel-lent corrosion resistance.For expanding use of aluminum alloys or replacing steels in many areas,however,there have been challenging formability problems for aluminum alloys to overcome.The formability of the aluminum alloys at room temperatures is generally lower than at both cryogenic and elevated temperatures.At cryogenic temperatures,the ten-sile elongation is significantly increased for many aluminum alloys especially5XXX series alloys and is related to the enhancement of work hardening,while at elevated temper-atures it is mainly due to the increased strain rate hardening. Forming at cryogenic temperatures is technologically more challenging than at high temperatures.At hot forming tem-peratures,other issues should also be taken into consideration such as creep mechanisms which may affect the forming deformation and cavitations at grain boundaries which may induce premature failure at low strain rates.2.Aluminum for passenger vehiclesLightweight vehicles have become a key target for car man-ufacturers due to increasing concerns about minimizing environmental impact and maximizing fuel economy without sacrificing the vehicle performance,comfort,and marketabil-ity(Cole and Sherman,1995).Aluminum will probably play an important role in the future car generations.Its material properties give some advantages and open the way for new applications in the automotive industry(Carle and Blount, 1999).As a result of the developments in the aluminum indus-try,improving the mechanical properties of the aluminum alloys by adding various alloying elements increased the application area of these alloys in automotive and aerospace industries(Richards,1900).Design of aluminum structures can also have a big influence on the sustainability of a car. Some of the important design aspects of a car which influence the environment are weight,aerodynamic and roll-resistance. DHV Environment and Transportation Final Report indicates that the material has a big influence on the car weight. (DHV Environment and Transportation Final Report,2005). Lightweight car consumes less material resources in the long run(300,000km),although it would cost about30%more than the conventional car.Therefore,its production would decrease employment in the car industry by about4%over a decade while increasing the employment in the short term (Fuhrmann,1979).Fig.1–Average use of aluminum(International Aluminum Institute(IAI),2002;Martchek,2006;Mildenberger and Khare,2000;Schwarz et al.,2001).Aluminum alloys are effective materials for the reduction of vehicle weight and are expanding their applications.Fig.1 illustrates the usage of aluminum for European and Ameri-can vehicles over years.In addition to USA and Europe,Japan has recently increased their aluminum alloy usage.Analysts expect that the aluminum alloys usage in Japan Automotive Industry will reach1.5million tons by2010.Assuming vehi-cle production holds steady at around10million units,the average yearly growth will be around2.5%(McCormick,2002). As shown in Fig.1,the amount of aluminum used in1960 is substantially low.The main reasons are forming difficul-ties of aluminum alloys at that time and the smaller range of alloys available.The demand for aluminum alloys as light weight materials has increased in recent years.Fig.2demon-strates the amount of produced aluminum products in the world.In the past,the main aluminum products were produced by casting such as engines,wheels,exhaust decor;how-ever nowadays wrought aluminum products arefinding more applications in sheets including exterior panels such as hoods and heat insulators,in extrusions including bumper beams, and in forgings including suspension parts Fig.2.One of the most important benefits of using aluminum alloys in automotive industry is that every kg of aluminum, which replaces2kg of steel,can lead to a net reduction of 10kg of CO2equivalents over the average lifetime of a vehi-cle(Ungureanu et al.,2007).In Fig.3,the effects of the car components on CO2emissions are shown.CO2emissionisFig.2–Aluminum products for automobile over years(Cole and Sherman,1995;Inaba et al.,2005;Patterson,1980; Miller et al.,2000;T urkish Statistical Institute,2004).j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y207(2008)1–123Fig.3–Effect of technical measures on the CO2emission(Mordike and Ebert,2001).1consumption(Mordike and Ebert,2001)Measure taken Potential saving(%)fuel Importance innovative materialsShort-medium term Long termLight constructions3–510–15++C w value24–6+Motor\gear control510±Resistance to rolling1–23+Motor preheating24–6±Equipment24±critical in terms of environmental pollution.Schwarz et al. (2001)inspected that the relationship between the usage of the aluminum products in new designs and the CO2emission and emphasized that the CO2emission ratio could be reduced by using lightweight materials such as aluminum in new transportation designs.Weight reduction of the car’s compo-nents influences fuel consumption considerably.In Table1, the effect of weight reduction on fuel savings is seen.Fuel economy improvements of around6–8%or as much as2.5 extra miles per gallon can be realized for every10%in weight reduction(Mordike and Ebert,2001).Recyclability of alloys has also become an important issue in view of energy and resource conservation.For example, recycling potential of the aluminum products is much bet-ter than the ferrous metals.Martchek(2006)and Mildenberger and Khare(2000)investigated the recycling potential and nec-essary energy to reproduce the aluminum products.According to Martchek(2006),increasing the recycled metal usage in the aluminum production consumes less energy and emits less greenhouse gas to produce the aluminum ingots.Sillekens et al.(1997)investigated the formability of recycled aluminum alloy5017.In their study,they focused on changes in the amounts of alloying elements(particularly iron)to see how they affect the formability of products.It is observed that the change in the iron content does not lead to a dramatic degen-eration in the performance of the material.Aluminum alloy sheets are widely used in the car,ship-building and aerospace industries as substitutes for steel sheets andfiber reinforced plastic(FRP)panels,due to their excellent properties such as high-strength,corrosion resis-tance,and weldability(Naka et al.,2001).The features of the most used aluminum–magnesium alloys in automotive application were summarized in Table2.Figs.4and5illus-trates aluminum and other materials usages in automotive and aerospace industry,respectively.Magnesium is one of the most effective and widely used alloying elements for aluminum,and is the principal element in the5XXX series alloys.These alloys often contain small additions of transition elements such as chromium or man-ganese,and less frequently zirconium to control the grain or subgrain structure and iron and silicon impurities that are usually present in the form of intermetallic particles(ASM2of several Al–Mg alloysStrength Formability Resistance to corrosion WeldabilityExcellent5454,5652––5454,5652 Highest5052–––High5456–54565083,5456 Good5154,52545005,5050,50835005,5050,5083,5254,56525154,5254,5557j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 207(2008)1–125P o l y c a r b o n a t e 12006554166,72,4–122110B S :115400P o l y e s t e r ,P B T 134********,82,4–135150S h o r e D :65380P o l y e s t e r e l a s t o m e r 12004537500,00,2–130500S h o r e D :30–82520A d h e s i v e s 10305553398,13,1–1252–250R u b b e r ,P V C 14404027777,83,1–13540S h o r e D :74–88120R u b b e r 9101516483,60,5–162600S h o r e A :30–90130T i –6A l –4V c4730900190274,81202359,510B S :3345845C u e x t r u t e d 891039043771,11204416.94B S :902323E p o x y /g l a s s S M C 1600260162500,0––3,65–––H 254p o l y e s t e r l a m i n a t e 16004125625,0––0,29–B a r c o l :30–T h e r m o s e t p o l y e s t e r 182********,8––1,2–B a r c o l :32–E p o x y /g l a s s S M C 1600260162500,0––3,65–––H 254p o l y e s t e r l a m i n a t e 16004125625,0––0,29–B a r c o l :30–aY i e l d p o i n t .b E l o n g a t i o n a t b r e a k %.cH e a t -t r e a t m e n t .Fig.4–Al alloys and its application for automotive industry (Sherman,2000;White,2006).Fig.5–Current material usages for Boing 757(Moscovitch,2005).Metal Handbook,1988).When magnesium is used as the major alloying element or combined with manganese,the result is a moderate to high-strength,non-heat-treatable alloy.Alloys in this series are readily weldable and have excellent resis-tance to corrosion,even in marine applications.Selection of suitable aluminum alloys,for several applications,requires a basic knowledge of heat treatment,corrosion resistance,and primarily,mechanical properties.Table 3summarizes features and applications of Al–Mg alloys.Three different material groups,their properties and applications were compared for material selection.3.Formability of aluminum–magnesium sheets3.1.The effects of blankholder force and drawbead geometryT ypical sheet metal forming processes are bending,deep drawing,and stretching.If a doubly curved product must be made from a metal sheet,the deep drawing process or the stretching process is used.The deep drawing process can reach production cycles of less than 10s,and is hence a suitable process for mass production.In deep drawing and stretching,the stresses normal to the sheet are usually very small compared to the in-plane stresses and are therefore neglected.T wo important failure modes limit the applicabil-ity of the deep drawing and stretching process:necking and wrinkling.Both are closely related to the material properties.The ability to accurately predict the occurrence of wrinkling is critical in the design of tooling and processing parame-ters (Xi and Jian,2000)like sheet thickness,blankholder force,6j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y207(2008)1–12and the local curvature of the sheet(Hutchinson and Neale, 1985).Ahmetoglu et al.(1997)determined wrinkling and frac-ture limits and developed blank holding force(BHF)control to eliminate these defects,improve part quality and increase the formability.A computer simulation model was developed to control aforesaid parameters.Jinta et al.(2000)examined wrinkle behavior in5XXX and6XXX series aluminum alloys and compared the results with wrinkle behavior of steel.Their results indicate that aluminum alloys generally forms more wrinkles then steel,especially6XXX series aluminum alloy has a tendency to more wrinkles than5XXX series.Gavas and Izciler(2007)examined the effect of blank holder gap(BHG) on deep drawing of square cup in order to investigate wrin-kling,tearing,and thickness distribution.As a result of their study,they observed that increasing of the BHG allows more material to be drawn into the die cavity without tearing or shape distortions.It is also noticed that it was impossible to use too large BHG because of excessive wrinkling and buck-ling at the straight sides which cause tearing.Lin et al.(2007) determined the drawing limit under constant and variable BHF.Drawbeads are directly related with wrinkling behavior of the materials.They are used to control theflow of sheet metal into the die cavity during the stretch forming of large panels. Beside that they reduce the BHF and minimize the blank size needed to make a part(Demeri,1993).The shape and position of drawbead and the amount of force on it are very impor-tant in terms of part quality.Samuel(2002)investigated the influence of drawbead geometry on the drawbead restraining force(DBRF)and BHF numerically and experimentally for alu-minum alloy.In his study,two kinds of drawbead geometry which are square female and round female were investigated. As a result,it is obtained that the DBRF and BHF for the square female bead are higher than those for the rounded female bead.He emphasized that this discrepancy are occurred due to the sharp corners.It is also observed that the total equiva-lent plastic strain and von Mises stresses at upper and lower surfaces of square female drawbead are higher than those for the round female drawbead.3.2.The effects of temperatures and strain ratesAlthough the aluminum alloys have high-strength to weight ratio and good corrosion resistance,the low formability of aluminum sheets limits their use in some products with complex shapes,such as automotive body parts.The warm forming process is intended to overcome this problem by using an elevated forming temperature which is below the recrystallization temperature(Tebbe and Kridli,2004).A typ-ical warm forming experimental set-up is shown in Fig.6.In the warm forming set-up,dies and blank holders are heated to 200–300◦C.In order to heat dies and blank holders,electrical heating rods that are located in these parts are used but there is a risk of necking during heating and cooling.Warm form-ing was studied for many years,e.g.in the1970s and1980s by Shehata et al.(1978)and Wilson(1988)with increasing atten-tion being dedicated to the subject in the last decade.The warm forming method improves the formability of the aluminum alloys.This improvement at the elevated tem-peratures is principal for the aluminum alloys such as5082 and5005alloys due to the increased strain ratehardening Fig.6–A typical warm forming set-up(Palumbo and Tricarico,2007).(Shehata et al.,1978).Schmoeckel(1994)and Schmoeckel et al.(1995)investigated the drawability of5XXX series alloys at the elevated temperatures.Temperature has a significant influence on the stamping process.Further investigation on forming showed that the formability with a partial heating in the holder or matrices area was much better when compared with the homogeneously heated tools(Schmoeckel,1994). Schmoeckel et al.(1995)showed that a significant increase in the limiting drawing ratio(LDR)for the aluminum alloy AlMg4.5Mn0.4can be achieved by a heated and lower strain rated hydromechanical stamp.Modeling of the deep draw-ing with a rotationally symmetrical tool(stamp diameter: 100mm)which was cooled from the stamp side by additional air ensured an increase in LDR.It was demonstrated that the formability is improved by a uniform temperature increase,but the best results are obtained by applying temperature gradients.The formabil-ity depends strongly on the composition of the aluminum alloy.Aluminum–magnesium alloys have a relatively good formability.A disadvantage is that these alloys suffer from stretcher lines,which gives an uneven surface after deforma-tion.Because of this reason,5XXX series aluminum is used for inner panels of vehicles.These undesired surface defects can be eliminated by the forming processes at the elevated temperatures(Van Den Boogaard et al.,2001).The aluminum which contains6%magnesium could give a300%total elonga-tion at about250◦C,finds more application in industry(Altan, 2002).Yamashita et al.(2007)numerically simulated circu-lar cups drawing process by using Maslennikov’s technique (Maslennikov,1957)which is also called“punchless drawing”. In this production technique,a rubber ring is used instead of the rigid punch.Browne and Battikha(1995)optimized the formability process by using aflexible die and optimized the process parameters to ensure a defect-free product.To accurately simulate warm forming of aluminum sheet, a material model is required that incorporates the temper-ature and strain-rate dependency(Van Den Boogaard and Hu´etink,2004).Because of this,the effect of temperature distribution on warm forming performance is very impor-tant.Van Den Boogaard and Hu´etink(2006)observed that the formability of the Al–Mg alloy sheets can be improved by increasing the temperature in some parts of the sheet andj o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y207(2008)1–127cooling the other parts when simulated by the cylindrical cup deep drawing at the different temperature gradients of the tools and the blanks.Chen et al.(2006)investigated combined isothermal/non-isothermalfinite element analysis(FEA)with design of experiments tools to predict appropriate warm form-ing temperature conditions for5083-O(Al–Mg)sheet metal blanks,deep drawing and two-dimensional stamping cases. To achieve increased degrees of forming,different tempera-ture levels should be assigned to the corner and body of the die and punch.25–250%elongation ranges were seen.They found that the formability of Al-5083alloy is greatly depen-dent on the temperature distribution of the die and punch. It is also observed that the optimal temperature distributions for warm deep drawing and warm two-dimensional stamping were not identical.Naka et al.(2003)investigated the effects of temperature on yield locus for5083aluminum alloy sheet.In their study, they have tried to determine the optimum condition of press forming for an aluminum alloy sheet,the effect of form-ing temperature on the yield locus.They obtained the yield locus for afine grain Al–Mg alloy(5083-O)sheet by performing biaxial tensile tests,using cruciform specimens,at temper-atures of30,100,170,250,and300◦C at10s−1strain rate. As expected,the size of the yield locus drastically decreased with increasing temperature.This can be exploited to improve press operations.Naka and Yoshida(1999)investigated the effects of temperature and forming rates on deep drawabil-ity of5083Al–Mg alloy.In their study,different temperature gradients from room temperature to180◦C and the form-ing speeds between0.2and500mm/min were performed. Results show that LDR increases mostly with increasing the die temperature because the deformation resistance inflange shrinkage decreases with an increase in temperature.Beside that the LDR becomes smaller with increasing the forming speed at all temperatures since theflow stress of the heated blank(at theflange)increases with increasing the strain rate. Moreover,the cooled blank at the punch corner becomes less ductile.Another comprehensive study for5XXX series alu-minum alloys was done by Bolt et al.(2001).In that study, the formability was compared for1050,5754and6016type aluminum alloys from100to250◦C by using the both box shaped and conical rectangular products.They observed that the minimum die temperature of6016alloy on the die pro-cess limits was lower than that of the5754alloy.Smerda et al.(2005)investigated the strain rate sensitivity of5754and 5182type aluminum alloy sheets at room temperature and elevated temperatures.In their study,the split Hopkinson bar apparatus was used to identify the constitutive response and the damage evolution in the aluminum alloys at high strain rates of600,1100and1500s−1.It was observed that the qua-sistatic and dynamic stress strain responses in the range of strain rates and temperatures were low for both alloys.AA5754 exhibited a mild increase inflow stress with strain rate,while AA5182appeared to be strain rate insensitive.The ductility of the materials showed little differences in the tempera-ture range between23and150◦C at a strain rate of1500s−1. However,thefinal elongation was decreased for both alu-minum alloys at300◦C and a strain rate of1500s−1when compared to that at lower temperatures.Picu et al.(2005) investigated the mechanical behavior of the commercial alu-minum alloy AA5182-O.The dynamic strain aging effect was observed at all temperature between−80and110◦C and at strain rates lower than10−1s−1.In addition,the strain rate sensitivity parameter was also determined as a function of temperature and plastic strain.Abedrabbo et al.(2007)devel-oped a temperature-dependent anisotropic material model for FEA and formability simulation for two automotive aluminum alloys,AA5182-O and AA5754-O.Multiple temperatures to simulate the formability of more complex automotive parts, where the temperatures of the different sections will be deter-mined automatically,can be found by this model.In addition to the temperature,the forming speed controlling the strain rate,the die and stamp corner radii and other geometric parameters of the die set-up determine the forming charac-teristics of aluminum alloy sheets.In thefinite element simulations,material models are quite important in order to evaluate accurately the formability of aluminum alloy sheets.Barlat models are commonly used to define aluminum alloy behaviors.Barlat and Lian(1989)developed a yield function that described the behavior of orthortropic sheets and metals exhibiting a planar anisotropy and subjected to plane stress conditions.This yield function showed similar results cal-culated by the Taylor/Bishop and Hill models.Barlat et al. (1991)extended this method to triaxial loading conditions by using a six-component yield function.Lian et al.(1989) used this yield function to study the influence of the yield surface shape on failure behavior of sheet metals.Yu et al. (2007)developed a ductile fracture criterion which is intro-duced by afinite element simulation.They carried out the simulations of aluminum alloy sheet forming based on Bar-lat’s yield function(Barlat and Lian,1989)and Hollomon’s hardening equation.In their study,the critical punch strokes of the aluminum alloy sheets of X611-T4,6111-T4and5754-O in a cylindrical complex forming in which deep drawing and stretching modes were calculated by the ductile fracture criterion.The results showed good agreements with the exper-imental results.Barlat et al.(1997)measured the yield surfaces for binary aluminum–magnesium sheet samples with differ-ent microstructures.A generalized plastic behavior of any aluminum alloy sheet yield description was proposed to pre-dict the behavior of the solute strengthened(precipitation hardened)aluminum alloy sheets.Barlat et al.(2003)proposed a plane stress yield function that describes the anisotropic behavior of the sheet metals,in particular,for aluminum alloy sheets.The anisotropy of the function was introduced in the formulation using two linear transformations on the Cauchy stress tensor.For the Al–5wt.%Mg and6016-T4alloy sheet samples,yield surface shapes,yield stress and r-value directionalities were compared with those of previously sug-gested yield functions by Yoon et al.(2004).Barlat et al.(2005) proposed anisotropic yield functions based on linear transfor-mations of the stress deviator in general terms.T wo specific convex formulations were given to describe the anisotropic behavior of metals and alloys for a full stress state(3D).Choi et al.(2007)developed analytical models for hydro-mechanical deep drawing tests to investigate the effects of process con-ditions such as temperature,hydraulic pressure,BHF and forming speed.According to their models,the experimental results show good agreement with the FE models.One of the8j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y207(2008)1–12most important problems in forming simulation programs is the method by which the defects are analyzed during the simulation.Forming conditions of the sheet metals were also investigated in implicit and explicitfinite element simulations by Van Den Boogaard et al.(2003).The study showed that the computation time for implicitfinite element analyses tended to increase disproportionally with increasing problem size.Sheet metal deformation is considered biaxial rather than tensile deformation.For this reason,biaxial data in material model should be evaluated.Li and Ghosh(2003)studied uni-axial tensile deformation behavior of three aluminum sheet alloys,Al5182+1%Mn,Al5754and Al6111-T4in the warm forming temperature range of200–350◦C and in the strain rate range of0.015–1.5s−1.It is found that the total elonga-tion in uniaxial tension increased with increasing temperature and decreased with increasing strain rate.They contributed to the enhanced ductility at elevated temperatures primarily from the post-uniform elongation which becomes dominant at elevated temperatures and/or at slow strain rates.The enhancement of strain rate sensitivity with increasing tem-perature accounts for the ductility improvement at elevated temperatures.In their study,the uniaxial tensile test is used as a screening test for ranking the relative formability among different sheet alloys.The strain hardened5XXX alloys(Al 5182+Mn and Al5754)have shown better formability than the precipitation hardened alloy(Al6111-T4).Li and Ghosh (2004)also investigated biaxial warm forming behavior in the temperature range200–350◦C for three automotive aluminum sheet alloys:Al5754,Al5182containing1%Mn and Al6111-T4.Formability was studied by forming rectangular parts at a rapid rate of1s−1using internally heated punch and die for both isothermal and non-isothermal conditions.It is observed that the formability for all the three alloys improves at ele-vated temperatures,the strain hardened alloys Al5754and Al 5182+Mn show considerably greater improvement than the precipitation hardened alloy Al6111-T4.Results show that temperature effects on drawing of the sheet have a large effect on formability.Setting die temperature slightly higher than punch temperature was favorable in promoting forma-bility.They also determined the forming limit diagram(FLD) under warm forming conditions which showed results con-sistent with the evaluation of part depth.Fig.7shows that the effects of temperature on FLDs of type5754,5182and 6111-T4aluminum alloy.As seen in thefigures,theformabil-Fig.7–The effect of warm temperatures on FLDs(Li and Ghosh,2004).。

TEM study of carbon®bre reinforced aluminium matrix composites: in¯uence of brittle phases and interface on mechanical propertiesncin*,C.MarhicPhysique Cristalline,Institut des MateÂriaux Jean Rouxel,BP32229Nantes cedex3,FranceReceived13May1999;received in revised form4November1999;accepted8November2000AbstractThe microstructure of A357aluminium alloy(7wt%Si+0.6wt%Mg)reinforced by1D-M40J carbon®bres is characterised using di erent techniques of transmission electron microscopy(di raction,HRTEM,EDX,EELS).The microstructure of the high modulus PAN based®bre and of the pyrolytic carbon coating(Cp)is fully characterised.Silicon and Mg2Si grains which determine matrix-reinforcement adhesion are investigated.On the basis of the microstructural features,the mechanical properties of the composites are discussed.The mechanical behaviour of composites prepared with and without Cp interphases corresponds to a brittle matrix reinforced by brittle®bres.In the case of the composite without Cp interphase,the most in¯uent parameter is the high resistance to sliding at the interface between silicon and®bres which leads to a strong®bre-matrix``bonding''and thus,to a weak and brittle material.The interfacial resistance to decohesion and to sliding is lower in the composite with Cp interphase resulting in higher strength and limited pull out.This lower interfacial resistance is due to the successive microporous and layered micro-structures of the pyrolytic carbon coating.#2000Elsevier Science Ltd.All rights reserved.Keywords:Al alloy;Al4C3;Carbon®bre;Composites;Electron microscopy;Interphase;Mechanical properties1.IntroductionCarbon®bres are used to increase strength and sti -ness of aluminium alloys while keeping low weight and good thermal and electrical conductivity.However, brittle phases which form or precipitate during the pro-cessing may have a deleterious in¯uence on the mechanical properties of the composites.1À11Studies are being performed to improve the processing and to obtain the required properties.C/Al composites pre-pared at ONERA-Chatillon for aerospace applications contain a high volume fraction of®bres(70%).12,13In one of these,an A357aluminium alloy is reinforced with high modulus M40J carbon®bres,sometimes coated with a pyrolytic carbon interphase(Cp).The M40J/A357composite without the Cp interphase is brittle and weak(%700MPa).13The M40J/Cp/A357 composite with the Cp interphase is stronger(%1360 MPa)but less so than expected from the law of mixtures and,moreover,it is rather brittle(limited pseudo-plastic deformation,low work of fracture),despite some®bre debonding.14Scanning electron microscopy(SEM)did not reveal any reaction layer in either type of composite.13 A transmission electron microscopy(TEM)study was thus performed to reach a full knowledge of the microstructural features of the materials and a better understanding of their mechanical behaviour.The mechanical tests and properties are fully described in a paper previously published.14In this paper,the microstucture of the high modulus carbon®bre and of the carbon coating is characterised. The amounts and locations of the brittle phases,alumi-nium carbides and silicon,are determined.Interdi u-sion,which is known to in¯uence the failure resistance at the interfaces,is also studied.The discussion of the microstructural features reveals important parameters and shows that the mechanical behaviour of the compo-sites does not depend on the metal matrix.Models devel-oped for ceramic±ceramic composites are successfully used to analyse the mechanical properties.2.Materials and experimental procedure1D-composites were fabricated with A357alloy (alloying elements:Si=7wt%,Mg=0.6wt%)and M40J®bres('r=4400MPa,E=377GPa)-by liquid metal in®ltration(LMI),under moderate pressure0955-2219/00/$-see front matter#2000Elsevier Science Ltd.All rights reserved. P I I:S0955-2219(00)00021-2Journal of the European Ceramic Society20(2000)1493±1503*Corresponding author now at:TECSEN Laboratory,case231, 13397Marseille Cedex20,France.E-mail address:lancin@matop.u-3mrs.fr(ncin).(15MPa),during 1min,at 918K.13Two composites were obtained using respectively bare ®bres or ®bres coated by pyrolytic carbon (Cp),both containing a high volume fraction of ®bres (%70%).Longitudinal and cross sections were prepared by mechanical grinding and ion thinning.During grinding and polishing,the cohesion of the samples was main-tained by using an appropriate glue and a Ti ring (Bal-zers).These necessary precautions and the occurrence of cracks sometimes observed between ®bres and matrix (Fig.1a)reveal the weak adhesion between ®bres and matrix in both composites.Ion thinning is carried out in a Balzers ion thinner using a low beam density and a sample holder cooled with liquid nitrogen.The composite micro-structure is not modi®ed by such ion thinning conditions.Phase identi®cation and localisation were performed by electron di raction patterns (EDP)and nanodi raction,high resolution transmission electron microscopy (HRTEM),energy dispersive X-rays spectroscopy (EDX)and electron energy loss spectroscopy (EELS).Experi-ments were carried out using a ®eld emission gun (FEG)HF2000Hitachi microscope (200KV,Scherzer resolu-tion=2.6A)equipped with a CCD camera (MSC Gatan),a Si-Li detector (Kevex Super Quantek)and a PEELS (Gatan).Analyses were performed using a cooling holder (Gatan),a 4±10nm probe size and a 12 tilt angle for EDX spectroscopy.3.ResultsMany microstructural features are shared by the two composites.We thus describe them and underline the di erences where theyoccur.Fig.1.Microstructure of a cross-section of the M40J/A357composite.Bright ®eld (a)and HRTEM (b)images of the Al/®bre interface.(b)Shows the folded aromatic layers surrounding porosities and the EDP characteristic of equiaxed {0002}carbon planes.ncin,C.Marhic /Journal of the European Ceramic Society 20(2000)1493±15033.1.ReinforcementThe mechanical properties of the composites depend on the microstructure of the reinforcement after the processing.The microstructural features of the ®bre and of the Cp were thus investigated on longitudinal and cross-sections of the M40J/A357and M40J/Cp/A357composites respectively. 3.1.1.Microstructure of M40J carbon ®breThe ®bres are regularly distributed in the composites.Their bean shaped cross-section (Fig.1a)is fringed with roughness ranging from a few nanometers to a few microns in size (Figs.1,4a,5,6,7,8).Their longitudinal section is more regular than their cross-section but the surface is wavy,with an amplitude which reaches 10nanometers or so (Fig.2a).Owing to theirbean-shapedFig.2.Microstructure of a longitudinal section of the M40J/Cp/A357composite.Bright ®eld image of the Al/Cp/®bre interface (a)HRTEM image of the microporous Cp (b),HRTEM images and EDPs of the layered Cp (c)the ®bre (d).ncin,C.Marhic /Journal of the European Ceramic Society 20(2000)1493±15031495cross-section,their¯uted surface and their dense pack-ing,the®bres are always in contact with two or three neighbours and are bridged to the others by an irregular and narrow matrix.As a result,the matrix is cut into microlingots and a residual porosity is observed where ®bres are in contact.The®bre section exhibits a uniform contrast in bright ®eld mode(Figs.4a,6)except in its outermost part where signi®cant porosity is revealed by its light con-trast.This porosity,which is not continuously observed all around the®bre,extends over10±200nm beneath the surface.The pore size ranges from4±20nm. Using M40J/A357cross-sections,the®bre micro-structure was characterised.The same di raction pat-tern is obtained in the®bre core and in the®bre periphery whatever its porosity(Fig.1b).This con-tinuous ring pattern is characteristic of equiaxed{0002} planes.HRTEM images con®rm these results and reveal the grain size(Fig.1b).Fibre nanocrystals consist of a stack of a few aromatic layers,one or two nanometers in length,named BSU(Basic Structural Unit15).BSUs of similar orientation are associated in clusters namedLMO(Local Molecular Orientation15)whose width and length rarely exceed10nanometers.BSUs are joined by disclinations where atoms are linked by twis-ted or stretched bonds or tetrahedral bonds as typical in carbon®bres.16Such an organisation results in many places in a strong curvature of the aromatic layers,in very sharp angles between them and in a residual por-osity.The pore size rarely exceeds a few nanometers in the®bre except near the surface as above mentioned.The folded organisation of the aromatic layers is observed throughout the®bre section even in its outermost part. HRTEM images of longitudinal sections reveal the length of BSUs and LMOs organised in micro®brils17 along the®bre axis(Fig.2d).As in the cross-sections,the BSU length is equal to one or two nanometers and the LMO width perpendicular to the®bre axis is about10nm. Unlike the cross-sections,micro®brils extend over several tens of nanometers along the®bre axis and the relative disorientation of aromatic layers is less substantial.EDPs show,indeed,that aromatic planes are aligned to within 17 of the®bre axis(Fig.2d).The space between micro®brils reveals the length of the pore structures (about50nm).Because of this microstructure,the®bre is categorised into the large family of turbostratic carbons. Interdi usion between aluminium and®bre was stu-died by EDX(Fig.3).Aluminium is revealed beneath the®bre surface over about®ve nanometers but ana-lyses may be in¯uenced by the®bre roughness.Traces of aluminium and silicon are also detected in the®bre just beyong this zone(Fig.3a),particularly in the pores. In aluminium,carbon is not detected without ambiguity either by EDX or by EELS.A layer of5nanometers or so in the matrix next to the Al/C interface contains some oxygen(Fig.3c).3.1.2.Microstructure of pyrolytic carbon coating(Cp) In the M40J/Cp/A357cross-sections,the Cp coating is identi®ed at the®bre/matrix interface by a light con-trast in bright®eld mode(Figs.4a,6b)and by EDPs and HRTEM images(Fig.4b)which distinguish®bre and Cp.Every®bre is continuously surrounded by the Cp coating,15±30nm thick.The Cp,in which aromatic lay-ers are cut into BSUs,exhibits a turbostratic structure. Near the®bres,BSUs tend to form continuous sheets which follow the roughness of the®bre surface and LMOs can hardly be distinguished.The orientation of the aromatic layers varies to within 35 with respect to the®bre axis.When the coating thickness exceeds about 10nanometers,LMOs are more and more distinct and their orientation less and less equiaxed resulting in a microporous microstructure near the matrix. HRTEM images and EDPs of longitudinal sections con®rm the change of the Cp microstructure versus thickness(Fig.2b,c).Along the®bre surface,over 10nm or so,we observe a layered structure.As in the cross-sections,LMOs or micro®brils cannot be dis-tinguished on HRTEM micrographs.In the layered zone,the aromatic layers follow the®bre surface and their disorientation versus the®bre( 25 )is only a lit-tle more evident than in the®bre.In the outermost part of the Cp,the microporous microstructure is in excellent agreement with that revealed by HTREM images on cross sections.In the microporous Cp coating,traces of aluminium, silicon and magnesium are detected by EDX but no other heteroatoms and in particular no oxygen atoms are found(Fig.3b).In the layered Cp coating,hetero-atoms are notdetected.Fig.3.EDX analysis.Spectra characteristic of the®bre at20nm from the interface(a)the microporous Cp(b)Al next to the Al/C interface (c)Al at20nm from the Al/C interface(d)Mg2Si(e)and Si next to the Si/C interface(f).ncin,C.Marhic/Journal of the European Ceramic Society20(2000)1493±15033.2.Brittle phasesAl4C3reaction product and matrix precipitates(Si or Mg2Si)are identi®ed on cross-sections by electron dif-fraction,HRTEM,EDX and by plasmon peak energy. Low loss edges permit a more systematic study of the phases than core edges because they are less sensitive to specimen thickness.Moreover,despite the variation of the plasmon peak energy of carbon as a function of the beam orientation versus the basal plane,it was possible to distinguish between the di erent phases because we found distinct values for the volume plasmonenergyFig.4.Microstructure of a cross-section of the M40J/Cp/A357composite.(a)Bright®eld image of a®bre/Cp/Al interface where the Cp and the ®bre porosities are revealed by a light contrast;(b)HRTEM image showing the layered Cp near the®bre and the microporous Cp near the matrix; EDP of the layered Cp.ncin,C.Marhic/Journal of the European Ceramic Society20(2000)1493±15031497(Fig.6):E Mg2Si =12.7eV,E Al =15eV,E Si =16.7eV,E carbide =18.9eV and for the di erent carbons 23.4eV in microporous Cp,24.3eV in layered Cp,23.7eV in the ®bre and 21eV in large ®bre pores.3.2.1.Si and Mg 2Si precipitatesThe matrix contains Si grains,inhomogeneously dis-tributed in the matrix and located predominantly next to the ®bres,generally in contact with two ®bres or between them (Fig.6).All these grains exhibit typical stacking faults;thus their contrast was most often used to identify them.It is noteworthy that the amount of stacking faults is the highest in the narrowest parts of the Si grains which bridge between ®bres.The Si grain size generally ranges between a hundred nanometers to about one micron but some nanometric grains are also detected.Si grains contain various amount of inclusions (Fig.6b,c)in which aluminium is detected but whose exact composition is under study.The largest inclusions were identi®ed as aluminium.Mg 2Si precipitates a few tens nanometers in size were on occasion identi®ed by EDX in the matrix (Fig.3e).Mg 2Si precipitates were never detected between the®bres probably because of the low concentration of Mg in the alloy as compared to Si.3.2.2.Aluminium carbidesIn both composites,aluminium carbide grains exhibiting a lath like morphology were revealed at interfaces using the techniques previously mentioned.In contact with the carbon,but not extending into the ®bres,the carbide grains are included in the matrix.They are generally located at Al/C interface (Figs.1b and 7a)but some of them are included in silicon grains (Fig.7b).The carbide grain length ranges from 100to 350nm and their aspect ratio (length/width)from 8to 15.The amount of carbide grains is di cult to assess because they are irregularly observed at the interfaces.From a systematic study of longitudinal and cross-sections,it is possible to conclude that carbides are rather rare in both composites but rarer at the Cp/matrix interface than at the ®bre/matrix interface.Moreover,their amount is clearly lower than that of the Si grains.3.2.3.Si±C interfacesHRTEM was performed to study the possible reaction between carbon and silicon resulting in SiCformationFig.5.EELS analysis:volume plasmon peaks.ncin,C.Marhic /Journal of the European Ceramic Society 20(2000)1493±1503which has been reported at low temperature(798K)in Al±Zn alloy.18No grains were detected at the Si/C interfaces in either composite(Fig.8)except for the Al4C3as above mentioned.Thus,the C±Si reaction,if any,is not detectable after liquid metal in®ltration. Interdi usion is also to a very limited extent detected by EDX in a narrow zone on either side of the interface. Carbon is not detected in the silicon beneath the inter-face(Fig.3f)whereas traces of silicon and aluminium are observed in the carbon,the spectrum being similar to that obtained at the Al/C interfaces.4.DiscussionBrittle phases have been identi®ed at interfaces in both composites.Their in¯uence depends on their brit-tleness as compared to that of the reinforcement and on the resistance to debonding and to sliding at the inter-face.Although the Cp coating does not change the strength and the Young modulus of the®bre,14it may modify crack propagation in the reinforcement.The resistance to debonding and to sliding between the rein-forcement and the di erent phases,aluminium,alumi-nium carbide or silicon,will accordingly be analysed using the microstructural features of the interfaces.Finally,we show that the conclusions derived from the micro-structural features are consistent with the mechanical behaviour of the composites4.1.Fibre and Cp microstructure and crack propagationThe turbostratic structure of the M40J®bre is more typical of the high strength PAN-based®bres than of the high modulus ones®rst described by Guigon et al.19 The core porosity for example,is close to that observed in ex-Pan high strength T300®bres by Lamouroux et al.20Moreover,in agreement with Feldho et al.,21weFig.6.Silicon grains.Localisation of Si grains in a cross-section thanks to their stacking faults(a).Bright®eld images of Si grains bridging®bres in M40J/Cp/A357(b)and in M40J/A357(c)composites:images reveal stacking faults and Al inclusions in Si and the continuous ribbon of Cp between Si and®bre in(b).ncin,C.Marhic/Journal of the European Ceramic Society20(2000)1493±15031499do not observe the layered structure of the outermost part of the®bres mentioned by Guigon et al.in high modulus®bres.The lack of layered structure could be due to some dissolution of the®bre during composite fabrication.However,in any case,the thickness of the possible lamellar layer must be limited in the M40J®bre because the short duration of the LMI and the small amount of aluminium carbide formed are not consistent with extensive dissolution.Fibre and Cp have di erent turbostratic structures but both of which could avoid crack propagation through the reinforcement.However,the impact of a possible surface notch is di erent whether it is applied on the®bre or on the Cp.The e ciency of pyrolytic carbon as a mechanical fuse is generally ascribed to a layered structure which promotes crack de¯ection and interface sliding.As for the crack de¯ection,the complex structure of the Cp is particularly e cient.The outermost part of the Cp exhibits indeed a microporous morphology where the LMOs are equiaxed.Such a structure is very suitable to promote crack branching on the disclinations as well as crack de¯ection in LMOs in a direction parallel to the {0002}planes and,thus,it is e cient to release stresses at crack tips.If blunted cracks could nevertheless reach the inner part of the Cp,where aromatic layers follow the®bre surface,they could be de¯ected in a direction parallel to the®bre surface more e ciently than the initial crack.As for the interface sliding,we have to consider the adhesion of the di erent phases and we discuss the in¯uence of this parameter later.When submitted to similar stresses than the Cp,the turbostratic structure of the®bre may also promote crack splitting and crack de¯ection in a direction parallel to its axis because folded aromatic layers are oriented to the®bre axis within 17 .However,during atensileFig.7.Identi®cation of A14C3grains.HRTEM images of A14C3grains at an Al/®bre interface(a)and included in a silicon grain identi®ed by its characteristic stacking faults and Al inclusions(b).In both A14C3grains the contrast of{0003}planes is visible.ncin,C.Marhic/Journal of the European Ceramic Society20(2000)1493±1503test,because the®bres are seldom aligned parallel to the applied stress,the notch tip will usually be submitted to shear stresses(failure in mode III22)which will promote propagation through the®bre and thus®bre failure. 4.2.AdhesionAdhesion of aluminium matrix to®bres(bare or coated)is determined by con¯icting e ects.During LMI,carbon is dissolved in liquid aluminium and alu-minium di uses into the carbon.After quenching,car-bon atoms remain in the matrix,even if they are undetectable by TEM,and Al is still observed in the carbon.Therefore,chemical di usion bonds which are said to be``good''bonds4can be considered to link aluminium to®bres.However,we have seen that the matrix is cut into micro-lingots by®bres in contact. During cooling,the shrinkage is determined by the ®bres.The coe cient of thermal expansion of alumi-nium(2310À6KÀ1)being higher than the axial(9À1 10À6KÀ1)and transverse(%À1010À6KÀ1)coe cients of thermal expansion of a®bre,Al/C interfaces are submitted to radial tension and transverse shear stress which counterbalance the chemical di usion bond.The aluminium matrix is,however,®xed here and there to the®bres by aluminium carbide grains.Aluminium carbide grains that are formed by reaction between liquid aluminium and carbon are linked to the reinforcement by a strong chemical reaction bond. However,carbide grains are also strongly linked to the matrix by mechanical bonds because they are embedded during the solidi®cation in aluminium that has a higher coe cient of thermal expansion than aluminium car-bide.Thus,Al4C3/C interfaces as well as Al/C interfaces are submitted to radial tension in opposition to the carbide±carbon reaction bonds.As for silicon-reinforcement bonding,chemical di u-sion bonds may exist because some silicon has di used into the®bres during LMI.However,we think that such bonds are of little importance compared to the mechanical bonds that link silicon to®bre if considera-tion is given to resistance to debonding and to sliding. The coe cient of thermal expansion of silicon(%2.92 10À6KÀ1)being low compared to the transverse coe -cient of thermal expansion of the®bre,silicon grains are ®rmly®xed on the®bre by the interlocking e ect pro-moted by the roughness.Such e ects are e cient enough to grip together phases as di erent as porous alumina and pivalic acid and to resist shear stresses during sample cutting.23Moreover,silicon grains are submitted to compression stresses when they bridge ®bres.Interlocking e ects and compression stresses limit,even if they do not prevent,sliding at the silicon/®bre interface when silicon grains bridge®bres.In silicon grains the gliding of Shockley partial dis-locations in{111}planes which causes stacking faults more numerous in the narrowest part of the silicon bridges,releases some thermal stresses during quench-ing.Residual thermal stresses may,however,decrease the strength to failure of the®bre near these bridges. The transverse stress to failure of the composites which results from these di erent bonds is weak as shown by the di cult thin®lm preparation(Section2and Fig.1a). Because Al4C3/C or Si/C interfaces are rather rare,this low adhesion is not exclusive of stronger bonding at a few points at the interfaces.4.3.In¯uence of the brittle phases on the mechanical behaviourIt is best®rst to discuss the in¯uence of silicon grains because they are more numerous than aluminiumcarbide Fig.8.Si/C.HRTEM images of the Si/®bre interface which shows that no Si±C reaction occurs during the composite fabrication.ncin,C.Marhic/Journal of the European Ceramic Society20(2000)1493±15031501grains and because they often bridge®bres.When®bres are bridged by a silicon grain,the composite has locally a brittle matrix.Because of the high volume fraction of ®bres and because neighbouring®bres are generally in contact,it is reasonable to suppose that the Young modulus of the composites(Ec)is determined by the volume fraction of®bres(Ec=0.7Ef=263GPa).Silicon has the lowest Young modulus(Esi=163GPa)when compared to®bre and composite.During a tensile test, the composite strain is higher than that of a®bre sub-mitted to the same stress and lower than that of silicon. The®bre/Si interface is thus submitted to shear stresses, the®bre being under greater tension than the silicon.These stresses are in opposite direction to the thermal stresses developed during cooling because of thermal mismatch. It is not clear if silicon grains are more or less brittle than®bres.The tensile strength to failure found by Pearson et al.(2GPa24)results in a strain to failure (12.210À3)higher than that of®bres(11.6710À3). However,the strength to failure of brittle materials depends on¯aw size.According to experimental results,14silicon grains seem to be more brittle than ®bres and,thus,we analyse the impact of a silicon crack on the mechanical behaviour in this light.It is helpful®rst to consider silicon grains linked to coated®bres in the M40J/Cp/A357composite.Silicon grains have the critical size to initiate a fracture in the reinforcement25if no mechanism releases the stresses at the crack tip.When silicon is linked to the Cp,stresses will be released by decohesion either at the silicon/Cp interface or in the Cp(Section4.1)depending on the debonding resistance as compared to the shear resis-tance of{0002}planes.After interface failure,a defor-mation of the outermost part of the coating under interfacial shear stresses can occur,despite radial com-pression stresses and surface roughness,because of its microporous microstructure and because®bres are under tension at the interface(see above).Microcracks and local shearing would release the stresses and may allow some sliding at the interface.These mechanisms delay the®bre and the composite failure.However, because these deformations are certainly limited and because the residual thermal stresses decrease the®bre strength,the composite breaks under a stress sub-stantially lower than the value predicted by the law of mixtures(1300instead of3000MPa)and with some pull out consistent with these model.When a silicon grain bridges bare®bres,even if the resistance to debonding at the Si/®bre interface were low,sliding is restricted if not impossible at the interface without notching the®bre.Therefore,whatever the resistance to decohesion,stresses cannot be released at the crack tip and they will induce®bre failure.The simultaneous failure of silicon grains produces the simultaneous failure of many®bres and thus a brittle failure of the composite under low stress.Owing to its well established deleterious e ect,alumi-nium carbide is certainly more brittle than®bres. Moreover,despite their limited size,previous results2À4 suggest that carbide grains can develop cracks which have the critical size to induce a®bre failure.When a crack notches the reinforcement,because of the turbos-tratic structure of the carbon and because of the axial tension at the Al4C3/C interface,it will be soon de¯ected in a direction parallel to the interface.In the case of coated®bres,the crack is de¯ected in the Cp.Sliding along the interfacial failure is possible despite the®bre roughness because of the limited width of the carbide grains(Section3.2.2)and again because of the axial tension at the Al4C3/C interface.Thus,a carbide failure does not induce a®bre failure and the carbide grains have no in¯uence on the strength of the M40J/Cp/A357 composite.In the case of a bare®bre,the notch will initiate a®bre failure when the tensile stress is high enough(Section4.1).Aluminium carbide grains con-tribute to the weakness and brittleness of the M40J/A357 composite,but,due their low amount as compared to silicon grains,do not provide the main in¯uence on M40J/A357strength.5.ConclusionThe main conclusions are as follows:.the microstructure of the high modulus PAN based M40J carbon®bre is fully characterised;it is in better agreement with the microstructure of a high strength than of a high modulus PAN based carbon®bre,.the pyrolytic carbon coating(Cp)used in the M40J/Cp/A357composite consists of two layers of distinct microstructure,a layered structure near the®bre and a microporous structure near the matrix,.the reaction of aluminium with carbon is limited during fabrication processing,particularly with the Cp;no reaction occurs between silicon and Cp or ®bre,.brittle phases,silicon and to a lesser extent alumi-nium carbide,are in contact with or bridge the ®bres,.the response of the composites to tensile stress parallel to the®bre axis is well described by the model of a brittle matrix reinforced by brittle ®bres:the most important parameter is the high resistance to sliding due to interlocking e ects between silicon grains and®bres and compression stresses in silicon bridges;it leads to a brittle and weak M40J/A357composite but its in¯uence is counterbalanced by the special structure of the Cp interphase in the M40J/Cp/A357composite.ncin,C.Marhic/Journal of the European Ceramic Society20(2000)1493±1503。

包含纳米CoSb3的Yb0.15Co4Sb12基复合材料的合成和热电性能1糜建立，赵新兵，朱铁军浙江大学材料科学与工程系，硅材料国家重点实验室，杭州 (310027)E-mail: zhaoxb@摘要：在块体材料中引入纳米组元构建微纳复合材料是热电研究的一个新方向。

本文合成了包含纳米CoSb3的Yb0.15Co4Sb12基复合材料，系统研究了不同含量的纳米CoSb3对Yb0.15Co4Sb12材料热电性能的影响。

将溶剂热合成的CoSb3粉末和熔炼/退火方法合成的Yb0.15Co4Sb12粉末进行放电等离子体烧结，制备了包含纳米CoSb3晶粒的Yb0.15Co4Sb12基微纳复合材料。

复合材料由小于100 nm的细小晶粒和微米级别的粗大晶粒组成。

少量纳米晶粒的掺入能够提高材料的Seebeck系数。

虽然掺入的纳米晶粒为未填充的CoSb3，但是它们的掺入仍然有效地降低了材料的热导率。

微纳复合结构引入了大量晶界，增强了声子散射，从而降低材料的热导率。

复合材料的最高ZT值达到0.9，与基体材料Yb0.15Co4Sb12相比增加了10%。

关键词：溶剂热合成，微纳复合，方钴矿，热电性能中图分类号：O 482, TN 371. 引言热电材料是一种能够实现电能与热能之间直接相互转换的功能材料。

热电材料的性能常　/　，其中　，　和　分别是Seebeck系数，电导率和用无量纲优值ZT来衡量，ZT = 2T热导率，T为绝对温度。

方钴矿化合物作为一种新型的热电材料，具有很好的应用前景[1, 2]。

在方钴矿结构中填入其它原子形成填充方钴矿材料能够有效降低材料的热导率。

Yb作为方钴矿化合物的填充元素已有广泛的研究。

其中Yb部分填充的CoSb3化合物被报道具有较好的热电性能[3]。

然而对于大多数填充方钴矿来说，其热导率与Bi2Te3基热电材料相比仍然较高。

纳米化能够有效降低材料的热导率，从而提高材料的热电性能[4]。

微纳复合是热电材料研究的一个新方向，对提高材料热电性能也具有很好的效果[5]。

1. Acta Biomaterialia2. ACS Applied Materials & Interfaces3. Advanced Engineering Materials4. Advanced Functional Materials5. American Mineralogist6. Biomacromolecules7. Biomaterials8. Biomedical Materials9. Biomedical Microdevices10. Biomedical Optics Express11. Biopolymers12. Biotechnology and Bioengineering13. Clinical Biomechanics16. Drug Development and Industrial Pharmacy17. e-Polymers18. Green Materials19. IEEE Transactions on Biomedical Engineering20. International Journal of Artificial Organs21. International Journal of Biological Macromolecules22. International Journal of Pharmaceutics23. Journal of Biomedical Materials Research Part A24. Journal of Biomedical Materials Research Part B25. Journal of Materials Science: Materials in Medicine26. Journal of Mechanical Behavior of Biomedical Materials27. Journal of Molecular Structure28. Journal of Nuclear Materials29. Macromolecular Bioscience31. Materials Science & Engineering C: Materials for Biological Applications32. Polymer International生物材料，也称为生物学材料，是一种可以用于医学、生物工程、机器人、非机械医学和生物信息学等科学领域的材料。

2024年的材料类SCI影响因子涵盖了多个不同类型的材料研究领域，包括金属、陶瓷、高分子、晶体、玻璃、复合材料等。

以下是2024年一些重要材料类SCI期刊的影响因子:1. Nature Materials - 影响因子: 36.503Nature Materials是一个综合性的材料科学期刊，涵盖了广泛的材料研究领域，包括材料设计、合成、结构、性能和应用。

它发表了一些具有重要影响力的研究论文和综述，被认为是材料科学领域的顶级期刊。

2. Advanced Materials - 影响因子: 15.409Advanced Materials是一个跨学科的材料科学期刊，涵盖了从基础研究到应用研究的各个方面，包括材料设计、合成、功能性能和应用。

它发表了一些具有创新性和高水平的研究论文，并致力于推动材料科学的发展。

3. Journal of Materials Chemistry - 影响因子: 6.626Journal of Materials Chemistry是一个涵盖了多个材料类别的综合性期刊，包括无机材料、有机材料和生物材料。

它发表了许多具有重要影响力和创新性的研究论文，并促进了材料科学的交叉学科研究。

4. Polymer - 影响因子: 3.766Polymer是一个专注于高分子材料研究的期刊，发表了很多关于高分子合成、结构、性能和应用的研究论文。

它在高分子领域具有很高的影响力，并且为高分子材料的发展做出了重要贡献。

5. Journal of Crystal Growth - 影响因子: 1.711Journal of Crystal Growth是一个专注于晶体生长和晶体学的期刊，发表了很多关于晶体生长机制、晶体结构和晶体生长技术的研究论文。

它在晶体领域具有一定的影响力，并且为晶体技术的发展做出了贡献。

6. Journal of the American Ceramic Society - 影响因子: 2.841Journal of the American Ceramic Society是一个专注于陶瓷材料研究的期刊，发表了许多关于陶瓷材料的合成、结构、性能和应用的研究论文。