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表面技术第53卷第5期H13钢表面激光熔覆NbC/Ni60复合涂层组织及高温耐磨损性能常倾城1，任利兵1，刘英1,2*，谢咏馨1，李卫1（1.暨南大学 先进耐磨蚀及功能材料研究院，广州 510632；2.暨南大学 韶关研究院，广东 韶关 512027）摘要：目的研究NbC颗粒的加入量对H13钢表面激光熔覆NbC/Ni60复合涂层的组织、硬度和耐磨性的影响。
方法将Ni60合金粉末与NbC碳化物粉末球磨混合，采用激光熔覆技术，在H13钢基体表面制备不同NbC含量（质量分数分别为0%、10%、20%、30%）增强的NbC/Ni60合金复合涂层。
采用电子扫描显微镜（SEM）、X射线衍射仪对复合涂层的微观组织和物相进行分析。
借助显微硬度计，研究复合涂层的截面显微硬度分布规律。
采用高温摩擦磨损试验机测试复合涂层在真空400 ℃下的摩擦磨损性能。
结果在激光熔覆NbC/Ni60复合涂层中，物相主要由γ-(Ni, Fe)固溶体、Ni2Si、CrB、Cr23C6、NbC组成；熔覆层以胞晶和枝晶为主，NbC含量对复合熔覆层组织及形态具有显著影响，加入少量NbC可使熔覆层组织细化；在NbC 的质量分数为20%时，大量弥散的NbC颗粒在枝晶间呈聚集趋势；在NbC的质量分数为30%时，熔覆层中NbC相呈现块状、花瓣状形貌。
NbC/Ni60复合涂层的硬度显著高于H13钢基体，随着NbC含量的增加，NbC/Ni60复合熔覆层的显微硬度逐渐升高，NbC的质量分数为30%的NbC/Ni60复合熔覆层的平均硬度高达848HV。
在真空400 ℃、压力100 N、转速100 r/min、时间7 200 s磨损工况下，NbC质量分数为20%的NbC/Ni60复合涂层的磨损量最小，因此其高温耐磨性最好。
NbC质量分数为10%的NbC/Ni60复合涂层的摩擦因数最小。
随着NbC含量的增加，复合涂层的摩擦因数反而升高。
结论NbC/Ni60复合涂层与H13钢基体具有很好的冶金结合，显著提高了高温耐磨性能；NbC颗粒硬质相具有较好的增强作用，能够明显提高NbC/Ni60复合涂层的硬度和耐磨性；粗大NbC相不利于复合涂层耐磨性的进一步提高。
第17卷 第9期强激光与粒子束Vol.17,No.9 2005年9月H IGH POWER L ASER AND PA R TICL E B EAMS Sep.,2005 文章编号: 100124322(2005)0921409205径向非均匀磁场下的磁控管工作性能模拟3陈 彦1, 杨中海1, 雷文强2(1.电子科技大学物理电子学院高能电子学研究所,四川成都610054;2.中国工程物理研究院应用电子学研究所,四川绵阳621900) 摘 要: 为了改善磁控管的输出频谱,使模拟结构更接近实际情况,考虑了磁钢和极靴的实际尺寸。
用MA FIA 对磁控管内的磁系统进行建模和模拟。
得到的磁场纵向分布有随着径向半径的增大而增大的趋势,这种径向不均匀性与实验测试结果一致。
将非均匀的径向阶梯形变化磁场分布带入腔体热测计算,模拟得到的π模工作频率2.437GHz ,实际工作频率2.450GHz ,相对偏差0.5%,并在高频谐波的抑制上获得了输出频谱的显著改善,使微波炉磁控管具有更好的电磁防护和实际应用。
关键词: 磁控管; 径向非均匀性磁场; 腔体热测计算; π模频率; 电磁防护 中图分类号: TN124 文献标识码: A 磁控管作为一类正交场器件在雷达技术和微波加热领域具有广泛的需求。
磁控管具有金属翼片型阳极块和阴极,为了获得正交型的均匀磁场还需要磁钢和极靴,腔体电磁场与外加磁场的相互作用使得电子回转成轮辐,最终形成振荡微波源。
由于磁控管结构的复杂性,国内外学者们大多采用仿真方法得到有规律的研究结果,从早期的二维六翼片腔体模拟[1]和三维旭日型腔体模拟[2],发展到十翼片双隔膜带的腔体三维模拟技术[3～5],计算结果越来越接近测试值,提高了模拟计算的精度。
但大多数的模拟方法中都没有考虑磁钢和极靴的实际结构,只是假定一个纵向均匀的磁场分布,这与实际情况并不相符。
近年来由于无线通讯大多占据2.4GHz 的频带,这与微波炉磁控管工作频率2.45GHz 非常接近,因此需要对磁控管的频谱噪声进行抑制,这已成为国际上的研究热点。
第15卷第6期精密成形工程袁美霞，柳校可，华明（北京建筑大学机电与车辆工程学院，北京 100044）摘要：选区激光熔化技术（SLM）被认为是极有前途的增材制造技术之一，但不可逆的溅射行为严重限制了SLM技术的应用。
从粉末熔池演变、加工工艺优化和飞溅颗粒动态特征监测等方面，总结了SLM过程中飞溅行为的研究现状，分析了飞溅行为的产生机制，探讨了激光–粉末–熔池相互作用下的熔池演变情况，表明金属蒸气、Marangoni效应和伯努利效应是诱发飞溅的主要因素；讨论了加工工艺与飞溅行为的相互关系，表明通过优化工艺参数和改善打印环境以抑制飞溅是行之有效的方法；阐述了飞溅诱导缺陷的机理，并讨论了SLM过程的监测方法，表明单一信号的局限性会导致监测结果失准，多信号融合监测是提升精准性的重要方法之一。
最后，针对飞溅行为存在的关键科学问题和技术难题，展望了SLM加工中飞溅行为的研究方向。
关键词：SLM技术；增材制造；飞溅机制；内部缺陷；飞溅监测DOI：10.3969/j.issn.1674-6457.2023.06.020中图分类号：TF121 文献标识码：A 文章编号：1674-6457(2023)06-0163-11Research Progress of Splashing Behavior in Selective Laser MeltingYUAN Mei-xia, LIU Xiao-ke, HUA Ming(School of Mechanical-electronic and Vehicle Engineering, Beijing University of Civil Engineering and Architecture,Beijing 100044, China)ABSTRACT: Selective Laser Melting (SLM) is considered one of the most promising additive manufacturing (AM) technolo-gies, but the irreversible sputtering behavior severely limits the application of SLM. The work summarized the research status of splashing behavior in SLM from the aspects of powder melt pool evolution, processing process optimization, and dynamic monitoring of splashed particles. The mechanism of splashing behavior was analyzed, and the evolution of melt pool under the interaction of the laser powder molten pool was explored. It showed that metal vapor, Marangoni effect, and Bernoulli effect were the main factors inducing splashing. The interaction between processing technology and splashing behavior was discussed, indicating that optimizing process parameters and improving the printing environment were effective methods to suppress splashing. The mechanism of splash-induced defects was elaborated, and the monitoring methods of SLM process were dis-cussed. It showed that the limitations of a single signal could lead to inaccurate detection results. Multi signal fusion monitoring was an important method to improve accuracy. Finally, in response to the key scientific and technical challenges in studying收稿日期：2023–03–07Received：2023-03-07基金项目：国家重点研发项目（2022YFC2406004）；北京建筑大学研究生创新项目（PG2022134）Fund：National Key Research and Development Projects(2022YFC2406004); Beijing University of Architecture Graduate In-novation Project(PG2022134)作者简介：袁美霞（1979—），女，博士，副教授，主要研究方向为增材制造。
表面技术第53卷第5期高温对含氢DLC涂层的微观结构及力学性能的影响贾伟飞1，梁灿棉2，胡锋1,2*（1.武汉科技大学 高性能钢铁材料及其应用省部共建协同创新中心，武汉 430081；2.广东星联精密机械有限公司，广东 佛山 528251）摘要：目的针对含氢DLC涂层热稳定性很差的问题，探究高温下含氢DLC涂层的微观组织变化特征，以及高温对其力学性能的影响。
方法采用等离子体强化化学气相沉积（Plasma Enhanced Chemical Vapor Deposition, PECVD）在S136模具不锈钢表面沉积以Si为过渡层的含氢DLC复合涂层，利用光学显微镜、扫描电镜、拉曼光谱、X射线电子衍射仪、三维轮廓仪研究DLC涂层的微观结构，采用划痕测试仪、往复式摩擦磨损试验机、纳米压痕仪研究DLC涂层的力学性能，并通过LAMMPS软件，利用液相淬火法建立含氢DLC模型，模拟分析经高温处理后涂层的组织变化特征和纳米压痕行为。
结果在400 ℃、2 h的退火条件下，拉曼谱峰强度I D/I G由未退火的0.7增至1.5，涂层发生了石墨化转变，同时基线斜率下降，H元素析出；XPS结果表明，在此条件下涂层中sp2杂化组织相对增加，氧元素增多，涂层粗糙度增大；在600 ℃、2 h退火条件下，DLC发生了严重氧化，LAMMPS模拟结果表明，在400 ℃高温下涂层的分子键长变短，表明sp3杂化组织在高温下吸收能量，并向sp2杂化转变。
纳米压痕模拟结果显示，在400 ℃下退火后，涂层的硬度下降。
结论在400 ℃下退火处理后，涂层中的H元素释放，涂层内应力减小，保证了涂层的强度；在600 ℃退火条件下，过渡层的Si和DLC在高温下形成了C—Si键，使得DLC薄膜部分被保留；LAMMPS 模拟结果表明，在高温下涂层发生了石墨化转变，涂层的硬度减小。
关键词：含氢DLC涂层；退火处理；微观组织；力学性能；LAMMPS模拟中图分类号：TB332 文献标志码：A 文章编号：1001-3660(2024)05-0174-10DOI：10.16490/ki.issn.1001-3660.2024.05.018Effect of High-temperature on Microstructure and MechanicalProperties of Hydrogen-containing DLC CoatingJIA Weifei1, LIANG Canmian2, HU Feng1,2*(1. Collaborative Innovation Center for Advanced Steels, Wuhan University of Science and Technology, Wuhan 430081,China; 2. Guangdong Xinglian Precision Machinery Co., Ltd., Guangdong Foshan 528251, China)ABSTRACT: The thermal stability of hydrogen-containing DLC coating is poor, and the work aims to explore the microstructure changes of hydrogen-containing DLC coating at high temperature and their impact on mechanical properties. The收稿日期：2023-01-09；修订日期：2023-05-18Received：2023-01-09；Revised：2023-05-18基金项目：中国博士后科学基金（2021M700875）Fund：China Postdoctoral Science Foundation (2021M700875)引文格式：贾伟飞, 梁灿棉, 胡锋. 高温对含氢DLC涂层的微观结构及力学性能的影响[J]. 表面技术, 2024, 53(5): 174-183.JIA Weifei, LIANG Canmian, HU Feng. Effect of High-temperature on Microstructure and Mechanical Properties of Hydrogen-containing DLC Coating[J]. Surface Technology, 2024, 53(5): 174-183.*通信作者（Corresponding author）第53卷第5期贾伟飞，等：高温对含氢DLC涂层的微观结构及力学性能的影响·175·hydrogen-containing DLC composite coating with Si as the transitional layer was deposited on the surface of S136 stainless steel by plasma enhanced chemical vapor deposition (PECVD). The microstructure of DLC coating was investigated by optical/scanning electron microscopy, Raman spectroscopy, XPS (X-ray photoelectron spectroscopy) and three-dimensional profiler, the mechanical properties of DLC coating were studied by scratch, reciprocating friction wear and nano-indentation experiment, and the nano-indentation experiment behavior of DLC coating was simulated by LAMMPS to analyze the microstructure characteristics in annealing. The coating was subject to annealing conditions of 400 ℃for 2 hours and 600 ℃for 2 hours. Under the former condition, Raman spectroscopy showed an increase in the intensity ratio of the I D/I G peaks from0.7 to 1.5, indicating graphitization transition, accompanied by a decrease in baseline slope and H element segregation. XPSanalysis revealed an increase in sp2 hybridization and oxygen content in the coating under this condition, as well as an increase in surface roughness. At 600 ℃, severe oxidation of the DLC coating was observed. Under that condition, the matrix stainless steel was also oxidized. Molecular dynamics simulations using LAMMPS suggested a decrease in molecular bond length at 400 ℃high temperature. The three-dimensional profile test showed that the roughness under the unannealed condition was mainly from the large particles produced during deposition. At 400 for 2℃h, the coating had the minimum surface roughness. At this time, some large particles in the coating structure fell off, and the coating was basically completely damaged at 600 for℃ 2 h. The roughness was mainly from the original stainless steel roughness. The scratch test showed that under the condition of 400 for℃2 h, due to the release of the internal stress of the coating and the tighter bonding of the transition layer, the coating had the bestbonding effect with the substrate and was the least likely to fall off. The statistical results of LAMMPS simulation showed that the chemical bonds of the original DLC model tended to become shorter after annealing at high temperature. Relative to the unannealed DLC coating, the mechanical properties of DLC coating were best under 400 for℃ 2 h. Under this condition, the precipitation of mixed H elements in the coating led to the transformation of the original C—H sp3 structure, which occupied a large space to the smaller C—C sp3 and C—C sp2 structure, releasing internal stress in the coating, while ensuring the strength.The nano-indentation experiments showed that the elastic recovery and hardness of the coating were the highest at 400 for℃ 2 h, compared with that at other annealing temperature. The structure of the DLC coating containing hydrogen changed due to the precipitation of H element at 400 ℃. On the one hand, the coating structure changed from sp3 to sp2 due to high temperature, and on the other hand, the precipitation of H element changed the original C—H sp3 to C—C sp3, reducing the internal stress of the coating and improving the mechanical properties. The coating is basically damaged at 600 for 2 h, but the substrate still℃retains part of the coating. This is because the transition layer Si reacts with the coating to improve the heat resistance of the remaining coating. Molecular dynamics simulations using LAMMPS showed that the coating undergoes a graphitization transition at high temperature, leading to a reduction in its hardness.KEY WORDS: hydrogen-containing DLC coating; annealing treatment; microstructure; mechanical properties; LAMMPS simulationDLC（Diamond-Like Carbon，类金刚石碳，简称DLC）涂层材料具有超高硬度、低摩擦因数、优良化学稳定性等特点，广泛应用于机械、电子、生物医学等领域[1-3]。
甲胺铅碘钙钛矿物性及制备过程的分子模拟陈超;赵伶玲;王镜凡【摘要】应用分子动力学方法分析了甲胺铅碘晶体的结构特征与机械性质等相关物性,模拟了用蒸气沉积法在TiO2基底上制备甲胺铅碘晶体的过程,探讨了生成的PbI4 2?、 PbI5 3?和PbI6 4?多面体的排布方式,结合周围CH3NH3+的分布筛选出满足结构要求的初生晶核,分析了前驱盐配比对甲胺铅碘初生晶核产量的影响.结果表明,在拉伸过程中,甲胺铅碘晶体经历弹性形变、塑性形变以及断裂三个阶段,拟合计算得到的弹性模量与实验值符合较好;大部分初生晶核以PbI3?5金字塔的结构存在.前驱盐配比对各系统中 PbIx多面体的总含量影响较小,但对其中排布有效的PbIx结构以及初生晶核的产量影响较大,二者产量随着配比PbI2∶CH3NH3I的增加而迅速减小,这一关系与研究者发现的实验现象相符.%Molecular dynamics simulation is used to investigate the structure characteristics and mechanical properties, and to discuss the vapor deposition of MAPbI3on the TiO2substrate under a temperature of 300 K and three precursor compositions of PbI2∶CH3NH3I=1∶2, 1∶1 and 2∶1, respectively. During the preparation processes, three polyhedral groups including PbI4 2? tetrahedra, PbI5 3? pyramids and PbI6 4?octahedra are produced. After classifying their arrangements and analyzing the distribution of CH3NH3+ cations, early CH3NH3PbI3nuclei consisting of well-connected PbIx(x=4, 5 or 6) polyhedral clusters and sufficient amounts of surroundingCH3NH3+cations were identified. The influence on early nuclei from the precursor compositions of PbI2:CH3NH3I were discussed. The results show that the calculated values of elastic modulus is in good agreement with theexperimental results. The PbI3?5pyramids dominate over other polyhedral groups during the vapor deposition simulations. Meanwhile, even though the total amounts of polyhedra have a small dependence on the precursor compositions, the populations of the well-connected clusters and the early nuclei decrease rapidly with increasing the PbI2∶CH3NH3I ratio. This is in consistent with the experimental finding which to some degree, adding more CH3NH3I will optimize the device performance.【期刊名称】《化工学报》【年(卷),期】2018(069)006【总页数】9页(P2380-2387,封2)【关键词】杂化钙钛矿;分子模拟;结晶;前驱盐配比;太阳能;蒸气沉积法【作者】陈超;赵伶玲;王镜凡【作者单位】东南大学能源与环境学院,能源热转换及其过程测控教育部重点实验室,江苏南京 210096;东南大学能源与环境学院,能源热转换及其过程测控教育部重点实验室,江苏南京 210096;东南大学能源与环境学院,能源热转换及其过程测控教育部重点实验室,江苏南京 210096【正文语种】中文【中图分类】TQ021.9引言杂化钙钛矿太阳能电池因具有制备成本较低和能量转换效率（power conversionefficiency，PCE）较高等优点得到广泛关注[1-5]。
第19卷第1期装备环境工程2022年1月EQUIPMENT ENVIRONMENTAL ENGINEERING·1·核电材料辐照损伤的多尺度高通量计算模拟薛飞1，王忆2,1，刘向兵1，赖文生2，季骅1,2，刘剑波2，柳百新2（1.苏州热工研究院有限公司，江苏 苏州 215004；2.清华大学，北京 100084）摘要：反应堆压力容器用钢等核电材料在持续服役中，由于中子辐照造成其内部缺陷不断累积，致使材料组织结构损伤、性能劣化，对核电安全运行形成潜在威胁。
多尺度计算模拟是探索辐照缺陷演化机理的有效手段，结合等效缺陷结构理论，有望实现核电材料服役行为的高效评价与预测。
文中综述了多尺度计算模拟在核电材料辐照缺陷演化相关研究领域的进展，并对缺陷结构的多尺度演化本质及相应的多尺度高通量计算模拟方法进行了分析讨论。
结果表明，通过缺陷结构特征能量等效传递的方法可以实现从第一性原理计算到缺陷扩散反应动力学等高通量计算模拟的跨尺度耦合；通过多尺度高通量计算模拟得到的缺陷演化热力学和动力学数据，可以搭建用于预测核电材料长期服役行为的材料基因工程数据库；在材料缺陷结构特征能量-组织结构-性能关联性探讨基础上，应用高通量计算模拟，辅以高通量实验数据验证，有望建立基于材料基因组结构能的服役安全工程模型。
关键词：核电材料；辐照缺陷；多尺度模拟；高通量计算；材料基因组结构能中图分类号：TG111 文献标识码：A 文章编号：1672-9242(2022)01-0001-10DOI：10.7643/ issn.1672-9242.2022.01.001. All Rights Reserved.Radiation Damage of Nuclear Power Materials: A Review of theMulti-Scale High-Throughput SimulationsXUE Fei1, WANG Yi2,1, LIU Xiang-bing1, LAI Wen-sheng2, JI Hua1,2, LIU Jian-bo2, LIU Bai-xin2(1.Suzhou Nuclear Power Research Institute, Suzhou 215004, China; 2.Tsinghua University, Beijing 100084, China)ABSTRACT: The nuclear power materials are subjected to chronic neutron irradiation, during which radiation defects accumu-late to degrade the material structure and properties, leading to potential threat of safety of nuclear power plants. The frameworkof multi-scale high-throughput simulations is a keystone on revealing the mechanisms of radiation defect evolution, which mayfulfill the life and performance prediction based on the concept of equivalent defect structures. In this paper, the recent devel-opment of multi-scale high-throughput simulations on the defect evolution in nuclear power materials is reviewed. First, themulti-scale nature of the evolution of defect structures is introduced. Then, the state-of-the-art multi-scale simulation techniques收稿日期：2021-05-25；修订日期：2021-07-20Received：2021-05-25；Revised：2021-07-20基金项目：国家重点研发计划资助项目（2017YFB0702200）Fund：Supported by the National Key Research and Development Program of China (2017YFB0702200)作者简介：薛飞（1975—），男，博士，研究员级高工，主要研究方向为核电站老化与寿命管理技术。
Journal of Materials Processing Technology 211(2011)750–758Contents lists available at ScienceDirectJournal of Materials ProcessingTechnologyj 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 ecPerformance of a cutting tool made of steel matrix surface nano-composite produced by in situ laser melt injection technologyO.Verezub a ,Z.Kálazi b ,A.Sytcheva c ,L.Kuzsella d ,G.Buza b ,N.V.Verezub e ,A.Fedorov f ,G.Kaptay g ,h ,∗aUni.Miskolc,Dep.Production Eng.,Egyetemvaros,3515Miskolc,HungarybBAY-ATI (RI for Materials Science and Engineering),Fehervari ut 130,Budapest,Hungary cBAY-NANO (RI for Nanotechnology),Dep.Nano-metrology,3515Miskolc,Egyetemvaros,E/7,Hungary dUni.Miskolc,Dep.Polimer Eng.,Egyetemvaros,3515Miskolc,Hungary eNational Technical University (“Kharkov Polytechnical Institute”),Frunze st.21,Dep.Integrated Technologies in Machine Building,Kharkov,Ukraine fARC -The Australian Reinforcing Company,518Ballarat Road,Sunshine,3020VIC,Melbourne,Australia gBAY-NANO (RI for Nanotechnology),Dep.Nano-composites and Uni.Miskolc,Egyetemvaros,E/7,3515Miskolc,Hungary hUniversity of Miskolc,Dep.Nanotechnology,Egyetemvaros,E/7,3515Miskolc,Hungarya r t i c l e i n f o Article history:Received 15May 2010Received in revised form 4December 2010Accepted 7December 2010Available online 15December 2010Keywords:Cutting toolMetal matrix nano-composite Laser processing Tool-lifea b s t r a c tSteel-matrix (105WCr6steel)surface nano-composites with (Ti,W)C micron-sized and (Fe,W)6C nano-sized carbide precipitates were produced by in situ laser melt injection technology with subsequent heat treatment.The microhardness of a 1mm thick nano-composite layer was found to be higher than that of the initial matrix.The machinability of the surface nano-composite by a cubic boron nitride (CBN)wheel was found lower,but still reasonable compared to the initial matrix.Cutting tools produced from our new nano-composite by the CBN wheel were found to have higher wear resistance,longer tool life and provided lower cutting forces against a C45steel workpiece compared to the initial matrix of the nano-composite.©2010Elsevier B.V.All rights reserved.1.IntroductionMaterial removal and machining processes play a key role in generating value-added activities to materials and machine parts since their introduction about 3centuries ago (Shaw,2005).Nev-ertheless,there is a constant interest in this field due to the high variety of newly developed materials (Biswas,2006)and superhard coatings (Veprek and Veprek-Heijman,2008)that can be used as cutting tools.The optimum combination of hard particles and duc-tile metallic matrices can lead to higher wear resistance.As this principle is widely recognized,particles reinforced metal matrix composites (MMCs)have been developed for cutting tools in a Al-matrix composites (Uday et al.,2009)and in steel matrix com-posites (Li et al.,2010).Among many requirements to cutting tools,cost is always at the top of the list.That is why this research was concentrated on the∗Corresponding author at:BAY-NANO (RI for Nanotechnology),Dep.Nano-composites and Uni.Miskolc,Egyetemvaros,E/7,3515Miskolc,Hungary.Tel.:+36304150002;fax:+3646362916.E-mail addresses:olga ver79@mail.ru (O.Verezub),kalazi@bzaka.hu (Z.Kálazi),kubaisy@mail.ru (A.Sytcheva),femkuzsy@uni-miskolc.hu(L.Kuzsella),buza@bzaka.hu (G.Buza),nikverezub@mail.ru (N.V.Verezub),FedorovA@.au (A.Fedorov),kaptay@ (G.Kaptay).cheapest possible steel matrix.To keep the cost of the cutting tool low,only the surface layer of the cutting tool will be improved,i.e.steel -matrix surface composites will be considered in this paper.It has been established by Iglesias et al.(2007)that the wear resistance of the composite increases with decreasing the size of the hard,reinforcing particles.That is why a special technology has been developed by Verezub et al.(2009)to ensure that the reinforc-ing particles are as small as ing smaller hard particles in the steel matrix of the cutting tool is expected to improve the performance of the cutting tool.There is a wide variety of reinforcing particles for the steel matrix.The most ‘popular’particles are TiC particles (Ala-Kleme et al.,2007)and WC particles embedded in the matrix of M2high-speed tool steels (Riabkina-Fishman et al.,2001)or in Ni–Cr matrix (St-Georges,2007).These particles are hard and thermodynami-cally stable.While TiC has superiority in both hardness (Kalpakjian,1995)and thermodynamic stability (Barin,1993)over WC,both of these particles are metallic in nature,and thus are well wettable by the steel matrix ensuring strong adhesion between the matrix and the reinforcing particles.This is essential to prevent reinforc-ing particles to turn out of the matrix without being worn,what is known to be one of the wear mechanisms of MMCs (Kaptay et al.,1997).In this respect WC is superior to TiC,as WC is perfectly wet-ted,while TiC is only moderately wetted by liquid steel as reviewed0924-0136/$–see front matter ©2010Elsevier B.V.All rights reserved.doi:10.1016/j.jmatprotec.2010.12.009O.Verezub et al./Journal of Materials Processing Technology211(2011)750–758751Table1Composition(in wt%)of the HVG steel(similar to105WCr6steel)(balance:Fe).C Mn Si S P Cu Cr Ni Al Ti Mo Nb W V O0.99 1.000.380.0140.0210.14 1.030.190.0280.0050.10.01 1.340.030.0056by Eustathopoulos et al.(1999)and shown theoretically by Kaptay (2005).Among the many possible technologies to produce steel matrix surface nano-composites,the laser melt injection(LMI)technology has been selected by the authors.This technology was developed 3decades ago by Ayers and Tucker(1980)to produce a surface composite layer.In this technique large(around100m)carbide particles are blown by a gas stream into a moving laser melted pool of a substrate metal.The method is superior to all coating technologies in providing perfect adhesion between the compos-ite and the substrate and also in providing large thickness(around 1mm)allowing to re-ground the cutting ter the LMI tech-nology has been proven to be efficient to produce WC particles reinforced steel matrix composites by Liu et al.(2008),using par-ticularly X40CrMoV5–1steel surface layer by Dobrzanski et al. (2005)and duplex stainless steels matrix by Do Nascimento et al. (2008).The combination of WC+Co particles was used by Bitay and Roósz(2006).TiC particles were added into liquid steel by Fábián et al.(2003).The drawback of LMI technology is that only large carbide par-ticles with a sufficiently large kinetic energy can break the high surface tension liquid metal/gas interface,as was proven for liq-uid steel by Farias and Irons(1985)and for liquid aluminum by Vreeling et al.(2000).This is especially true for low-density TiC particles(compared to the density of liquid steel)that are not‘per-fectly’wetted by liquid steel as shown by Verezub et al.(2005). In fact,incorporation problems for the TiC/liquid steel couple was mentioned in the veryfirst paper by Ayers and Tucker(1980)and was later confirmed by Králik et al.(2003).Thus,the steel rein-forcing matrix,the TiC particles,the LMI technique and the desire to produce surface nano-composite seem to be contradictory.To solve this problem,an in situ LMI technology was developed by Verezub et al.(2009)to produce steel matrix carbide reinforced surface nano-composites.The in situ production route of steel-matrix TiC reinforced composites has been known since the work by Terry and Chinyamakobvu(1991).This method has been developed further by using reaction casting by Feng et al.(2005)and high-energy electron beam irradiation by Lee et al.(2006).The method was extended to produce Fe/(TiW)C composite powder by Correa et al. (2007).Good tribological behaviour of TiC–ferrous matrix com-posites was shown by Kattamis and Suganuma(1990).The Fe/TiC composites were found to have excellent wear properties by Galgali et al.(1999),confirmed also for elevated temperatures by Degnan et al.(2001).The samefinding was extended by Dogan et al.(2001) for cast chromium steels reinforced by TiC particles.Nevertheless, the in situ production of Fe/TiC composites and the LMI technology was combined for thefirst time by Verezub et al.(2009).The goal of the present paper is to evaluate the machinability(upon producing a cutting tool from it)and also the performance as a cutting tool of a steel matrix(TiW)C reinforced surface nano-composite produced on a cheap steel matrix by the in situ LMI process.2.Materials and methods2.1.MaterialsLow-alloyed tool steel plates of grade HVG(Russian GOST5950-73,1973being the analogue of steel105WCr6)have been selected as a base material for the current research.Detailed chemical com-position of the HVG substrate is given in Table1.The initial size of the substrates was8mm×8mm×4mm.Additionally,tungsten carbide and metallic titanium powders of chemical purity,both with a particle size of45–70m were used.These two powders were mixed at a1:1molar ratio.This molar ratio was chosen as the most stable carbides in the Ti–W–C system are the TiC and WC car-bides,and in this way the exchange reaction Ti+WC=W+TiC can be ensured between them.Of-course,the C-content of the original steel will also play some role(see below).2.2.Production of the nano-compositesSchematic diagram of LMI-equipment used in the current study is shown in Fig.1.The upper8mm×8mm plane of the HVG sub-strate was coated by a thin layer of graphite to increase laser beam absorption efficiency during the LMI process.The other side of the steel substrate was brazed onto a large,water-cooled Cu-plate to ensure fast cooling of the substrate.The top surface of the substrate was melted by a2.5kW CO2continuous wave laser(type Trumph TLC105),with a laser spot of2mm in diameter.The laser spot was moving along the sample with a scanning speed of400mm/min. The(WC+Ti)powder mixture was blown into the melted pool at an angle of45◦using argon as carrier gas.The following three pow-der feeding rates were used during our experiments:1.3g/min, 2.3g/min and3.8g/min.Several laser tracks were drawn parallel to each other with a50%overlapping.After the LMI process,the rapidly cooled samples were heat treated under the following conditions:austenitizing at a tempera-ture of1000–1050◦C during10–15s in a high frequency induction furnace,followed by rapid cooling and tempering at a temperature of350◦C for1h.The second round of tempering was performed during1h at560◦C.For reasons of more correct comparison,the initial HVG samples were heat treated under usual conditions (hardening at840◦C and tempering at170◦C).The samples were grinded,polished,etched and analyzed uti-lizing special techniques.The microstructure of the substrates was observed using an AMRAY1810i SEM(Scanning Electron Microscopy with micro resolution),equipped with EDS(Energy Dispersive X-ray Spectroscopy).The identification of nano-sized particles was performed by a high resolution SEM(HitachiS-4800,Fig.1.The schematic diagram of the laser melt injection(LMI)equipment(1–laser, 2–powder nozzle,3–steel substrate to be treated,4–copper cooling plate,5–working table,6–cooling water input and output).752O.Verezub et al./Journal of Materials Processing Technology211(2011)750–758Japan).Quantitative analysis of the samples was performed by ImageJ software.In different parts of the paper the following short sample names are used:i.“LMI”is the sample produced here by the LMI procedure includ-ing heat treatment.ii.“HVG”is the original HVG sample(see Table1)heat treated as described above.iii.“HSS”is a commercially available high speed steel sample with 6%W+5%Mo.2.3.Microhardness measurement of the nano-compositeThe microhardness profiles were measured using TUKON2100B equipment(Wilson Instr.)using load of500g and time of pressing of10s.The samples were polished and etched before the micro-hardness measurements.Microhardness was scanned along two lines:(i)perpendicular to the surface,as function of depth,and(ii) parallel along the surface,at the depth of0.40mm.2.4.Machinability of the nano-composite by cubic BN wheelA cubic boron nitride(CBN)grindingflaring cup wheel of type L010(125/100)–100%–B1–58(Russian standard)which cor-responds to grinding wheel B120C100vitrified bond(Stephenson and Agapiou,2006)was used to remove small quantities of the sur-face composite material to produce the required shape and surface quality for the cutting tool insert.Additionally,the grinding ratio of the CBN wheel was studied by removing the same thickness of 1mm from each substrate(HVG,LMI,HSS).The grinding ratio G,is defined as the ratio of worn mass of the grinding wheel(mg)to the mass of the removed material(g).The CBN wheel was studied by SEM+EDS after the grinding experiments.2.5.Performance of a cutting tool made of LMI nano-composite materialSteel C45(0.45%C+0.6%Mn+0.25%Si)was used as a workpiece for the cutting experiments.Machining of the steel C45workpiece was performed by the HVG,LMI and HSS cutting tools.All the exper-iments were run with the following cutting conditions:cutting speed V=20–60m/min,feed f=0.05–0.3mm/rev and depth of cut d=0.25–1.75mm.The cutting force components were measured by a piezoelectric dynamometer(Kistler).SEM and EDS analysis of the cutting tool and the removed chips were applied after the cutting experiments.3.Results and discussion3.1.Structure and composition of the nano-compositeFig.2shows SEM pictures of cross section of the characteristic LMI sample.Fig.2a shows that the depth of the surface composite layer is approximately1mm.Due to multiple scanning by the laser beam,the depth of the melted layer shows a certain pattern in Fig.2a,with minima in the depth separated by a distance of about 1mm(what is half of the2mm laser spot diameter due to50% of overlapping).As one can see from Fig.2,the microstructure of the melted layer seems to be macroscopically homogeneous.This is due to the high velocity of Marangoni convection of the laser melted pool during the LMI process.Fig.3shows enlarged SEM pictures of the LMI sample with two types of precipitates.The several micron sized(Ti,W)C carbide pre-cipitates(Fig.3a)formed during fast cooling at the latest stage ofthe Fig.2.SEM pictures of the cross sections of the steel substrate after the LMI treat-ment(a)and the general view of the microstructure within the laser treated zone (b)(powder feeding rate is2.3g/min).LMI process by in situ precipitation from the molten steel matrix. The core of these precipitates is rich in Ti,while the outer region of the precipitates is rich in W.This is so due to higher thermodynamic stability of TiC compared to WC.The second type of precipitates is below100nm in diameter and is formed only during the subse-quent heat treatment.These nano-particles are(Fe,W)6C carbides (Fig.3a and b),precipitating from the supersaturated solid steel matrix(for more details see Verezub et al.,2009).The volume%of micron sized(Ti,W)C particles are shown in Fig.4as function of the powder feeding rate.The theoretical maximum,shown in Fig.4was calculated from the technologi-cal parameters and from the cross section of the melted zone(see Fig.2a).One can see that in the as-received LMI samples the amount of incorporated(Ti,W)C particles is somewhat lower compared to the theoretical maximum.The incorporation ratio decreases from 89%(for1.3g/min)to76%(for3.8g/m)with increasing the pow-der feeding rate.It is probably due to the gradual increase in the effective viscosity of the suspension with increasing its solid con-tent,what makes further incorporation and dissolution of(Ti+WC) particles more difficult.During heat treatment of the LMI samples, the volume%of micron-sized(Ti,W)C particles is decreased further by about20%.This is due to the partial dissolution of the W-rich outer region of the micron-sized(Ti,W)C precipitates.The amount of nano-sized(Fe,W)6C particles is found around25±5vol%,being independent of the powder feeding rate.These nano-sized particles form during the heat treatment,from the over-saturated matrix and partially from the dissolved outer regions of the micron-sized precipitates.As follows from materials balance,the majority of the content of these(Fe,W)6C nano-particles originate from the mate-rial of the matrix.Further investigation is needed to clarify howO.Verezub et al./Journal of Materials Processing Technology 211(2011)750–758753Fig.3.SEM micrographs of the cross section of the LMI sample in two different magnifications (powder feeding rate is 2.3g/min).the conditions of heat treatment influence the micro-and nano-structure of the composite and the amount and size distribution of (Fe,W)6C particles.It should be mentioned that at the highest powder feeding rate of 3.8g/min the LMI samples appeared to be cracked.This is probably due to the too high volume %of the carbide phase in the matrix.The two other samples (produced at the powder feeding rates of 1.3and 2.3g/min)are free of cracks.The latter is more promising as the higher amount of carbide phase leads to improved mechanical properties of the composite,if the formation of cracks is avoided.3.2.Microhardness of the LMI nano-composite sampleThe depth profile of microhardness of the LMI nano-composite sample with powder feeding rate of 2.3g/min is shown in Fig.5.All measurements are made after the heat treatment described in the experimental part.The depth profile can be divided into three regions:010*******12345powder feeding rate, g/min(T i ,W )C , v o l %Fig.4.The volume %of the micron-sized (Ti,W)C particles as function of powder feeding rate after the LMI process (before and after the heat treatment procedure).020040060080010001200140000.51 1.52M i c r o h a r d n e s s , HVDistance from surface, mmFig.5.Depth profile of microhardness of the LMI nano-composites (powder feeding rate is 2.3g/min).i.The upper surface layer of about 500m thickness has a highest microhardness of about 1200HV.ii.The initial substrate (below 1mm from the top surface)has a lowest microhardness of about 1000HV.iii.There is a transition zone between 500and 1000m measuredfrom the top surface,within which the microhardness gradually changes between the above mentioned limits.In evaluation of these results let us remind that carbon can diffuse from the non-melted part of the substrate into the melted LMI part of the substrate during the heat treatment.The increased microhardness of the upper surface layer of LMI nano-composite sample is obviously due to the precipitated micron-sized (Ti,W)C and nano-sized (Fe,W)6C hard carbide par-ticles.The existence of the intermediate zone could be due to the interplay between solidification rate (solidification goes from the bottom of the melted zone upwards)and the feeding and mixing rates of the added powder mixture (powder mixture is added to the top and the incorporated particles together with the dissolved atoms move downwards mainly by the Marangoni convection).In Fig.6a the measured microhardness is shown parallel along the sample surface,at the depth of about 0.4mm for both the as received LMI sample and the heat treated LMI sample.One can see that the microhardness of the LMI samples increase due to the heat treatment,what is probably due to the formation of (Fe,W)6C02004006008001000120014000.511.522.53M i c r o h a r d n e s s , H VDistance, mm00.20.40.60.8100.511.52 2.53d e p t h , m mdistance, mmFig.6.Microhardness scanned parallel along the surface,at the depth of about 0.4mm for both as received LMI sample and the heat treated LMI sample (a)and the depth of the melted pool as function of the same path (b)(see the pattern in Fig.2a).754O.Verezub et al./Journal of Materials Processing Technology 211(2011)750–7580.0340f , m m /p a s sv f ,/m i nFig.7.The grinding ratio of the LMI nano-composite sample (powder feeding rate is 2.3g/min,V =25m/s).nanoparticles.It is also obvious that the heat treatment flatters out the large fluctuations in the microhardness of the as received LMI sample.The minima in the microhardness fluctuations (Fig.6a)approximately coincide with the minima in the depth of the melted zone (see Fig.6b and the pattern in Fig.2a).This can be explained by Fig.5,measured at the largest depth of the melted zone.The smaller is the depth of the melted zone,the higher becomes the relative depth of the same absolute depth of 0.4mm,and thus,in accordance with Fig.5,the smaller is the microhardness.As follows from Figs.5–6,the microhardness of the produced nano-composite layer is around 12GPa.For this value the optimum grinding wheel and the optimum workpiece to be machined should be selected such that the ratio of microhardnesses of the machining and that of the to be machined materials should be at least 3.As a machining tool,CBN (cubic boron-nitride)has been selected with its microhardness of about 50GPa (Kalpakjian,1995)being about 4.2times stronger compared to the hardness of our LMI sample.On the other hand,the C45workpiece has been selected with its microhardness of about 2.7GPa,being about 4.4times less strong compared to our LMI sample.Thus,the microhardness of our LMI nano-composite sample is positioned almost in the middle (in a logarithmic scale)of the interval between the microhardness val-ues of the machining CBN tool and that of the to be machined C45workpiece.3.3.Machinability of the LMI nano-composite by cubic BN wheel During the LMI treatment the surface of the substrate melts,and thus it becomes quite uneven after solidification (the subsequent heat treatment does not provide any significant improvement).As a result,the as-received LMI nano-composite sample cannot be used as a cutting tool.Therefore,the as-received LMI nano-composite sample was grinded by a CBN wheel to obtain the shape and surface quality required for cutting tools.Fig.7shows the grinding ratio of the LMI nano-composite sam-ple as function of the feed rate of the workpiece (v f ,m/min)and a feed (f ,mm/pass).The combination of a feed of f =0.04m/pass and a feed rate of v f =3m/min leads to a maximum grinding ratio of about G =45–50mg/g.Based on the results shown in Fig.7,the optimal grinding conditions are selected as:f =0.01–0.02mm/pass,v f =1–2m/min and V =25m/s.Under these conditions the grinding ratio can be kept at a reasonable level of G =8–15mg/g.In compari-son,under the same conditions the grinding ratio for the HVG steel was found to be 6.8mg/g,while the grinding ratio for HSS is known to be about 5–6mg/g (Lisanov,1978).The increased grinding ratio of our LMI sample is obviously due to the hard (Ti,W)C and (Fe,W)6C particles in the surface of newly developed material.Fig.8.The EDS spectra of CBN wheels after grinding HVG (a)and LMI (b)samples (powder feeding rate is 2.3g/min).In Fig.8the energy dispersive X-ray spectra of two CBN wheels are compared after identical grinding runs of the HVG and LMI samples.In addition to the C-and Fe-peaks after grinding the HVG sample,large W and Ti peaks are observed after grinding the LMI nano-composite sample.This can be explained by stabilisation of the C-content of the initial steel substrate by added Ti and by the attraction between (Ti,W)and (B,N)atoms,respectively,being due to the existence of stable titanium boride,titanium nitride and tungsten boride compounds as reported by Barin (1993).Thus,dur-ing the grinding process part of the Ti-and W-content of the LMI nano-composite substrate adheres to the CBN surface.The pres-ence of solid titanium and tungsten carbides causes loosening of the CBN grains and their fallout,leading to intensive wear of the wheel,also shown by Klimenko et al.(1996).Forming the rake and flank surfaces during grinding of the LMI nano-composite samples resembles the grinding of high-speed steel cutting tools as shown by Mamalis et al.(2002).When the high quality alloyed layer is achieved,cutting edge without visible chip-ping is obtained.At the same time the edge roughness,as well as the radius of the cutting edge are higher for the HVG substrate com-pared to the LMI nano-composite substrate (Table 2).Increase of the feed rate and that of the feed lead to further increase in roughness of the tool’s cutting edge.Table 2Roughness of tool’s cutting edge and surfaces after grinding by CBN wheels (param-eters:V =25m/s,v f =2m/min,f =0.01mm/pass).Tool material Cutting edgeroughness R a ,m Roughness of rake and flank surfaces R a ,m HSS 1.2–1.30.15–0.18LMI 1.3–1.50.17–0.20HVG1.4–1.60.21–0.24O.Verezub et al./Journal of Materials Processing Technology 211(2011)750–75875500,10,20,30,40,5050100150200machining time, minV B , m mFig.9.The influence of machining time on flank wear for different cutting tool materials (V =25m/min,f =0.1mm/rev,d =0.5mm).Curves correspond to the HVG steel,LMI nano-composite produced with different powder feeding rates (figures on curves correspond to the unit of g/min),and HSS.Overall it can be concluded that CBN wheels can be used with optimum grinding parameters of f =0.01–0.02mm/pass,v f =1–2m/min and V =25m/s to convert the as-received LMI nano-composite into the cutting tool.The required shape and roughness of the cutting tool can be obtained with a reasonable grinding ratio of about G =8–15mg/g.3.4.Tool life of the cutting tool made of our nano-composite During testing of a new LMI nano-composite cutting tool on C45workpiece,crater wear was found to be negligible compared to flank wear.These two types of wear are the most common measured forms of tool wear.Thus,the tool life of this newLMI2040608010012014016018001234powder feeding rate, g/mint o o l l i f e , m i nFig.10.Tool life as function of the powder feeding rate during the LMI process (V =25m/min,f =0.1mm/rev,d =0.5mm)(the point at zero feeding rate refers to a different heat history of a sample,that is why this point is connected to other points by a thin line).nano-composite cutting tool is determined from the measured flank wear.Fig.9shows flank wear measurements for HVG steel used as a base material,LMI nano-composite produced with different pow-der feeding rates and HSS.The critical flank wear of 0.45mm was chosen based on values recommended for replacing or re-grounding alloyed tool materials (Kalpakjian,1995).The machining time during which the actual flank wear achieves the critical value is called tool life (T ,min).Tool life as function of the powder feeding rate is shown in Fig.10.It shows that an optimum value of the powder feeding rate exists for the maximum tool life.When Fig.10is rationalized in combination with Fig.4,it can be seenthat the volume %of carbide particles in the compositeFig.11.SEM images of the cutting tool made of the LMI sample after its service (a–c)and EDS spectrum (d)of the worn surface.756O.Verezub et al./Journal of Materials Processing Technology211(2011)750–758Fig.12.Removed chip from steel C45by the LMI cutting tool(a)and its EDS spec-trum(b).gradually increases with the increase of powder feeding rate and,as a consequence,tool life also increases.However,as was mentioned above,cracks were formed in the substrate,made by the powder feeding rate of3.8g/min.As a result,a cutting tool made of this substrate has a lower tool life.One can suppose that there is an optimum feeding rate in the interval between2.3and3.8g/min, when the volume%of carbide particles is somewhat larger than for the2.3g/min feeding rate,but still without crack formation.The SEM images of the LMI nano-composite cutting tool faces after machining of C45steel are shown in Fig.11.Fig.11a shows the overlapping of the LMI tracks and the traces of theflank wear(mea-sured as0.45mm).In Fig.11b–c carbide particles being similar to those shown in Figs.2–3are shown.The difference is that the steel matrix is worn away in between the hard carbide particles after machining compared to the initial state of the LMI nano-composite substrate(compare Fig.11b–c to Figs.2–3).Therefore,it is evident that theflank wear is the result of abrasive wear of the LMI cutting tool.The EDS spectrum(Fig.11d)of the worn surface shows Fe,Ti, W as basic components.In Fig.12the SEM picture and EDS spectrum of the removed chip from the C45workpiece is shown,after its machining by the LMI nano-composite cutting tool.It can be seen that the removed chip is continuous,and the main elements of the nano-composite (Ti and W)are missing from its X-ray spectrum.Thus,there was no adhesion of Ti and/or W to the C45steel workpiece during its machining by the LMI nano-composite cutting tool.In order to position our cutting tool made of the LMI substrate on a tool-life scale,tool-life tests have been conducted.The effect of cutting speed V on tool-life T has been assessed using Taylor’s tool life equation(Eq.(1))(Taylor,1907)and the results are shown Table3Experimental tool life(T,min)of different cutting tool materials against C45work-piece(f=0.1mm/rev and d=0.5mm).V,m/min T,min(experimental)HVG LMI HSS20385500535 2550160190 30157088 3542550 40–1726 45–715 50–47 60–24in Table3.C1=V·T n(1) Eq.(1)is widely used and recognized in the industry.It relates tool life to the cutting speed through empirical tool life constants n and C1.Table4shows the range of values n and C1for different cutting tool materials obtained from the data in Table3.The data(Table4)indicate that the LMI process of inserting (TiW)C particles into HVG substrate improved tool life of the base material by300–400%when cutting speed V was25–35m/min. However,this new material was felt short to surpass tool life of HSS cutting material by just20%in the same cutting speed range.No sig-nificant difference between tool life of HSS and LMI was observed during machining at20m/min.Performance of cutting tools made of LMI nano-composite is similar to the performance of HSS and is limited by wear resistance at cutting speeds above40–45m/min (Stephenson and Agapiou,2006).Cutting tools made of HVG steel can be used at speeds up to30m/min.3.5.Cutting force componentsDuring machining of the C45workpieces by the cutting tool made of HVG and LMI substrates,the two main force components F z (N)and F x(N)have been measured as function of the depth of cut d (mm)and feed f(mm/rev).The effects of depth of cut and feed on the measured F z and F x force components for the two different cutting materials(HVG and LMI nano-composite)are shown in Figs.13–14. The cutting speed increase within limits of V c=20–60m/min does not sufficiently influence the value of the cutting force(Fedorov, 2005)and therefore has not been tested in this paper.The effect of depth of cut d on the measured force components for two different cutting tool materials is shown in Fig.13a and b. Forces F z and F x increase with the increase in depth of cut because the increase in depth of cut leads to increase in the area of cut and length of the cutting edge in contact.The influence of feed f on the forces F z and F x is shown in Fig.14a and b.The increases in feed lead to increase in cut thickness,which,in turn,increases the area of cut and as a consequence,the force components.Figs.13–14show that machining with LMI nano-composite cut-ting tool material decreases cutting forces F z and F x in comparison with HVG cutting tool material.However,significant force reduc-tions can only be observed when depth of cut is greater than1mm and feed is greater than0.2mm/rev.The force components can be described as function of parameters d and f by the followingTable4Values of n and C1for different tool materials(f=0.1mm/rev,d=0.5mm).Cutting tool material n C1HSS0.2280.89 LMI0.1967.18 HVG0.1241.58。
大能量中空光束大气传输的仿真与实验比对研究赵琦;樊红英;李轶国;蒋泽伟;胡绍云;赖庚辛;黄燕琳;耿旭【摘要】To study the characteristics of high energy pulse laser propagating through turbulent atmosphere , the propagation of high energy laser in atmosphere was numerically simulated , and the far-field beam quantity in terms of the power in the bucket was compared with the experiment performed by using an Nd ∶glass laser at output of 800J.The error between the numerical simulation and experiment was analyzed in detail .The numerical simulation and experiments were studied under different conditions of visibility , output energy , and turbulence .It is shown that the numerical simulation can be used to predict the experimental results by a suitable choice of structure parameters of the phase screen .The results show the significance to guide the application and development of high energy pulse lasers .%为了研究高能脉冲激光大气传输特性，采用桶中功率和光斑半径作为远场光束质量评价标准，与800J钕玻璃激光器的实验结果进行了比较，仿真计算和实验中分别考虑了不同能见度、出射光能量和湍流强度下远场光斑半径的变化趋势；对数值模拟与实验结果分别进行了分析和比对，并对两者之间的误差做了详细分析。
29.F.F.Balakirev et al.,Phys.Rev.Lett.102,017004(2009).30.F.Rullier-Albenque et al.,Phys.Rev.Lett.99,027003(2007).31.T.Senthil,Phys.Rev.B78,035103(2008).32.A.Kanigel et al.,Nat.Phys.2,447(2006).33.J.W.Loram,K.A.Mirza,J.R.Cooper,J.L.Tallon,J.Phys.Chem.Solids59,2091(1998).34.T.Yoshida et al.,J.Phys.Condens.Matter19,125209(2007).35.J.Zaanen,Nature430,512(2004).36.J.W.Loram,J.Luo,J.R.Cooper,W.Y.Liang,J.L.Tallon,J.Phys.Chem.Solids62,59(2001).37.C.Panagopoulos et al.,Phys.Rev.B67,220502(2003).38.H.J.A.Molegraaf,C.Presura,D.van der Marel,P.H.Kes,M.Li,Science295,2239(2002).39.S.Chakraborty,D.Galanakis,P.Phillips,/abs/0807.2854(2008).40.P.Phillips,C.Chamon,Phys.Rev.Lett.95,107002(2005).41.We acknowledge technical and scientific assistance fromS.L.Kearns,J.Levallois,and N.Mangkorntang andcollaborative support from H.H.Wen.This work wassupported by Engineering and Physical Sciences ResearchCouncil(UK),the Royal Society,Laboratoire National desChamps Magnétiques Pulsés,the French AgenceNationale de la Recherche IceNET,and EuroMagNET.Supporting Online Material/cgi/content/full/1165015/DC1Materials and MethodsFigs.S1and S2References22August2008;accepted21November2008Published online11December2008;10.1126/science.1165015Include this information when citing this paper.Revealing the Maximum Strengthin Nanotwinned CopperL.Lu,1*X.Chen,1X.Huang,2K.Lu1The strength of polycrystalline materials increases with decreasing grain size.Below a critical size,smaller grains might lead to softening,as suggested by atomistic simulations.The strongest size should arise at a transition in deformation mechanism from lattice dislocation activities to grain boundary–related processes.We investigated the maximum strength of nanotwinned copper samples with different twin thicknesses.We found that the strength increases with decreasing twin thickness,reaching a maximum at 15nanometers,followed by a softening at smaller values that is accompanied by enhanced strain hardening and tensile ductility.The strongest twin thickness originates from a transition in the yielding mechanism from the slip transfer across twin boundaries to the activity of preexisting easy dislocation sources.T he strength of polycrystalline materials increases with decreasing grain size,asdescribed by the well-known Hall-Petch relation(1,2).The strengthening originates from the fact that grain boundaries block the lattice dislocation motion,thereby making plastic defor-mation more difficult at smaller grain sizes.How-ever,below a certain critical size,the dominating deformation mechanism may change from lattice dislocation activities to other mechanisms such as grain boundary–related processes,and softening behavior(rather than strengthening)is expected (3,4).Such a softening phenomenon has been demonstrated by atomistic simulations,and a crit-ical grain size of maximum strength has been predicted(5–7).In pure metals,an impediment to determining the grain size that yields the highest strength is the practical difficulty of obtaining sta-ble nanostructures with extremely small structural domains(on the order of several nanometers).The driving force for growth of nanosized grains in pure metals,originating from the high excess en-ergy of numerous grain boundaries,becomes so large that grain growth may take place easily even at ambient temperature or below.Coherent twin boundaries(TBs),which aredefined in a face-centered cubic structure as the(111)mirror planes at which the normal stackingsequence of(111)planes is reversed,are known tobe as effective as conventional grain boundariesin strengthening materials.Strengthening has beenobtained in Cu when high densities of nanometer-thick twins are introduced into submicrometer-sized grains(8–10).In addition,coherent TBs aremuch more stable against migration(a fundamen-tal process of coarsening)than conventional grainboundaries,as the excess energy of coherent TBs isone order of magnitude lower than that of grainboundaries.Hence,nanotwinned structures areenergetically more stable than nanograined coun-terparts with the same chemical composition.Thestable nanotwinned structure may provide samplesfor exploring the softening behavior with very smalldomain sizes.Here,we prepared nanotwinned pureCu(nt-Cu)samples with average twin thicknessranging from a few nanometers to about100nm.High-purity(99.995%)Cu foil samples com-posed of nanoscale twin lamellae embedded insubmicrometer-sized grains were synthesized bymeans of pulsed electrodeposition.By increasingthe deposition rate to10nm/s,we succeeded inrefining the mean twin thickness(i.e.,the meanspacing between adjacent TBs,hereafter referredto as l)from a range of15to100nm down to arange of4to10nm(see supporting online ma-terial).The as-deposited Cu foils have an in-planedimension of20mm by10mm and a thicknessof30m m with a uniform microstructure.Shownin Fig.1,A to C,are transmission electron mi-croscopy(TEM)plane-view images of three as-deposited samples with l values of96nm,15nm,and4nm,respectively.The TEM images indicatethat some grains are irregular in shape,but low-magnification scanning electronic microscopyimages,both cross section and plane view,showthat the grains are roughly equiaxed in three di-mensions.Grain size measurements showed asimilar distribution and a similar average diam-eter of about400to600nm for all nt-Cu samples.Twins were formed in all grains(see the electrondiffraction pattern in Fig.1D),and observationsof twins in a large number of individual grainsrevealed no obvious change in the twin densityfrom grain to grain.Note that in all samples,theedge-on twins that formed in different grains arealigned randomly around the foil normal(growth)direction(8,11),in agreement with a strong[110]texture determined by x-ray diffraction(XRD).Foreach sample,twin thicknesses were measured froma large number of grains,which were detected fromnumerous TEM and high-resolution TEM(HRTEM)images,to generate a distribution.Figure1E illus-trates the492measurements for the sample withthe finest twins;the majority yielded spacings be-tween twins smaller than10nm,with a mean of4nm.For simplicity,each nt-Cu sample is iden-tified by its mean twin thickness;for example,thesample with l=4nm is referred to as nt-4.Figure2shows the uniaxial tensile true stress–true strain curves for nt-Cu samples of various lvalues.Also included are two stress-strain curvesobtained from a coarse-grained Cu(cg-Cu)andan ultrafine-grained Cu(ufg-Cu)that has a sim-ilar grain size to that of nt-Cu samples but is freeof twins within grains.Two distinct features areobserved with respect to the l dependence of themechanical behavior of nt-Cu.The first is the oc-currence of the l giving the highest strength.Allstress-strain curves of nt-Cu samples in Fig.2,Aand B,are above that of the ufg-Cu,indicating astrengthening by introducing twins into the sub-micrometer grains.However,such a strengthen-ing does not show a linear relationship with l.Forl>15nm(Fig.2A),the stress-strain curves shiftupward with decreasing l,similar to the strength-ening behavior reported previously in the nt-Cu(9,11)and nanocrystalline Cu(nc-Cu)(12–15)samples(Fig.3A).However,with further de-1Shenyang National Laboratory for Materials Science,Institute of Metal Research,Chinese Academy of Sciences,Shenyang 110016,P.R.China.2Center for Fundamental Research:Metal Structures in Four Dimensions,Materials Research Department, RisøNational Laboratory for Sustainable Energy,Technical Uni-versity of Denmark,DK-4000Roskilde,Denmark.*To whom correspondence should be addressed.E-mail: llu@ o n J a n u a r y 3 0 , 2 0 0 9 w w w . s c i e n c e m a g . o r g D o w n l o a d e d f r o mcreases of l down to extreme dimensions (i.e.,less than 10nm),the stress-strain curves shift down-ward (Fig.2B).As plotted in Fig.3A,the mea-sured yield strength s y (at 0.2%offset)shows a maximum value of 900MPa at l ≈15nm.The second feature is a substantial increase in tensile ductility and strain hardening when l <15nm.As seen in Fig.2,the tensile elongation of the nt-Cu samples increases monotonically with decreasing l .When l <15nm,the uniform ten-sile elongation exceeds that of the ufg-Cu sample,reaching a maximum value of 30%at the finest twin thickness.Strain-hardening coefficient (n )values were determined for each sample by fitting the uniform plastic deformation region to s =K 1+K 2e n ,where K 1represents the initial yield stress and K 2is the strengthening coefficient (i.e.,the strength increment due to strain hardening at strain e =1)(16,17).The n values determined for all the nt-Cu samples increase monotonically with de-creasing l (Fig.3B),similar to the trend of uniform elongation versus l .When l <15nm,n exceeds the value for cg-Cu (0.35)(16,17)and finally reaches a maximum of 0.66at l =4nm.The twin refinement –induced increase in n is opposite to the general observation in ultrafine-grained and nanocrystalline materials,where n continuously decreases with decreasing grain size (Fig.3B).The strength of the nt-Cu samples has been considered to be controlled predominantly by the nanoscale twins via the mechanism of slip trans-fer across the TBs (10,18),and it increases with decreasing l in a Hall-Petch –type relationship (9)similar to that of grain boundary strength-ening in nanocrystalline metals (12).Our re-sults show that such a relationship breaks down when l <15nm,although other structural pa-rameters such as grain size and texture are un-changed.The grain sizes of the nt-Cu samples are in the submicrometer regime,which is too large for grain boundary sliding to occur at room temperature,as expected for nanocrystalline ma-terials with grain sizes below 20nm (3).There-fore,the observed softening cannot be explained by the initiation of grain boundary –mediated mechanisms such as grain boundary sliding and grain rotation,as proposed by molecular dynam-ics (MD)simulations for nanocrystalline mate-rials (3).To explore the origin of the twin thickness giv-ing the highest strength,we carried out detailed structural characterization of the as-deposited sam-ples.HRTEM observations showed that in each sample TBs are coherent S 3interfaces associated with the presence of Shockley partial dislocations (as steps),as indicated in Fig.1D.These partial dislocations have their Burgers vector parallel to the twin plane and are an intrinsic structural fea-ture of twin growth during electrodeposition.The distribution of the preexisting partial dislocations is inhomogeneous,but their density per unit area of TBs is found to be rather constant among sam-ples with different twin densities.This suggests that the deposition parameters and the twin re-finement have a negligible effect on the nature ofTBs.Therefore,as a consequence of decreasing l ,the density of such TB-associated partial dislocations per unit volume increases.We also noticed that grain boundaries in the nt-Cu samples with l ≤15nm are characterized by straight segments (facets)that areoftenTrue strain (%)T r u e s t r e s s (M P a )True strain (%)Fig.2.Uniaxial tensile true stress –true strain curves for nt-Cu samples tested at a strain rate of 6×10−3s −1.(A )Curves for samples with mean twin thickness varying from 15to 96nm;(B )curves for samples with mean twin thickness varying from 4to 15nm.For comparison,curves for a twin-free ufg-Cu with a mean grain size of 500nm and for a cg-Cu with a mean grain size of 10m m areincluded.C200 nmA BbbDFig.1.TEM images of as-deposited Cu samples with 15nm.(C )l =4nm.(D )The same sample as (C)but at higher resolution,with a corresponding electron diffraction pattern (upper right inset)and a HRTEM image of the outlined area showing the presence of Shockley partials at the TB (lower right inset).(E )Distribution of the lamellar twin thicknesses determined from TEM and HRTEM images for l =4nm.REPORTSo n J a n u a r y 30, 2009w w w .s c i e n c e m a g .o r g D o w n l o a d e d f r o massociated with dislocation arrays (19),whereas in samples with coarser twins,grain boundaries are smoothly curved,similar to conventional grain boundaries.The microstrain measured by XRD was a negligible 0.01%for samples with l ≥15nm,but increased gradually from 0.038%to 0.057%when l decreased from 10to 4nm,which also indicates a gradual increase in the de-fect density.Recent experimental studies and MD simula-tions (3,20,21)have shown that an increase in the density of preexisting dislocations in nano-scale materials will cause softening.In the nt-Cu samples studied,both the dislocation arrays asso-ciated with the grain boundaries and the steps associated with the preexisting partial dislocations along TBs could be potential dislocation sources,which are expected to affect the initiation of plas-tic deformation (22)and to provide the disloca-tions required for the dislocation-TB interactions that cause work hardening.The preexisting par-tial dislocations can act as readily mobile dislo-cations,and their motion may contribute to the plastic yielding when an external stress is applied to the sample.The plastic strains induced by the motion of preexisting partial dislocations can be estimated as e =r 0b s d/M (where r 0is the initial dislocation density,b s is the Burgers vector of Shockley partial dislocation,d is the grain size,and M is the Taylor factor).Calculations showed that for the samples with l >15nm,the preexist-ing dislocations induce a negligibly small plastic strain (<0.05%).However,for the nt-4specimen,a remarkable amount of plastic strain,as high as 0.1to 0.2%,can be induced just by the motions of high-density preexisting dislocations at TBs (roughly 1014m −2),which could control the mac-roscopic yielding of the sample.The above anal-ysis suggests that for extremely small values of l ,a transition in the yielding mechanism can result in an unusual softening phenomenon in which the preexisting easy dislocation sources at TBs andgrain boundaries dominate the plastic deforma-tion instead of the slip transfer across TBs.Shockley partial dislocations are always in-volved in growth of twins during crystal growth,thermal annealing,or plastic deformation.Shock-ley partials might be left at TBs when the twin growth is interrupted.Therefore,the presence of Shockley partials at some TBs is a natural phe-nomenon.Although these preexisting dislocations may have a small effect on the mechanical be-havior of the samples with thick twins,the effect will be much more pronounced in the samples with nanoscale twins and/or with high preexist-ing TB dislocation densities such as those seen in deformation twins (23).To understand the extraordinary strain hard-ening,we analyzed the deformation structures of the tensile-deformed samples.In samples with coarse twins,tangles and networks of perfect dis-locations were observed within the lattice between the TBs (Fig.4A),and the dislocation density was estimated to be on the order of 1014to 1015m −2.In contrast,high densities of stacking faults and Shockley partials associated with the TBs were found to characterize the deformed structure of the nt-4sample (Fig.4,B and C),indicating the interactions between dislocations and TBs.Recent MD simulations (18,24,25)showed that when an extended dislocation (two Shockley partials connected by a stacking fault ribbon)is forced by an external stress into a coherent TB,it recom-bines or constricts into a perfect dislocation con-figuration at the coherent TB and then slips through the boundary by splitting into three Shockley par-tials.Two of them glide in the slip plane of the adjacent twin lamella,constituting a new extended dislocation,whereas the third one,a twinning par-tial,glides along the TB and forms a step.It is expected that with increasing strain,such an in-teraction process will generate a high density of partial dislocations (steps)along TBs and stack-ing faults that align with the slip planes in the twin lamellae,which may (or may not)connect to the TBs.Such a configuration of defects was observed,as shown in Fig.4C.The density of partial dislocations in the deformed nt-4sampleFig.3.Variation of (A )yield strength and (B )strain hardening coef-ficient n as a function of mean twin thickness for the nt-Cu samples.For comparison,the yield strength and n values fornc-Cu [▲(12),◀(13),▶(14),and ◆(15)],ufg-Cu[▾(9)],andcg-Cu samples reported in the literature are included.A maximum in the yieldstress is seen for thent-Cu with l =15nm,but this has not beenobserved for the nc-Cu,even when the grain size is as small as 10nm.0.00.20.40.60.8nor d (nm)0200400600800σy (M P a )λ or d (nm)020406080100120110100100010000λ2 nmTTB200 nmA CFig.4.(A )A typical bright TEM image of the deformed nt-96sample showing the tangling of lattice dislocations.(B )An HRTEM image of the nt-4sample tensile-deformed to a plastic strain of 30%,showing a high density of stacking faults (SF)at the TB.(C )The arrangement of Shockley partials and stacking faults at TBs within the lamellae in the nt-4sample.Triangles,Shockley partial dislocations associated with stacking faults;⊥,partials with their Burgers vector parallel to the TB plane.REPORTSo n J a n u a r y 30, 2009w w w .s c i e n c e m a g .o r g D o w n l o a d e d f r o mwas estimated to be 5×1016m −2on the basis of the spacing between the neighboring partials and l .This is two orders of magnitude higher than that of the preexisting dislocations and the lattice dislocations stored in the coarse twins.Such a finding suggests that decreasing the twin thick-ness facilitates the dislocation-TB interactions and affords more room for storage of dislocations,which sustain more pronounced strain hardening in the nt-Cu (26,27).These observations suggest that the strain-hardening behavior of nt-Cu samples is governed by two competing processes:dislocation-dislocation interaction hardening in coarse twins,and dislocation-TB interaction hardening in fine twins.With a refining of l ,the contribution from the latter mech-anism increases and eventually dominates the strain hardening,as revealed by the continuous increase of n values (Fig.3B).However,the former hard-ening mechanism usually leads to an inverse trend,diminishing with size refinement (17).Twins are not uncommon in nature,and they appear in various metals and alloys with different crystallographic structures.Extremely thin twin lamellae structures can possibly be achieved under proper conditions during crystal growth,plastic deformation,phase transformations,or thermal annealing of deformed structures.Our finding of the twin thickness giving maximum strength il-lustrates that the scale-dependent nature of plastic deformation of nanometer-scale materials is not necessarily related to grain boundary –mediated processes.This finding also provides insight into the development of advanced nanostructured materials.References and Notes1.E.O.Hall,Proc.Phys.Soc.London Ser.B 64,747(1951).2.N.J.Petch,J.Iron Steel Inst.174,25(1953).3.J.Schiøtz,K.W.Jacobsen,Science 301,1357(2003).4.S.Yip,Nature 391,532(1998).5.M.A.Meyers,A.Mishra,D.J.Benson,Prog.Mater.Sci.51,427(2006).6.P.G.Sanders,J.A.Eastman,J.R.Weertman,Acta Mater.45,4019(1997).7.C.C.Koch,K.M.Youssef,R.O.Scattergood,K.L.Murty,Adv.Eng.Mater.7,787(2005).8.L.Lu et al .,Acta Mater.53,2169(2005).9.Y.F.Shen,L.Lu,Q.H.Lu,Z.H.Jin,K.Lu,Scr.Mater.52,989(2005).10.X.Zhang et al .,Acta Mater.52,995(2004).11.L.Lu,Y.Shen,X.Chen,L.Qian,K.Lu,Science 304,422(2004);published online 18March 2004(10.1126/science.1092905).12.J.Chen,L.Lu,K.Lu,Scr.Mater.54,1913(2006).13.S.Cheng et al .,Acta Mater.53,1521(2005).14.Y.Champion et al .,Science 300,310(2003).15.Y.M.Wang et al .,Scr.Mater.48,1851(2003).16.A.Misra,X.Zhang,D.Hammon,R.G.Hoagland,Acta Mater.53,221(2005).17.M.A.Meyers,K.K.Chawla,in Mechanical Behavior of Materials ,M.Horton,Ed.(Prentice Hall,Upper Saddle River,NJ,1999),pp.112–135.18.Z.H.Jin et al .,Scr.Mater.54,1163(2006).19.X.H.Chen,L.Lu,K.Lu,J.Appl.Phys.102,083708(2007).20.X.Huang,N.Hansen,N.Tsuji,Science 312,249(2006).21.Z.W.Shan,R.K.Mishra,S.A.Syed Asif,O.L.Warren,A.M.Minor,Nat.Mater.7,115(2008).22.K.Konopka,J.Mizera,J.W.Wyrzykowski,J.Mater.Process.Technol.99,255(2000).23.Y.S.Li,N.R.Tao,K.Lu,Acta Mater.56,230(2008).24.S.I.Rao,P.M.Hazzledine,Philos.Mag.A 80,2011(2000).25.Z.H.Jin et al .,Acta Mater.56,1126(2008).26.M.Dao,L.Lu,Y.Shen,S.Suresh,Acta Mater.54,5421(2006).27.T.Zhu,J.Li,A.Samanta,H.G.Kim,S.Suresh,Proc.Natl.Acad.Sci.U.S.A.104,3031(2007).28.Supported by National Natural Science Foundation ofChina grants 50431010,50621091,50725103,and 50890171,Ministry of Science and Technology of China grant 2005CB623604,and the Danish National Research Foundation through the Center for FundamentalResearch:Metal Structures in Four Dimensions (X.H.).We thank N.Hansen,Z.Jin,W.Pantleon,and B.Ralph for stimulating discussions,X.Si and H.Ma for sample preparation,S.Zheng for TEM observations,and Y.Shen for conducting some of the tensile tests.Supporting Online Material/cgi/content/full/323/5914/607/DC1Materials and Methods Table S1References24October 2008;accepted 30December 200810.1126/science.1167641Control of Graphene ’s Properties by Reversible Hydrogenation:Evidence for GraphaneD.C.Elias,1*R.R.Nair,1*T.M.G.Mohiuddin,1S.V.Morozov,2P.Blake,3M.P.Halsall,1A.C.Ferrari,4D.W.Boukhvalov,5M.I.Katsnelson,5A.K.Geim,1,3K.S.Novoselov 1†Although graphite is known as one of the most chemically inert materials,we have found that graphene,a single atomic plane of graphite,can react with atomic hydrogen,which transforms this highly conductive zero-overlap semimetal into an insulator.Transmission electron microscopy reveals that the obtained graphene derivative (graphane)is crystalline and retains the hexagonal lattice,but its period becomes markedly shorter than that of graphene.The reaction with hydrogen is reversible,so that the original metallic state,the lattice spacing,and even the quantum Hall effect can be restored by annealing.Our work illustrates the concept of graphene as a robust atomic-scale scaffold on the basis of which new two-dimensional crystals with designed electronic and other properties can be created by attaching other atoms and molecules.Graphene,a flat monolayer of carbon atoms tightly packed into a honeycomb lattice,continues to attract immense interest,most-ly because of its unusual electronic properties and effects that arise from its truly atomic thick-ness (1).Chemical modification of graphene has been less explored,even though research on car-bon nanotubes suggests that graphene can be al-tered chemically without breaking its resilient C-C bonds.For example,graphene oxide is graphene densely covered with hydroxyl and other groups (2–6).Unfortunately,graphene oxide is strongly disordered,poorly conductive,and difficult to reduce to the original state (6).However,one can imagine atoms or molecules being attached to the atomic scaffold in a strictly periodic manner,which should result in a different electronic struc-ture and,essentially,a different crystalline mate-rial.Particularly elegant is the idea of attaching atomic hydrogen to each site of the graphene lattice to create graphane (7),which changes the hybridization of carbon atoms from sp 2into sp 3,thus removing the conducting p -bands and open-ing an energy gap (7,8).Previously,absorption of hydrogen on gra-phitic surfaces was investigated mostly in con-junction with hydrogen storage,with the research focused on physisorbed molecular hydrogen (9–11).More recently,atomic hydrogen chem-isorbed on carbon nanotubes has been studied theoretically (12)as well as by a variety of exper-imental techniques including infrared (13),ultra-violet (14,15),and x-ray (16)spectroscopy and scanning tunneling microscopy (17).We report the reversible hydrogenation of single-layer graphene and observed dramatic changes in its transport properties and in its electronic and atomic struc-ture,as evidenced by Raman spectroscopy and transmission electron microscopy (TEM).Graphene crystals were prepared by use of micromechanical cleavage (18)of graphite on top of an oxidized Si substrate (300nm SiO 2)and then identified by their optical contrast (1,18)and distinctive Raman signatures (19).Three types of samples were used:large (>20m m)crystals for Raman studies,the standard Hall bar de-vices 1m m in width (18),and free-standing mem-branes (20,21)for TEM.For details of sample fabrication,we refer to earlier work (18,20,21).1School of Physics and Astronomy,University of Manchester,M139PL,Manchester,UK.2Institute for Microelectronics Tech-nology,142432Chernogolovka,Russia.3Manchester Centre for Mesoscience and Nanotechnology,University of Manches-ter,M139PL,Manchester,UK.4Department of Engineering,Cambridge University,9JJ Thomson Avenue,Cambridge CB3OFA,UK.5Institute for Molecules and Materials,Radboud University Nijmegen,6525ED Nijmegen,Netherlands.*These authors contributed equally to this work.†To whom correspondence should be addressed.E-mail:Kostya@REPORTSo n J a n u a r y 30, 2009w w w .s c i e n c e m a g .o r g D o w n l o a d e d f r o m。
