Synthesis of tetragonal BaTiO3 film on Ti substrate by micro-arc oxidation
第八届国际凝聚态理论与计算材料学会议
大会将请专家对以下领域作专题性的特邀报告:
(1). 凝聚态物理的最新进展:
A. 自旋电子学
B. 纳米材料
C. 固体量子信息和计算
D. 玻色－爱因斯坦凝聚
E. 强关联电子系统
F. 高温超导
G. 量子霍尔效应
H. 磁学
I 表面和界面
J. 半导体物理
K. 低维凝聚态物理
L. 介观物理
M. 软凝聚态物质
N. 生物物理
structure Metallic Phases in Solid Germane (GeH4) under Pressure 锗烷在
压力下的金属相
11:45 -12:15
12:15 -13:00
午餐
张振宇 橡树岭国家实验室
柳百新
清华大学
林海青 香港中文大学 牛谦 美国德克萨斯大学
7 月 16 日
凝聚态理论专题邀请报告 I 主持人：谢心澄 地点：芙蓉厅
14:30 -14:55 14:55 -15:20 15:20 -15:45 15:45 -16:10 16:10 -16:35 16:35 -16:50
7 月 16 日
Manipulating Magnetization States of Nanostructures Two-spin decoherence in semiconductor quantum dots
稀磁半导体(Al,Cr)N 中的氮空位与高温铁磁性的第一原理研究
蒋青 刘邦贵
吉林大学 中科院物理所
15:20 -15:45 15:45 -16:10 16:10 -16:35
Theoretical and experimental studies of semiconductor dilute nitrides and devices
砷化镓热氧化法制备β-Ga2O3体块薄膜
第43卷第7期2022年7月Vol.43No.7July,2022发光学报CHINESE JOURNAL OF LUMINESCENCE砷化镓热氧化法制备β‑Ga2O3体块薄膜刁肇悌，陈威，董鑫*，焦腾，李政达（吉林大学电子科学与工程学院集成光电子学国家重点联合实验室，吉林长春130012）摘要：通过对p型砷化镓（p‑GaAs）单晶衬底高温热氧化方法制备了β‑Ga2O3体块薄膜。
探讨了高温氧化过程中O2流量对β‑Ga2O3体块薄膜形貌的影响。
通过对体块薄膜的晶体质量、结构特性、光致发光特性的测试分析，可以发现在高温高氧环境下GaAs转变为β‑Ga2O3体块薄膜的过程与Langmuir蒸发相关。
当O2流量较低时（0.2L/min），GaAs衬底处于缺氧状态，所制备样品呈现纳米线状形貌；而当通入O2流量超过0.4L/min时，GaAs 衬底被完全氧化为具有纳米岛状结构的β‑Ga2O3体块薄膜，且晶体质量得到显著提高。
本文提出的砷化镓单晶热氧化工艺可以高效、低成本地获得较高结晶质量的纳米结构β‑Ga2O3体块薄膜，对于β‑Ga2O3材料的应用具有极大的丰富作用。
关键词：热氧化；纳米岛状体块薄膜；β-Ga2O3；氧气流量中图分类号：O482.31文献标识码：A DOI：10.37188/CJL.20210399β-Ga2O3Bulk Films Prepared by Thermal Oxidation of GaAs DIAO Zhao-ti，CHEN Wei，DONG Xin*，JIAO Teng，LI Zheng-da（State Key Laboratory of Integrated Optoelectronics，College of Electronic Science and Engineering，Jilin University，Changchun130012，China）*Corresponding Author，E-mail：dongx@Abstract：In this paper，β-Ga2O3bulk films were prepared by high-temperature thermal oxidation of p-type gallium arsenide（p-GaAs）single crystal substrate.The effect of O2flow rate on the morphol‑ogy ofβ-Ga2O3bulk films was also discussed.By testing and analysing the crystal quality，structural characteristics and photoluminescence characteristics of the bulk films，we found that the transfor‑mation of GaAs intoβ-Ga2O3bulk films was related to Langmuir evaporation.When the O2flow rate was low（0.2L/min），GaAs substrate was in anoxic state，and the prepared samples showed nanolin‑ear morphology.However，when the flow rate of O2exceeds0.4L/min，GaAs substrate is completely oxidized intoβ-Ga2O3bulk film with nanometer island structure，and the crystal quality is signifi‑cantly improved.The thermal oxidation process based on the GaAs single crystal proposed in this ar‑ticle can obtain nano-structuredβ-Ga2O3bulk films of high crystal quality with high efficiency and low cost.This process enriches the application ofβ-Ga2O3materials greatly.Key words：thermal oxidation；nano island bulk film；β-Ga2O3；oxygen flow rate1引言作为第三代半导体材料，氧化镓（Ga2O3）近年来得到了业界广泛的关注。
BaTiO3薄膜的制备及其在电光调制器的应用
第52卷第4期2023年4月人㊀工㊀晶㊀体㊀学㊀报JOURNAL OF SYNTHETIC CRYSTALS Vol.52㊀No.4April,2023BaTiO 3薄膜的制备及其在电光调制器的应用任怡静,马新国,张㊀锋,陆晶晶,张㊀力,王㊀晗(湖北工业大学芯片产业学院,武汉㊀430068)摘要:BaTiO 3凭借其较高的电光系数㊁优良的压电性质和非线性光学性质,成为制备高性能电光调制器的关键材料㊂其制备方法㊁实验条件和衬底的选择等决定了薄膜的质量㊁生长取向和电光系数等,进而影响电光调制器的传播损耗㊁半波电压㊁消光比等性能㊂本文从电光调制器的工作原理出发,围绕BaTiO 3薄膜的制备方法㊁实验条件和薄膜衬底等,探讨了成膜的取向和质量的影响因素,分析了各个BaTiO 3薄膜制备方法的优缺点,论述了波导的结构及参数,并展望了未来优化BaTiO 3薄膜制备和波导制作工艺的方向㊂关键词:BaTiO 3薄膜;电光调制器;波导;电光系数;化学气相沉积法;分子束外延法中图分类号:O484;O439㊀㊀文献标志码:A ㊀㊀文章编号:1000-985X (2023)04-0688-13Preparation of BaTiO 3Thin Film and Its Application in Electro-Optic ModulatorREN Yijing ,MA Xinguo ,ZHANG Feng ,LU Jingjing ,ZHANG Li ,WANG Han (School of Chip Industry,Hubei University of Technology,Wuhan 430068,China)Abstract :BaTiO 3has become the key materials for the preparation of high-performance electro-optical modulators due to its high electro-optical coefficient,excellent piezoelectric properties and nonlinear optical properties.The quality,growth orientation and electro-optic coefficient of the film are determined by the preparation method,experimental condition and the selection of substrate,and further affecting the propagation loss,half-wave voltage and extinction ratio of the electro-optical modulator.Based on the working principle of electro-optical modulator,the factors affecting the orientation and quality of film forming were discussed by focusing on the preparation methods,experimental conditions and film substrates of BaTiO 3thin film in this paper.The advantages and disadvantages of each process of BaTiO 3thin film preparation were analyzed,the structure and performance of waveguide were discussed.Finally,the direction of optimizing BaTiO 3thin film fabrication and waveguide fabrication process in the future were summarized in this paper.Key words :BaTiO 3thin film;electro-optic modulator;waveguide;electro-optic coefficient;chemical vapor deposition method;molecular beam epitaxy method ㊀㊀收稿日期:2022-12-07㊀㊀基金项目:国家自然科学基金(51472081);结构化学国家重点实验室科学基金(20210028);湖北工业大学绿色工业引领计划杰出青年基金(JCRC2021003)㊀㊀作者简介:任怡静(1998 ),女,山东省人,博士研究生㊂E-mail:renyij1@ ㊀㊀通信作者:马新国,博士,教授㊂E-mail:maxg@0㊀引㊀㊀言光互联具有功耗小㊁成本低㊁集成度高,以及通信容量大等优势,在信息传输领域承担重要角色[1-2]㊂电光调制器作为光互联的关键组成部分,成为近年来光通信领域的研究热点㊂电光调制器性能的好坏与波导层材料的本征性质㊁微结构㊁包层间折射率和波导结构有密切关系,因此选择合适的波导层和包层材料,以及优化其制备工艺,成为降低半波电压和提高电光带宽等参数的主要手段㊂电光调制器波导材料较为常用的是铌酸锂(LiNbO 3,LNO)㊁砷化镓(GaAs)和钽酸锂(LiTaO 3)㊂其中LNO 材料电光系数达到30.8pm /V [3],其电光调制器需要的驱动电压仍然较高㊂相比之下,钛酸钡(BaTiO 3,BTO)的电光系数理论值可达到㊀第4期任怡静等:BaTiO 3薄膜的制备及其在电光调制器的应用689㊀1640pm /V,实验上获取的BTO 材料的最佳电光系数为923pm /V [4]㊂因此,由BTO 材料制作的电光调制器将具有更高的数据传输速率和更低的操作电压[5],使其有潜力应用于 超高速低电压 电光调制器[6-7]㊂国内之前关于BTO 的研究主要集中在薄膜的电学性能方面,探究了介电和铁电性能与衬底结构㊁晶体微观结构㊁温度变化和掺杂的密切关系,可以应用在铁电随机存取存储器等器件上[8]㊂之后研究热点为BTO 铁电薄膜畴壁表面电势分布的温度特性以及畴壁运动的基本机理[9-11],在外电场的作用下,晶畴所带的动能转化为势能后形成的内部电场直接影响电光效应,这对电光调制器的应用有着实际意义㊂最近,科研人员发现BTO 薄膜厚度和衬底也影响着薄膜的光学性能和薄膜质量[12-14],并论证了BTO 晶体薄膜电光调制器的100GHz 单通道调制带宽的可行性[15-19]㊂国内大部分研究停留在薄膜制备的层面,对BTO 电光调制器的系统性探索还比较少㊂总体而言,从基于BTO 晶体薄膜制作的电光调制器研究进展来看,我国与欧美等技术强国的工作基础还有一定的差距㊂本文从电光调制器工作原理出发,综述了BTO 薄膜的制备方法㊁工艺条件和薄膜衬底等对成膜的质量㊁生长取向和电光系数的影响,以及波导的结构对电光调制器性能的影响㊂1㊀电光调制器的工作原理电光调制器是利用某些电光晶体的电光效应制成的调制器㊂因为晶体折射率的各向异性与组成晶体的原子或分子的排列方式㊁相互作用有关,电压形成的电场影响到晶体中原子㊁分子的排列和它们之间的相互作用㊂这种内部变化,使晶体在宏观上表现出晶体折射率的改变,产生双折射效应㊂根据功能不同,可分为相位调制和强度调制㊂相位调制不改变光的偏振态,强度调制是通过改变晶体的折射率来改变光偏振态㊂根据加在晶体上电场的方向与光束在晶体中传播方向的不同,电光调制可分为纵向调制和横向调制㊂电场方向与光的传播方向平行,称为纵向电光调制;电场方向与光的传播方向垂直,称为横向电光调制[20]㊂同时拥有横向和强度调制特征的电光调制称为横向强度电光调制㊂图1为横向强度电光调制的过程,线偏振光垂直入射穿过电光晶体时,利用驱动电路向晶体提供一个电压信号,在外加电场的作用下晶体折射率发生变化,导致光波偏振状态发生变化,使得输入的线偏振光经过晶体后以椭圆偏振光输出㊂通过晶体后的出射光经过1/4波片,可以使电光调制器工作在线性区域㊂椭圆偏振光经过检偏器,恢复到线偏振光,完成光的强度调制㊂图1㊀横向电光强度调制原理图Fig.1㊀Schematic diagram of transverse electro-optic intensity modulation 2㊀薄膜制备工艺制备高质量的BTO 薄膜是获得高性能电光调制器的基础㊂BTO 单晶薄膜的制备方法和工艺条件直接影响着其生长取向㊁薄膜应力㊁表面粗糙度㊁晶粒尺寸㊁厚度均匀性和缺陷浓度,进而影响BTO 薄膜电光系数[21]㊂表1中总结了不同方法制备的BTO 薄膜质量参数和电光系数㊂BTO 薄膜的制备方法包括化学气相沉积(chemical vapor deposition,CVD)法㊁脉冲激光沉积(pulsed laser deposition,PLD)法㊁射频磁控溅射(radio frequency magnetron sputtering,RFMS)法㊁溶胶-凝胶(sol-gel)法和分子束外延(molecular beam epitaxy,MBE )法等㊂相对来说,实验室研究多采用CVD 法㊁PLD 法和MBE 法,而产业上多采用RFMS 法㊂690㊀综合评述人工晶体学报㊀㊀㊀㊀㊀㊀第52卷表1㊀不同方法制备的BTO薄膜质量参数和电光系数Table1㊀Quality and electro-optical coefficient of BTO films prepared by different methodsMethod Substrate a/c RefractiveindexBTOthickness/nm Roughness/nm OrientationElectro-opticcoefficient∗/(pm㊃V-1)ReferenceMOCVD MgO0.99 2.35700.23aγeff=150[22] MOCVD MgO 2.3/2.31620a/cγeff=360~420[23] PLD SOI+STO 1.01 2.2197 cγeff=37[21] MBE SOI+STO 17010 γeff=380[24] MBE SOI+STO 130<0.5aγeff=148[25] MBE SOI+STO 1.01 2.6800.4 γeff=213ʃ49[26] MBE SOI+STO 1.01 2.27780.5cγeff=140[21] MBE SOI+STO 50 aγeff=624[27]MBE SOI+STO 2250.4γ42=923ʃ215γ33=342ʃ93[4] RF-MS SOI+STO 7500.863cγeff=110[5] RF-MS SOI+STO 1.01 500 aγeff=131ʃ36[28] RF-MS SOI+STO 1.02 2.2196 cγeff=41[21]㊀㊀Note:SOI is silicon on insulator;STO is SrTiO3;∗γeff is effective electro-optic coefficient;γ42andγ33are components of the electro-optic tensor.2.1㊀化学气相沉积法该方法分为有机化学气相沉积法㊁金属化学气相沉积法和高真空化学气相沉积法等㊂在薄膜生长过程中,生长方式对成膜质量和速度有重要影响[29]㊂两步外延生长法分为低温成核和高温生长两个阶段,是当前化学气相沉积法的主流㊂美国西北大学Towner等[30]采用金属有机化学气相沉积法,通过低温生长和高温成核两步外延生长阶段,在MgO衬底上制备了厚度为100nm结晶度高㊁粗糙度低的BTO薄膜,优于他们以前文章报道的单阶段生长薄膜[31]㊂该方法通过第一步低温生长形成的氧化物缓冲层,能够使晶格更好地弛豫,有效降低界面处的热应力,从而改善了MgO与BTO之间5%晶格失配的问题㊂为了进一步改善晶格失配的问题,应当选择合适的衬底,如SiO2㊁Si㊁SrTiO3(STO)和绝缘衬底上硅(silicon on insulator,SOI)等㊂其中以STO为衬底的薄膜制备成为当前研究重点,有理论计算发现相对于MgO衬底,在STO衬底上制备的BTO薄膜能得到更高的电光系数[32]㊂此外,前驱体材料对制备高质量BTO薄膜起到关键性作用㊂Shuster等[33]采用金属化学气相沉积法, M(dfhd)2作为Ba和Sr的前驱体,在STO衬底上制备了1μm厚的高结晶度BTO-STO超晶格薄膜㊂相对于M(hfa)2[23],M(dfhd)2有更低的熔点㊁更好的热稳定性和存放稳定性,将其用作前驱体,可以获得结晶性更好的BTO薄膜㊂Reinke等[34]选用水㊁氧和臭氧为氧化剂,以异丙基环戊二烯基(Ba(iPr3Cp)2)㊁异丙氧化钛(TTIP)为前驱体,采用高真空化学气相沉积法在厚度为100nm Si衬底上制备了BTO薄膜㊂发现在370ħ低温下,使用H2O㊁O2作为氧化剂,可以得到Ba/Ti化学计量比为1ʒ1的BTO薄膜,其表面颗粒较大,相对光滑㊂然而,使用臭氧作氧化剂时,在BTO薄膜中看到BaCO3(002)XRD衍射峰㊂这可能是因为臭氧的氧化性强,使前驱体Ba(iPr3Cp)2被氧化进入BTO薄膜中,形成晶体BaCO3㊂所以,前驱体材料和配比等都会对薄膜的成分和结晶质量产生很大的影响㊂另一方面,高真空化学气相沉积不需要过高的生长温度㊂有研究表明,使用普通的化学气相沉积法时,薄膜在740ħ下生长结晶性能最好[32]㊂但是,高真空化学气相沉积法可以在370ħ生长温度下外延BTO薄膜,且不需要额外的退火步骤,显著降低了在互补金属氧化物半导体(complementory metal oxide semiconductor, CMOS)电路上获得功能BTO薄膜的总热预算㊂Reinke等[35]使用高真空化学气相沉积法探讨了前驱体撞击速率对BTO薄膜形貌和生长取向的影响㊂前驱体在高撞击速率下制备的BTO薄膜形貌呈现多晶性质和小角度晶界,这种形貌的形成可能是STO层为缓解不同结构之间的晶格应变,形成了缺陷㊂另外,BTO薄膜在高生长速率沉积时是a取向的,而低生长速率沉积时既有a取向又有c取向㊂可见,虽然前驱体撞击速率较高时会使BTO薄膜呈现高占比a轴取向,但形貌出现缺陷㊂该研究进一步探讨了实验条件对BTO薄膜生长的影响,为后续的研究提供了支撑㊂㊀第4期任怡静等:BaTiO3薄膜的制备及其在电光调制器的应用691㊀化学气相沉积法成膜装置简单㊁制备过程灵活性较大,但沉积薄膜的速率不高,低于溅射镀膜,因此应用上受到一定限制㊂可见,采用化学气相沉积法制备薄膜的工艺不适用于产业化㊂2.2㊀脉冲激光沉积法在该方法中,激光脉冲和衬底温度同时影响了薄膜的生长取向㊂Petraru等[36-37]采用脉冲激光沉积法,在MgO衬底上沉积了多晶BTO薄膜,探究了激光功率和衬底温度对薄膜生长取向和电光系数的影响㊂当激光功率为1100mJ/脉冲,衬底温度为800ħ时,制备出c轴取向的薄膜,在波长为632nm处,有效电光系数为86pm/V;当激光功率为300mJ/脉冲,衬底温度为850ħ时,制备出a轴取向的薄膜,在波长为632nm 处,得到较大的有效电光系数734pm/V㊂该研究发现了使用脉冲激光沉积法时适合BTO薄膜生长的激光功率和衬底温度,并且得到了较为理想的电光系数,但是没有深入研究脉冲激光沉积法的各项参数如何单独影响薄膜质量和性能㊂提高衬底温度有利于晶粒均匀生长,改善结晶度,进而影响BTO的晶格参数[38]㊂2018年,Lyu等[39]在LaNiO3㊁CeO2㊁YSZ缓冲层和Si衬底上,制备了厚度为110nm的BTO薄膜㊂薄膜在所有衬底温度(T s)下都呈现横向岛状生长结构,如图2(a)~(c)所示㊂衬底温度为375ħ时,表面形态较均匀㊂衬底温度达到400ħ时,薄膜的表面不平整,粗糙度上升㊂700ħ的衬底温度下,岛屿较高,结构密集㊂这是因为在375ħ衬底温度下,薄膜刚开始沉积,处于非结晶状态㊂达到400ħ时,薄膜处于结晶良好与非结晶态区域共存的形态,导致岛屿高度差异很大㊂700ħ时,形成结晶良好的薄膜㊂不同温度下结晶度对比如图2(d)所示㊂在深入研究了衬底温度对BTO薄膜生长的影响后,发现最佳的衬底温度为700ħ,并得到较为平整的薄膜㊂图2㊀不同衬底温度对BTO薄膜生长的影响㊂(a)~(c)BTO薄膜在不同衬底温度下生长的形貌图;(d)BTO薄膜在不同衬底温度下生长的XRD图谱Fig.2㊀Effect of different substrate temperatures on BTO films growth.(a)~(c)Morphology of BTO film growing at different substrate temperatures;(d)XRD patterns of BTO film growing at different substrate temperatures2022年,Behera等[40]在Si和SiO2衬底上制备了400nm厚的BTO薄膜,发现了与Lyu等[39]相似的薄膜形貌和结晶规律㊂并进一步发现在700ħ衬底温度下,薄膜的化学计量比最接近理想值,薄膜缺陷最小㊂另外,衬底温度变化时,a轴晶格参数变化的幅度很小,而c轴晶格参数变化的幅度很大,说明衬底温度对c轴晶格参数的影响更大㊂所以可以利用衬底温度来操纵a㊁c轴晶格参数和铁电极化,从而得到高电光系数的BTO薄膜㊂在较低的衬底温度下,原子的迁移率较低,致使衬底表面扩散能力太低而不能成核,这样得到的薄膜表面比较粗糙㊂而衬底温度过高时,表面粒子迁移率过高,衬底表面原子结合能力下降,从而引入大量缺陷㊂合适的衬底温度下,衬底表面粒子有足够的能量迁移到易于成核的位置,易凝聚成核,因此薄膜的取向一致,结晶度高㊂在脉冲激光沉积法中,衬底温度是一个重要的实验参数㊂除衬底温度外,氧分压也是决定薄膜的粗糙度㊁晶粒尺寸和生长取向的关键因素㊂a轴取向薄膜的光轴垂直于外加电场,电光效应依赖于电光张量中较大的垂直γ51分量,所以a轴取向占比越高,薄膜的电光效应越强[41]㊂在低氧分压环境下易得到c轴取向的薄膜,并拥有高度双折射效应;在高氧分压环境下易得到a 轴取向的薄膜,但粗糙度较高[42]㊂Estrada等[43]采用脉冲激光沉积法,在STO衬底上制备了BTO薄膜㊂发现氧分压过小,薄膜粗糙度很大,随着氧分压的增加,晶粒尺寸增大,样品的粗糙度也减小㊂但氧分压增加到692㊀综合评述人工晶体学报㊀㊀㊀㊀㊀㊀第52卷0.1mbar(1mbar=0.1kPa)之后,粗糙度随着氧分压继续增加而增加㊂Lyu等[44]采用脉冲激光沉积法,在Si 和钙钛矿衬底上制备了厚度为94~112nm的BTO薄膜㊂氧分压过高时,薄膜粗糙度变高,并且生长取向由c轴向a轴转变㊂Wang等[45]在Si衬底上制备了BTO薄膜,探讨了氧分压对BTO极化方向和晶格常数大小的影响,得到了与Lyu等[44]相同结论㊂该研究阐明了BTO生长方向随氧分压变化的机理,实现了氧分压作为调节旋钮控制BTO薄膜的极化方向的目的㊂所以,过低或者过高的氧分压都会对薄膜的生长产生不利的影响㊂这是因为氧分压的增大会使沉积速率减小,原子从而有足够的时间在衬底表面迁移,有利于核的形成和晶粒尺寸的生长,有效减少缺陷㊂但若氧分压过大,会迅速减慢粒子向衬底过渡,不利于薄膜的制备㊂在脉冲激光沉积法中,合适的氧分压才能得到质量高且a轴取向的膜㊂BTO本身是各向异性晶体,存在双折射效应,在BTO薄膜波导的电光调制理论中,对晶体薄膜的电光特性产生很大影响㊂罗梦希[46]在MgO衬底上制备了a轴和c轴生长方向的BTO薄膜,建立了γ51和双折射率b eo的同步测试与分析模型,探究了电光系数γ51和双折射b eo与调制电场间的有效关系,得到的测量误差在ʃ5%以内㊂以往的很多研究忽略了双折射的影响,导致其实验结果不准确,而将BTO本身双折射效应的影响考虑进去会使γ51的测量更加精确㊂脉冲激光沉积法优点是沉积速率高㊁可以原位引入多种气体,以及提高薄膜的质量等㊂并且该方法对腔体内的大气环境要求低,可以制备出许多与靶材成分接近的薄膜[47]㊂但是此方法制备的薄膜厚度不均匀,且易形成缺陷,即薄膜存在表面颗粒问题,大面积地均匀沉淀比较困难㊂2.3㊀磁控溅射法磁控溅射法分为直流磁控溅射法和射频磁控溅射法,对于绝缘靶材或导电性差的非金属靶材,需使用射频磁控溅射法[48]㊂为了制备高质量薄膜,应当根据材料性质和技术指标来探究最合适的溅射参数,其中包括射频功率㊂射频功率过低,将会导致BTO在基底表面的入射能量不够,使BTO薄膜结晶困难㊂BTO XRD 衍射峰的强度会随着射频功率增加而增加,在射频功率到达一定高度时获得最大强度,但是若射频功率继续增加,会形成富含Ti的第二相BaTi2O5,降低结晶度[49]㊂工作气压影响着薄膜的结晶性能㊂BTO XRD衍射峰的强度会随着工作气压的增加而增加,在一定工作气压处达到最大后,又随着工作气压的进一步增加而减小㊂这是因为在合适的工作气压下,高能量粒子垂直轰击靶材,溅射粒子到达衬底时不会造成显著的能量损失,形成高结晶薄膜㊂而在较高的工作气压下,粒子散射引起低能粒子以斜角轰击,BTO薄膜不再是定向生长,而是随机生长[48]㊂溅射气体的组成也是关键溅射参数之一㊂相对于纯氩气的环境,在20%(体积分数,下同)氧气和80%氩气的混合气体中沉积时,BTO薄膜具有更大的残余极化和胁迫电压㊂因为纯氩气中沉积的BTO薄膜存在缺氧问题,导致泄漏电流增加,铁电性能下降㊂在混合气体环境沉积薄膜的过程中,氧气减少氧空位数量,从而提高BTO薄膜的性能㊂在10%氧分压条件下沉积的BTO薄膜具有a轴和c轴混合取向;20%氧分压下沉积的薄膜呈现高度c轴取向㊂在氧分压达到30%时,BTO相变为BaTi2O5[49]㊂虽然BaTi2O5的形成会降低薄膜结晶度,改变介电常数,但是氧会填充柱状BTO晶体间的孔隙,因此BaTi2O5的形成有利于降低薄膜孔隙率,改善薄膜的电学性能[50]㊂薄膜厚度与BTO薄膜的粗糙度有关㊂但Kim等[5]发现薄膜厚度在达到1μm时,粗糙度变高,结晶度降低㊂750nm厚和1μm厚的BTO薄膜表面粗糙度的均方根分别为0.863nm和5.16nm㊂2021年,德克萨斯大学奥斯汀分校Posadas等[28]以MgO为衬底,采用离轴射频磁控溅射法制备了500nm与1μm厚的BTO 薄膜,同样得到1μm厚的薄膜质量较差的结果㊂但是在700ħ高温结晶后,1μm厚的BTO薄膜的粗糙度得到了极大的改善㊂衬底对薄膜的生长取向㊁结晶度影响较大㊂近年来硅衬底电光调制器拥有的各种优良性能,使其成为开发光电集成芯片最有吸引力的研究方向之一㊂在SOI衬底上制备BTO薄膜成为研究热点,BTO在没有界面反应的情况下,不能直接在SOI上成核,所以一般在BTO与SOI之间使用缓冲层,STO作为缓冲层沉积的BTO 薄膜结晶度最好[50]㊂另外,衬底与薄膜的自由载流子浓度㊁内置电压和泄漏电流都有密切的联系[18-19]㊂缓冲层的厚度也与BTO薄膜的生长取向有着密切关系㊂Posadas等[28]采用射频溅射法,在SOI衬底上制备了厚度2㊁4和6nm的STO缓冲层以及300nm厚的薄膜㊂结果表明,厚度6nm STO上的BTO薄膜外延㊀第4期任怡静等:BaTiO 3薄膜的制备及其在电光调制器的应用693㊀表面平坦,如图3(a)所示㊂厚度4nm STO 上的BTO 薄膜生长稳定,但是出现了岛状生长的情况,如图3(b)所示㊂厚度2nm 的STO 层被溅射生长过程中相对恶劣的高温和高氧分压破坏,使其上面生长的BTO 薄膜有大量多晶组分,如图3(c)所示㊂缓冲层越厚BTO 薄膜质量越高,然而由图3(d)~(f)可以看出,缓冲层越薄,越有利于a 轴取向的BTO 薄膜生长㊂所以,为了得到较平坦且高a 轴取向的BTO 薄膜,4nm 的STO 缓冲层是最好的选择㊂若使用射频溅射法,通过非晶生长和再结晶两步结晶,可得到100%沿a 轴生长的薄膜,但其有大量空隙㊂而直接高温溅射可以产生无孔薄膜,提高薄膜的结晶度,从而较少泄漏电流,该方法解决了在较低薄膜厚度的情况下,100%沿a 轴生长薄膜的问题,并探究出了最适合BTO 生长的缓冲层厚度㊂图3㊀STO 模板层厚度对在700ħ下生长的厚度为300nm BTO 质量的影响[28]㊂(a)~(c)BTO 薄膜在6㊁4和2nm 厚STO 层上生长的RHEED 图;(d)~(f)BTO 薄膜在6㊁4和2nm 厚STO 层上生长的XRD 图谱Fig.3㊀Effect of STO template layer thickness on quality of 300nm BTO grown at 700ħ[28].(a)~(c)RHEED images after growth of BTO films on 6,4and 2nm STO layers;(d)~(f)XRD patterns of BTO films grown on 6,4and 2nm STO layers 溅射法的优势在于薄膜沉积时工艺稳定㊁易重复,以及在中低温就可以生长出结晶性良好的薄膜,在晶体质量上可与分子束外延生长的薄膜相媲美,且生长速度可快10倍,更加适合产业化㊂但是此方法有设备维修和靶材成本高㊁等离子体不稳定等缺点㊂2.4㊀溶胶凝胶法水热法可以在低温条件下结晶,并且结晶性良好,结合水热法和溶胶凝胶法可得到一种被称为溶胶凝胶水热法(sol-gel hydrothermal,SGHM)的新方法㊂相对于普通的溶胶凝胶法[51],SGHM 明显地降低了结晶温度,提高了BTO 的结晶度和纯度[52]㊂退火工艺与薄膜的结晶度和质量有密切的联系㊂Chinchamalatpure 等[53]研究发现,在600ħ的退火温度下,BTO 膜呈现四方相,薄膜的结晶度增强㊂室温C -V 特性表明,积累区有较大的频率色散,并且理想因子接近统一㊂2020年,Edmondson 等[54]在SOI 衬底STO 缓冲层上制备了85nm 厚的BTO 薄膜㊂在600ħ的温度下退火1h 后,同样得到呈四方相且结晶度较好的BTO 薄膜㊂对退火前后BTO 薄膜的电光系数进行对比,结果表明晶粒尺寸的大小和薄膜缺陷密度的高低㊁有效电光系数的大小有密切联系㊂之后,该研究组采用另一种BTO /Si 的退火程序,即在750ħ温度和流动氧气的环境中退火10h,可以使退火后BTO 薄膜的电光系数增加3到4倍[55]㊂这是因为高温退火可以使晶体结构重组,改变晶粒表面松弛的状态,使薄膜的结晶度更高㊂溶胶凝胶法有设备易于操作㊁实验成本低㊁成膜效率高,以及可大面积成膜等优势,已经在薄膜制备方面被广泛应用㊂但通过溶胶凝胶法制备的薄膜均匀性比较差,控制溶胶的表面张力可以解决这一问题㊂2.5㊀分子束外延法除了提高晶体薄膜的结晶性能㊁降低粗糙度和选取特定生长取向外,抑制薄膜缺陷也可以提高薄膜的电694㊀综合评述人工晶体学报㊀㊀㊀㊀㊀㊀第52卷光性能㊂Abel [56]在STO 缓冲层和SOI 衬底上沉积了50nm 的BTO 薄膜,为了使Ba 和Ti 之间的化学计量比达到1ʒ1以抑制BTO 薄膜中的缺陷,开发了封闭共沉积生长技术㊂该技术的要点在于通过调节通入Ba 和Ti 的时间间隔Δt 来控制Ba 和Ti 的比例,并建立反馈机制进行补偿㊂该研究通过逐层封闭共沉积生长方法,突破了在实验中无法精确控制化学计量的局限性㊂使用分子束外延法可以制备出高质量的超晶格结构㊂最早,日本科学家Tsurumi 等[57]首先研制了BTO-STO 超晶格薄膜,并在不同的周期和原胞层数下,研究了超晶格的结构和介电性能㊂在超晶格周期叠加过程中,电光效应随着叠加周期的降低而增加,达到最大值后,再次降低叠加周期就会导致有效电光系数降低㊂2019年,Merckling 等[58]在STO 缓冲层和Si 衬底上,制备了5个周期高质量㊁无位错的BTO-STO 超晶格㊂通过纳米电子衍射分析实验,得到STO /Si 衬底上弛豫的临界厚度为5nm,这是因为使用由堆叠[BTO /STO]双分子层膜组成的应变介导超晶格可防止错配位错的形成㊂超晶格结构能够提高薄膜的电光性能,增大折射率,以及增强非线性光学效应,还可以通过改变其周期来改变材料的性质㊂在使用分子束外延法制备薄膜时,后处理工艺极大地影响了其电光性能㊂Hsu 等[41]发现在分子束外延法生长BTO 薄膜的过程中,BTO 层被引入了许多氧空位,导致了较大的泄漏电流㊂在600ħ的氧环境下进行30min 的后退火,可有效抑制泄漏电流,以确保器件运行过程中在BTO 层中建立足够的电场㊂Abel 等[4]采用STO 为缓冲层,在Si 衬底上制备了225nm 厚的BTO 薄膜㊂用Al 2O 3作为结合界面,将BTO 层转移到SOI 上,两种晶片的表面粗糙度均低于0.4nm㊂在低损耗的热处理后,通过蚀刻移除施主晶圆,得到BTO 薄膜的电光系数为923pm /V,是较为理想的值㊂该研究通过在BTO 薄膜与衬底之间添加氧化层避免光泄漏,极大地改善了薄膜的电光性能,得到了BTO 在实验上最大的电光系数㊂综上所述,分子束外延法相比于其他方法能够得到较高的电光系数,是因为采用MBE 法得到的薄膜有较低的孔隙率,更容易制备出四方的㊁致密的和结晶度高的薄膜㊂并且通过这种方法可以准确地控制薄膜的微量成分㊁厚度,以及掺杂量,还可以原位观测外延薄膜的具体情况,但是用此方法制备薄膜时生长速度较慢,对衬底要求高,不适用于产业化生产㊂2.6㊀Smart Cut 技术智能剥离技术可以制备均匀㊁较薄的SOI 衬底㊂它解决了硅片键合和反面腐蚀技术中减薄困难问题的同时,兼顾了注氧隔离技术和键合技术的优点,但该技术制备出的薄膜较厚㊁粗糙度较大㊂针对具有更大惰性和易损性的金属氧化物,研究人员结合传统Smart Cut 开发出一种获得金属氧化物单晶薄膜的技术,即离子注入剥离(crystal ion slicing,CIS),可降低薄膜的粗糙度,主要步骤如图4所示㊂Izuhara 等[59]使用CIS 制备出厚度8μm 的薄膜,首先,将He 2+注入BTO 材料中㊂该步骤能够在材料下方引入缺陷层,将薄膜与供体晶片分离,膜厚取决于注入能量,0.26㊁1.2和2MeV 的植入能量分别可制备出厚度0.5㊁4.7和8μm 的薄膜㊂注入时,环境温度保持在低温50ħ,避免缺陷层的动态退火㊂在第二步中,将植入的BTO 晶圆粘合在覆盖粘合层的衬底之上㊂接下来,样品逐渐加热至220ħ㊂热致应力使牺牲层能够从供体晶圆上剥离薄膜,之后在300ħ的温度下退火数小时㊂最后,通过Ar +离子蚀刻使薄膜的表面光滑㊂应用机械抛光可将其表面粗糙度降低到1nm 以下,制备X -CUT 和Z -CUT 的薄膜,应用于不同的电光器件㊂图4㊀CIS 制备薄膜的步骤[60]Fig.4㊀The steps of preparing thin film by CIS [60]。
大学物理实验报告 英文版
大学物理实验报告Ferroelectric Control of Spin PolarizationABSTRACTA current drawback of spintronics is the large power that is usually required for magnetic writing, in contrast with nanoelectronics, which relies on “zero-current,” gate-controlled operations. Efforts have been made to control the spin-relaxation rate, the Curie temperature, or the magnetic anisotropy with a gate voltage, but these effects are usually small and volatile. We used ferroelectric tunnel junctions with ferromagnetic electrodes to demonstrate local, large, and nonvolatile control of carrier spin polarization by electrically switching ferroelectric polarization. Our results represent a giant type of interfacial magnetoelectric coupling and suggest a low-power approach for spin-based information control.Controlling the spin degree of freedom by purely electrical means is currently an important challenge in spintronics (1, 2). Approaches based on spin-transfer torque (3) have proven very successful in controlling the direction of magnetization in a ferromagnetic layer, but they require the injection of high current densities. An ideal solution would rely on the application of an electric field across an insulator, as in existing nanoelectronics. Early experiments have demonstrated the volatile modulation of spin-based properties with a gate voltage applied through a dielectric. Notable examples include the gate control of the spin-orbit interaction in III-V quantum wells (4), the Curie temperature T C (5), or the magnetic anisotropy (6) in magnetic semiconductors with carrier-mediated exchange interactions; for example, (Ga,Mn)As or (In,Mn)As. Electric field–induced modifications of magnetic anisotropy at room temperature have also been reported recently in ultrathin Fe-based layers (7, 8).A nonvolatile extension of this approach involves replacing the gate dielectric by a ferroelectric and taking advantage of the hysteretic response of its order parameter (polarization) with an electric field. When combined with (Ga,Mn)As channels, forinstance, a remanent control of T C over a few kelvin was achieved through polarization-driven charge depletion/accumulation (9, 10), and the magnetic anisotropy was modified by the coupling of piezoelectricity and magnetostriction (11, 12). Indications of an electrical control of magnetization have also been provided in magnetoelectric heterostructures at room temperature (13–17).Recently, several theoretical studies have predicted that large variations of magnetic properties may occur at interfaces between ferroelectrics and high-T C ferromagnets such as Fe (18–20), Co2MnSi (21), or Fe3O4 (22). Changing the direction of the ferroelectric polarization has been predicted to influence not only the interfacial anisotropy and magnetization, but also the spin polarization. Spin polarization [i.e., the normalized difference in the density of states (DOS) of majority and minority spin carriers at the Fermi level (E F)] is typically the key parameter controlling the response of spintronics systems, epitomized by magnetic tunnel junctions in which the tunnel magnetoresistance (TMR) is related to the electrode spin polarization by the Jullière formula (23). These predictions suggest that the nonvolatile character of ferroelectrics at the heart of ferroelectric random access memory technology (24) may be exploited in spintronics devices such as magnetic random access memories or spin field-effect transistors (2). However, the nonvolatile electrical control of spin polarization has not yet been demonstrated.We address this issue experimentally by probing the spin polarization of electrons tunneling from an Fe electrode through ultrathin ferroelectric BaTiO3 (BTO) tunnel barriers (Fig. 1A). The BTO polarization can be electrically switched to point toward oraway from the Fe electrode. We used a half-metallic La0.67Sr0.33MnO3(LSMO) (25) bottom electrode as a spin detector in these artificial multiferroic tunnel junctions (26, 27). Magnetotransport experiments provide evidence for a large and reversible dependence of the TMR on ferroelectric polarization direction.Fig. 1(A) Sketch of the nanojunction defined by electrically controlled nanoindentation. A thin resist is spin-coated on the BTO(1 nm)/LSMO(30 nm) bilayer. The nanoindentation is performed with a conductive-tip atomic force microscope, and the resultingnano-hole is filled by sputter-depositing Au/CoO/Co/Fe. (B) (Top) PFM phase image of a BTO(1 nm)/LSMO(30 nm) bilayer after poling the BTO along 1-by-4–μm stripes with either a negative or positive (tip-LSMO) voltage. (Bottom) CTAFM image of an unpoled area of a BTO(1 nm)/LSMO(30 nm) bilayer. Ω, ohms. (C) X-ray absorption spectra collected at room temperature close to the Fe L3,2 (top), Ba M5,4 (middle), and TiL3,2 (bottom) edges on an AlO x(1.5 nm)/Al(1.5 nm)/Fe(2 nm)/BTO(1 nm)/LSMO(30 nm)//NGO(001) heterostructure. (D) HRTEM and (E) HAADF images of the Fe/BTO interface in a Ta(5 nm)/Fe(18 nm)/BTO(50 nm)/LSMO(30 nm)//NGO(001) heterostructure. The white arrowheads in (D) indicate the lattice fringes of {011} planes in the iron layer. [110] and [001] indicate pseudotetragonal crystallographic axes of the BTO perovskite.The tunnel junctions that we used in this study are based on BTO(1 nm)/LSMO(30 nm) bilayers grown epitaxially onto (001)-oriented NdGaO3 (NGO) single-crystal substrates (28). The large (~180°) and stable piezoresponse force microscopy (PFM) phase contrast (28) between negatively and positively poled areas (Fig. 1B, top) indicates that the ultrathin BTO films are ferroelectric at room temperature (29). The persistence of ferroelectricity for such ultrathin films of BTO arises from the large lattice mismatch with the NGO substrate (–3.2%), which is expected to dramatically enhance ferroelectric properties in this highly strained BTO (30). The local topographical and transport properties of the BTO(1 nm)/LSMO(30 nm) bilayers were characterized by conductive-tip atomic force microscopy (CTAFM) (28). The surface is very smooth with terraces separated by one-unit-cell–high steps, visible in both the topography (29) and resistance mappings (Fig. 1B, bottom). No anomalies in the CTAFM data were observed over lateral distances on the micrometer scale.We defined tunnel junctions from these bilayers by a lithographic technique based on CTAFM (28, 31). Top electrical contacts of diameter ~10 to 30 nm can be patterned by this nanofabrication process. The subsequent sputter deposition of a 5-nm-thick Fe layer, capped by a Au(100 nm)/CoO(3.5 nm)/Co(11.5 nm) stack to increase coercivity, defined a set of nanojunctions (Fig. 1A). The same Au/CoO/Co/Fe stack was deposited on another BTO(1 nm)/LSMO(30 nm) sample for magnetic measurements. Additionally, a Ta(5 nm)/Fe(18 nm)/BTO(50 nm)/LSMO(30 nm) sample and a AlO x(1.5 nm)/Al(1.5nm)/Fe(2 nm)/BTO(1 nm)/LSMO(30 nm) sample were realized for structural and spectroscopic characterizations.We used both a conventional high-resolution transmission electron microscope (HRTEM) and the NION UltraSTEM 100 scanning transmission electron microscope (STEM) to investigate the Fe/BTO interface properties of the Ta/Fe/BTO/LSMO sample. The epitaxial growth of the BTO/LSMO bilayer on the NGO substrate was confirmed by HRTEM and high-resolution STEM images. The low-resolution, high-angle annular dark field (HAADF) image of the entire heterostructure shows the sharpness of theLSMO/BTO interface over the studied area (Fig. 1E, top). Figure 1D reveals a smooth interface between the BTO and the Fe layers. Whereas the BTO film is epitaxially grown on top of LSMO, the Fe layer consists of textured nanocrystallites. From the in-plane (a) and out-of-plane (c) lattice parameters in the tetragonal BTO layer, we infer that c/a = 1.016 ± 0.008, in good agreement with the value of 1.013 found with the use of x-ray diffraction (29). The interplanar distances for selected crystallites in the Fe layer [i.e.,~2.03 Å (Fig. 1D, white arrowheads)] are consistent with the {011} planes ofbody-centered cubic (bcc) Fe.We investigated the BTO/Fe interface region more closely in the HAADF mode of the STEM (Fig. 1E, bottom). On the BTO side, the atomically resolved HAADF image allows the distinction of atomic columns where the perovskite A-site atoms (Ba) appear as brighter spots. Lattice fringes with the characteristic {100} interplanar distances of bcc Fe (~2.86 Å) can be distinguished on the opposite side. Subtle structural, chemical, and/or electronic modifications may be expected to occur at the interfacial boundarybetween the BTO perovskite-type structure and the Fe layer. These effects may lead to interdiffusion of Fe, Ba, and O atoms over less than 1 nm, or the local modification of the Fe DOS close to E F, consistent with ab initio calculations of the BTO/Fe interface (18–20).To characterize the oxidation state of Fe, we performed x-ray absorption spectroscopy (XAS) measurements on a AlO x(1.5 nm)/Al(1.5 nm)/Fe(2 nm)/BTO(1 nm)/LSMO(30 nm) sample (28). The probe depth was at least 7 nm, as indicated by the finite XAS intensity at the La M4,5 edge (28), so that the entire Fe thickness contributed substantially to the signal. As shown in Fig. 1C (top), the spectrum at the Fe L2,3 edge corresponds to that of metallic Fe (32). The XAS spectrum obtained at the Ba M4,5 edge (Fig. 1C, middle) is similar to that reported for Ba2+ in (33). Despite the poor signal-to-noise ratio, the Ti L2,3 edge spectrum (Fig. C, bottom) shows the typical signature expected for a valence close to 4+ (34). From the XAS, HRTEM, and STEM analyses, we conclude that theFe/BTO interface is smooth with no detectable oxidation of the Fe layer within a limit of less than 1 nm.After cooling in a magnetic field of 5 kOe aligned along the [110] easy axis of pseudocubic LSMO (which is parallel to the orthorhombic [100] axis of NGO), we characterized the transport properties of the junctions at low temperature (4.2K). Figure 2A (middle) shows a typical resistance–versus–magnetic field R(H) cycle recorded at a bias voltage of –2 mV (positive bias corresponds to electrons tunneling from Fe to LSMO). The bottom panel of Fig. 2A shows the magnetic hysteresisloop m(H) of a similar unpatterned sample measured with superconducting quantuminterference device (SQUID) magnetometry. When we decreased the magnetic field from a large positive value, the resistance dropped in the –50 to –250 Oe range and then followed a plateau down to –800 Oe, after which it sharply returned to thehigh-resistance state. We observed a similar response when cycling the field back to large positive values. A comparison with the m(H) loop indicates that the switching fields in R(H) correspond to changes in the relative magnetic configuration of the LSMO and Fe electrodes from parallel (at high field) to antiparallel (at low field). The magnetically softer LSMO layer switched at lower fields (50 to 250 Oe) compared with the Fe layer, for which coupling to the exchange-biased Co/CoO induces larger and asymmetric coercive fields (–800 Oe, 300 Oe). The observed R(H) corresponds to a negative TMR = (R ap–R p)/R ap of –17% [R p and R ap are the resistance in the parallel (p) and antiparallel (ap) magnetic configurations, respectively; see the sketches in Fig. 2A]. Within the simple Jullière model of TMR (23) and considering the large positive spin polarization of half-metallic LSMO (25), this negative TMR corresponds to a negative spin polarization for bcc Fe at the interface with BTO, in agreement with ab initio calculations (18–20).Fig. 2(A) (Top) Device schematic with black arrows to indicate magnetizations. p, parallel; ap, antiparallel. (Middle) R(H) recorded at –2 mV and 4.2 K showing negative TMR. (Bottom) m(H) recorded at 30 K with a SQUID magnetometer. emu, electromagnetic units. (B) (Top) Device schematic with arrows to indicate ferroelectric polarization. (Bottom) I(V DC) curves recorded at 4.2 K after poling the ferroelectric down (orange curve) or up (brown curve). The bias dependence of the TER is shown in the inset.As predicted (35–38) and demonstrated (29) previously, the tunnel current across a ferroelectric barrier depends on the direction of the ferroelectric polarization. We also observed this effect in our Fe/BTO/LSMO junctions. As can be seen in Fig. 2B, after poling the BTO at 4.2 K to orient its polarization toward LSMO or Fe (with a poling voltage of VP–≈ –1 V or VP+≈ 1 V, respectively; see Fig. 2B sketches),current-versus-voltage I(V DC) curves collected at low bias voltages showed a finite difference corresponding to a tunnel electroresistance as large as TER = (I VP+–I VP–)/I VP–≈ 37% (Fig. 2B, inset). This TER can be interpreted within an electrostatic model (36–39), taking into account the asymmetric deformation of the barrier potential profile that is created by the incomplete screening of polarization charges by different Thomas-Fermi screening lengths at Fe/BTO and LSMO/BTO interfaces.Piezoelectric-related TER effects (35, 38) can be neglected as the piezoelectric coefficient estimated from PFM experiments is too small in our clamped films (29). TER measurements performed on a BTO(1 nm)/LSMO(30 nm) bilayer with the use of a CTAFM boron-doped diamond tip as the top electrode showed values of ~200%(29). Given the strong sensitivity of the TER on barrier parameters and barrier-electrode interfaces, these two values are not expected to match precisely. We anticipate that the TER variation between Fe/BTO/LSMO junctions and CTAFM-based measurements is primarily the result of different electrostatic boundary conditions.Switching the ferroelectric polarization of a tunnel barrier with voltage pulses is also expected to affect the spin-dependent DOS of electrodes at a ferromagnet/ferroelectric interface. Interfacial modifications of the spin-dependent DOS of the half-metallic LSMO by the ferroelectric BTO are not likely, as no states are present for the minority spins up to ~350 meV above E F (40, 41). For 3d ferromagnets such as Fe, large modifications of the spin-dependent DOS are expected, as charge transfer between spin-polarized empty and filled states is possible. For the Fe/BTO interface, large changes have been predicted through ab initio calculations of 3d electronic states of bcc Fe at the interface with BTO by several groups (18–20).To experimentally probe possible changes in the spin polarization of the Fe/BTO interface, we measured R(H) at a fixed bias voltage of –50 mV after aligning the ferroelectric polarization of BTO toward Fe or LSMO. R(H) cycles were collected for each direction of the ferroelectric polarization for two typical tunnel junctions of the same sample (Fig. 3, B and C, for junction #1; Fig. 3, D and E, for junction #2). In both junctions at the saturating magnetic field, high- and low-resistance states are observed when the ferroelectric polarization points toward LSMO or Fe, respectively, with a variation of ~ 25%. This result confirms the TER observations in Fig. 2B.Fig. 3(A) Sketch of the electrical control of spin polarization at the Fe/BTO interface.(B and C) R(H) curves for junction #1 (V DC = –50 mV, T = 4.2 K) after poling the ferroelectric barrier down or up, respectively. (D and E) R(H) curves for junction #2 (V DC = –50 mV, T= 4.2 K) after poling the ferroelectric barrier down or up, respectively.More interestingly, here, the TMR is dramatically modified by the reversal of BTO polarization. For junction #1, the TMR amplitude changes from –17 to –3% when the ferroelectric polarization is aligned toward Fe or LSMO, respectively (Fig. 3, B and C). Similarly for junction #2, the TMR changes from –45 to –19%. Similar results were obtained on Fe/BTO (1.2 nm)/LSMO junctions (28). Within the Jullière model (23), these changes in TMR correspond to a large (or small) spin polarization at the Fe/BTO interface when the ferroelectric polarization of BTO points toward (or away from) the Fe electrode. These experimental data support our interpretation regarding the electrical manipulation of the spin polarization of the Fe/BTO interface by switching the ferroelectric polarization of the tunnel barrier.To quantify the sensitivity of the TMR with the ferroelectric polarization, we define a term, the tunnel electromagnetoresistance, as TEMR = (TMR VP+–TMR VP–)/TMR VP–. Largevalues for the TEMR are found for junctions #1 (450%) and #2 (140%), respectively. This electrical control of the TMR with the ferroelectric polarization is repeatable, as shown in Fig. 4 for junction #1 where TMR curves are recorded after poling the ferroelectric up, down, up, and down, sequentially (28).Fig. 4TMR(H) curves recorded for junction #1 (V DC = –50 mV, T = 4.2 K) after poling the ferroelectric up (VP+), down (VP–), up (VP+), and down (VP–).For tunnel junctions with a ferroelectric barrier and dissimilar ferromagnetic electrodes, we have reported the influence of the electrically controlled ferroelectric barrier polarization on the tunnel-current spin polarization. This electrical influence over magnetic degrees of freedom represents a new and interfacial magnetoelectric effect that is large because spin-dependent tunneling is very sensitive to interfacial details. Ferroelectrics can provide a local, reversible, nonvolatile, and potentially low-power means of electrically addressing spintronics devices.Supporting Online Material/cgi/content/full/science.1184028/DC1Materials and MethodsFigs. S1 to S5References∙Received for publication 30 October 2009.∙Accepted for publication 4 January 2010.References and Notes1. C. Chappert, A. Fert, F. N. Van Dau, The emergence of spin electronics indata storage. Nat. Mater. 6,813 (2007).2.I. Žutić, J. Fabian, S. Das Sarma, Spintronics: Fundamentals andapplications. Rev. Mod. Phys. 76,323 (2004).3.J. C. Slonczewski, Current-driven excitation of magnetic multilayers. J.Magn. Magn. Mater. 159, L1(1996).4.J. Nitta, T. Akazaki, H. Takayanagi, T. Enoki, Gate control of spin-orbit interaction in an inverted In0.53Ga0.47As/In0.52Al0.48Asheterostructure. Phys. Rev. Lett. 78, 1335 (1997).5.H. Ohno et al., Electric-field control offerromagnetism. Nature 408, 944 (2000).6. D. Chiba et al., Magnetization vector manipulation by electricfields. Nature 455, 515 (2008).7.M. Weisheit et al., Electric field–induced modification of magnetism inthin-film ferromagnets. Science315, 349 (2007).8.T. Maruyama et al., Large voltage-induced magnetic anisotropy changein a few atomic layers of iron.Nat. Nanotechnol. 4, 158 2009).9.S. W. E. Riester et al., Toward a low-voltage multiferroic transistor:Magnetic (Ga,Mn)As under ferroelectric control. Appl. Phys.Lett. 94, 063504 (2009).10.I. Stolichnov et al., Non-volatile ferroelectric control of ferromagnetismin (Ga,Mn)As. Nat. Mater. 7, 464(2008).11. C. Bihler et al., Ga1−x Mn x As/piezoelectric actuator hybrids: A modelsystem for magnetoelastic magnetization manipulation. Phys. Rev.B 78, 045203 (2008).12.M. Overby, A. Chernyshov, L. P. Rokhinson, X. Liu, J. K. Furdyna, GaMnAs-based hybrid multiferroic memory device. Appl. Phys.Lett. 92, 192501 (2008).13. C. Thiele, K. Dörr, O. Bilani, J. Rödel, L. Schultz, Influence of strain on themagnetization and magnetoelectric effect inLa0.7A0.3MnO3∕PMN-PT(001)(A=Sr,Ca). Phys.Rev.B 75, 054408 (2007).14.W. Eerenstein, M. Wiora, J. L. Prieto, J. F. Scott, N. D. Mathur, Giantsharp and persistent converse magnetoelectric effects in multiferroic epitaxial heterostructures. Nat. Mater. 6, 348 (2007).15.T. Kanki, H. Tanaka, T. Kawai, Electric control of room temperatureferromagnetism in a Pb(Zr0.2Ti0.8)O3/La0.85Ba0.15MnO3 field-effect transistor. Appl.Phys. Lett. 89, 242506 (2006).16.Y.-H. Chu et al., Electric-field control of local ferromagnetism using amagnetoelectric multiferroic. Nat. Mater. 7, 478 2008).17.S. Sahoo et al., Ferroelectric control of magnetism in BaTiO3∕Feheterostructures via interface strain coupling. Phys. Rev. B 76, 092108 (2007). 18. C.-G. Duan, S. S. Jaswal, E. Y. Tsymbal, Predicted magnetoelectric effectin Fe/BaTiO3 multilayers: Ferroelectric control of magnetism. Phys. Rev.Lett. 97, 047201 (2006).19.M. Fechner et al., Magnetic phase transition in two-phase multiferroicspredicted from first principles.Phys. Rev. B 78, 212406 (2008).20.J. Lee, N. Sai, T. Cai, Q. Niu, A. A. Demkov, preprint availableat /abs/0912.3492v1.21.K. Yamauchi, B. Sanyal, S. Picozzi, Interface effects at ahalf-metal/ferroelectric junction. Appl. Phys. Lett. 91, 062506 (2007).22.M. K. Niranjan, J. P. Velev, C.-G. Duan, S. S. Jaswal, E. Y. Tsymbal, Magnetoelectric effect at the Fe3O4/BaTiO3 (001) interface: A first-principles study. Phys. Rev. B 78, 104405 (2008).23.M. Jullière, Tunneling between ferromagnetic films. Phys. Lett.A 54, 225 (1975).24.J. F. Scott, Applications of modern ferroelectrics. Science 315, 954 (2007).25.M. Bowen et al., Nearly total spin polarization in La2/3Sr1/3MnO3 fromtunneling experiments. Appl. Phys. Lett. 82, 233 (2003).26.J. P. Velev et al., Magnetic tunnel junctions with ferroelectric barriers:Prediction of four resistance states from first principles. Nano Lett. 9, 427 (2009).27. F. Yang et al., Eight logic states of tunneling magnetoelectroresistancein multiferroic tunnel junctions.J. Appl. Phys. 102, 044504 (2007).28.Materials and methods are available as supporting materialon Science Online.29.V. Garcia et al., Giant tunnel electroresistance for non-destructivereadout of ferroelectric states. Nature460, 81 (2009).30.K. J. Choi et al., Enhancement of ferroelectricity in strained BaTiO3 thinfilms. Science 306, 1005(2004).31.K. Bouzehouane et al., Nanolithography based on real-time electricallycontrolled indentation with an atomic force microscope for nanocontactelaboration. Nano Lett. 3, 1599 (2003).32.T. J. Regan et al., Chemical effects at metal/oxide interfaces studied byx-ray-absorption spectroscopy.Phys. Rev. B 64, 214422 (2001).33.N. Hollmann et al., Electronic and magnetic properties of the kagomesystems YBaCo4O7 and YBaCo3M O7 (M=Al, Fe). Phys. Rev. B 80, 085111 (2009).34.M. Abbate et al., Soft-x-ray-absorption studies of the location of extracharges induced by substitution in controlled-valence materials. Phys. Rev.B 44, 5419 (1991).35. E. Y. Tsymbal, H. Kohlstedt, Tunneling across aferroelectric. Science 313, 181 (2006).36.M. Ye. Zhuravlev, R. F. Sabirianov, S. S. Jaswal, E. Y. Tsymbal, Giantelectroresistance in ferroelectric tunnel junctions. Phys. Rev.Lett. 94, 246802 (2005).37.M. Ye. Zhuravlev, R. F. Sabirianov, S. S. Jaswal, E. Y. Tsymbal, Erratum:Giant electroresistance in ferroelectric tunnel junctions. Phys. Rev.Lett. 102, 169901 2009).38.H. Kohlstedt, N. A. Pertsev, J. Rodriguez Contreras, R. Waser, Theoreticalcurrent-voltage characteristics of ferroelectric tunnel junctions. Phys. Rev.B 72, 125341 (2005).39.M. Gajek et al., Tunnel junctions with multiferroic barriers. Nat.Mater. 6, 296 (2007).40.M. Bowen et al., Spin-polarized tunneling spectroscopy in tunneljunctions with half-metallic electrodes.Phys. Rev. Lett. 95, 137203 (2005).41.J. D. Burton, E. Y. Tsymbal, Prediction of electrically induced magneticreconstruction at the manganite/ferroelectric interface. Phys. Rev.B 80, 174406 (2009).42.We thank R. Guillemet, C. Israel, M. E. Vickers, R. Mattana, J.-M. George,and P. Seneor for technical assistance, and C. Colliex for fruitful discussions on the microscopy measurements. This study was partially supported by theFrance-U.K. Partenariat Hubert Curien Alliance program, the French RéseauThématique de Recherche Avancée Triangle de la Physique, the European Union (EU) Specific Targeted Research Project (STRep) Manipulating the Coupling inMultiferroic Films, EU STReP Controlling Mesoscopic Phase Separation, U.K. Engineering and Physical Sciences Research Council grant EP/E026206/I, French C-Nano Île de France, French Agence Nationale de la Recherche (ANR) Oxitronics, French ANR Alicante, the European Enabling Science and Technology through European Elelctron Microscopy program, and the French Microscopie Electronique et Sonde Atomique network. X.M.acknowledges support from Comissionat per a Universitats i Recerca (Generalitat de Catalunya).。
BaTiO3 thin films obtained by sol–gel spin coating
Surface and Coatings Technology 151–152(2002)118–1210257-8972/02/$-see front matter ᮊ2002Elsevier Science B.V .All rights reserved.PII:S0257-8972Ž01.01602-4BaTiO thin films obtained by sol–gel spin coating3M.Manso-Silvan *,L.Fuentes-Cobas ,R.J.Martın-Palma ,M.Hernandez-Velez,J.M.Martınez-a ,b aa ´´´´´Duart aDepartment Fısica Aplicada,Universidad Autonoma de Madrid,28049Madrid,Spaina´´Laboratorio de Ceramicas Piezoelectricas,CIMA,S.C.,Chihuahua,C.P .31110,Mexicob´´AbstractBaTiO thin film deposition techniques are the subject of many research studies mainly due to their influence on the optical3and electrical properties of this material,which are of increasing interest for the processing of opto-and microelectronic devices.In this work,BaTiO thin films were grown onto silicon substrates by using a sol–gel spin coating procedure from Ba (OH )and 32tetraisopropyl–orthotitanate (TIPT )-based precursors.The as-grown films were annealed at different temperatures.Both the preparation of the starting solutions through a multi-step process together with an activation energy reached by sintering at approximately 8008C,were the most important parameters necessary to obtain the ferroelectric phase in these films.The evolution of the network bonds and the structural characterization of the films was studied by Fourier transform infrared (FTIR )and X-ray diffraction (XRD )techniques,respectively.In order to ascertain the suitability of BaTiO films for the processing of electronic 3devices,the electrical behaviour (I –V curves )of different thin films was also determined.ᮊ2002Elsevier Science B.V .All rights reserved.Keywords:Ferroelectrics;Sol–gel;BaTiO3;Thin films1.IntroductionThe structural properties of BaTiO (BTO )determine 3their optical and electrical properties that have induced the study of BTO materials in fields such as non-linear-optics,photo refractivity,ferroelectricity and piezoelec-tricity among others.Previous experimental results suggest that BTO thin films are suitable for the process-ing of opto-and microelectronic devices such as wave guides,high-k gates and non-volatile memories w 1–4x .Several techniques have been used to form bare or doped BTO thin films.In that sense molecular beam epitaxy w 1,2x ,plasma sputtering w 3x ,metal organic chem-ical vapour deposition w 4x and pulsed laser deposition w 5x are probably the most referenced techniques.Sol–gel BTO coatings have also been previously prepared in order to study their chemical and microstructural prop-erties w 6x .However,in this work,the perovskite phase was not reported.The sol–gel method presents several advantages for coatings processing;the control of com-position,surface morphology engineering and low tem-*Corresponding author.Tel.:q 34-913-974-919;fax:q 34-913-973-969.E-mail address:miguel.manso@uam.es (M.Manso-Silvan).´perature processing,which allows the use of thermallyfragile substrates,are its principal merits w 7x .The aim of this work is to verify that BaTiO films can be 3prepared by sol–gel spin coating and their physicochem-ical and electrical properties can be improved by sinter-ing at different temperatures.2.Experimental procedure 2.1.Sample preparationSilicon wafers (100,polished on one side )were used as substrates for the deposition of BaTiO coatings.3Before deposition,the Si wafers were treated in air at 5008C for 2h to eliminate adsorbed impurities and to form a thermally stable SiO surface film.The precursor 2solution was prepared in a two-step process from tetrai-sopropyl-orthotitanate (TIPT )and Ba (OH )(BOH ),2both from Fluka chemicals.A 0.2M TIPT solution diluted in ethanol with a TIPT y water ratio r s 0.82and w a pH s 1.27was prepared as described previously w 8x .Secondly,a 0.2M BOH solution in ethanol y acetic acid (50%)was added to an equivalent volume of the TIPT solution,thus obtaining a Ba y Ti molar ratio equal to 1.119M.Manso-Silvan et al./Surface and Coatings Technology 151–152(2002)118–121´Fig.1.XRD diffractograms of:(a )a BTO500and (b )a BTO800coating at highertemperatures.Fig.2.FTIR spectra of:(a )a BTO300,(b )a BTO500and (c )a BTO800coating.This solution was dispersed on the substrates and then spin coated at 3000rev.y min for 45s.The as-synthesi-sed xerogel films were sintered at 300,500and 8008C for 15min.In order to prepare thicker BaTiO films,a 3multilayer deposition procedure was followed,i.e.the deposition y heat-treatment cycle was repeated up to 10times.The coatings were finally sintered for 1h at the corresponding temperature.In this work,the samples were labelled BTO300,BTO500and BTO800in corre-spondence with the sintering temperature.2.2.Structural and chemical characterisationThe coatings characterisation was carried out using the following techniques.The samples thickness ranged between 0.5and 1m m as determined by profilometry.X-Ray diffraction (XRD )was performed in a Siemens diffractometer using u y 2u configuration with 0.028scan step and 6s integration time.The diffractograms were calculated by the Rietveld refinement method by using FULLPROF computing program w 9x after deletion of the Si substrate contribution.Fourier transform infrared (FTIR )spectra were recorded in a transmission scheme using a Bio-Rad FTS165spectrometer (100scans at 20Hz,4cm resolution ).y 12.3.I–V characteristicsIn order to study the electrical properties of barium titanate,Ti y BTO y Si sandwich structures were made.The metal contacts (f s 3mm )were deposited on both sides by sputtering.Electrical characterisation of these structures (I –V curves,y 1V to 1V range )was carried out in the dark,by using a Hewlett-Packard pA meter y DC voltage source,Model 4140B.3.Results3.1.Structure and compositionThe XRD diffractograms corresponding to BTO300coatings present no diffraction peaks indicating that the phase grown is amorphous.The film densification was drastically activated by sintering at 5008C XRD shoul-ders related to short range order could be observed in the diffractogram of BTO500in the 2u range of 25–308and 42–508(Fig.1a ).Sintering at 8008C promoted the crystallisation of the coatings.Multiple peaks appear-ing in the XRD diffractograms of Fig.1b confirm the polycrystalline structure of BTO800samples.The obser-vation of the (110)and (101)peaks is in agreement with the formation of a BaTiO perovskite phase 3(JCPDS 83-1880).The residual peak at 33.38,corre-sponds to the Si substrate and was omitted in order to refine the coating structure by the Rietveld method.The refinement was carried out taking account of 12reflections.The a and c cell parameters were found tobe 4.006and 4.031"0.001A,respectively.The analysis ˚of preferred orientations showed that the relative inten-sities of the diffraction peaks are in agreement with those obtained from BaTiO standards,proving that the 3substrate does not induce preferential directions in the film growth.From the FTIR spectra (Fig.2)it can be observed that in the case of BTO300coatings,two weak and wide bands are detected at 310and 460cm which y 1have been previously identified with Ti–O absorptions in a non-developed network.The weak but well-defined band at 630cm appears to be a common feature for y 1BTO300and BTO500samples.The band at 820cm y 1may arise from a carbonate trace in the BTO300coating.Several weak bands associated with titanium–alkoxi120M.Manso-Silvanet al./Surface and Coatings Technology 151–152(2002)118–121´Fig.3.Experimental I –V curves of the Ti y BTO y Si structures.Inset:general rectifying behaviour in the y 1,1V range.Fitted curves of:(a )BTO300,(b )BTO500and (c )BTO800structures.Table 1Saturation current,series resistance and ideality coefficient of the different Ti y BTO y Si structures,determined from current–voltage measurementsBTO300BTO500BTO800Saturation current,I (A )s (4.64"0.07)=10y 9(1.45"0.04)=10y 8(1.25"0.07)=10y 10Ideality coefficient,n 8.05"0.074.5"0.15.3"0.2Series resistance,R (V )s (2.81"0.05)=106(5.3"0.4)=104(8.8"0.9)=105bonds w 8x are also detected at 1040,1070and 1125cm .The intense bands at 1360and 1445cm are y 1y 1assigned to carbonate bands partially bonded to barium ions w 10x .A high content of carboxyl groups is also detected at 1560cm .The wide band observed at y 13600–3100cm is assigned to hydroxyl group asym-y 1metric vibration modes.The spectrum of the BTO500coating is very similar to the spectrum of BTO300coatings.It is only worth emphasising the decreasing absorption of some bands,mainly those assigned to carboxyl groups and hydroxyl groups,which are elimi-nated for increasing sintering temperatures.The coatings sintered at 8008C present deeper mod-ifications in their spectra.The intense band at the edge of detection (260cm )is assigned to Ti–O modes in y 1a developed network w 11x .Several components appear in the Ti–O–Ti band region.The principal ones are centred at 450cm for these coatings.A broad and y 1strong band at 500cm ,characteristic of alkaline y 1titanates w 12x ,appears with a shoulder at 560cm y 1which confirms that TiO octahedra have been formed 6w 10x .Several new features can be observed in the spectra of these coatings between 600and 1500cm .The y 1bands between 600and 900cm are sharper but weaker y 1than the corresponding bands in the BTO500film.Furthermore,the BTO500spectrum presented bands at 1405and 1450cm ,which point out the presence ofy 1a carbonated compound that was not identified in theBTO800spectrum.3.2.I–V characteristicsCurrent–voltage (I –V )measurements of the different Ti y BTO y Si structures were taken in the y 1V to 1V range.The structures analysed show a rectifying behav-iour (Fig.3,inset ).The experimental I –V curves were fitted in the 0–0.5V range to the following general expression w 13x :NŽ.I s I exp q V y IR y n kT y 1(1)w x Ž.s,i s i µ∂8i s 1where I is the saturation current,R the series resistance s,i s and n the ideality coefficient.The aim of the non-linear i analysis is to resolve the total current through the various Ti y BTO y Si structures,into several contributions due to different carrier transport mechanisms.In the three cases investigated in the present study,the experimental data could be fitted considering only one contribution (Fig.3a–c ).The non-linear least squares fitting in the 0–0.5V range was carried out by employing a Levenberg–Marquardt algorithm.Table 1shows the values of I ,n s i and R of the three structures analysed.It can be s observed that they show quite high values of the ideality coefficient,as it is usual when the density of interface states is large.This is supposed to be the case of sol–gel derived thin films.However,it is interesting to note the high values of the series resistance,which is typical of bulk barium titanate.4.Discussion and conclusionsIt has been found that the structure of BaTiO coatings 3deposited by the spin sol–gel process strongly depend on post thermal treatments of the as deposited xerogels.FTIR and XRD analysis are very useful to analyse the sequence of processes taking place during sintering.It was shown that a significant activation of the Ti–O–Ba network condensation is produced when a tempera-ture of 5008C is reached.Further sintering at tempera-tures of 8008C provided the activation of a polycrystalline phase.Moreover,a perovskite phase withlattice parameters a s 4.006and c s 4.031"0.001Awas ˚found.Although the lattice splitting does not reach thevalues obtained for bulk samples (a s 3.9945A,c s ˚4.0335A)it is shown that BTO ferroelectric coatings ˚121 M.Manso-Silvan et al./Surface and Coatings Technology151–152(2002)118–121´can be obtained by sol–gel spin coating processing from BOH based precursors.Sol–gel BTO coatings preparedfrom BOH as a starting precursor and sintered at8008C have been previously reported.They were shown to crystallise in the cubic phase so that no ferroelectricproperties could be inferred w6x.It is demonstrated inthis work that a BTO ferroelectric phase can be obtainedby using the same precursor.This result can be attributedto the solution preparation(i.e.two-step preparationinfluencing composition,viscosity),which determinethe evolution of the polycondensation process.DifferentTi y BTO y Si structures were elaborated in order to deter-mine the electrical parameters(saturation current,seriesresistance and ideality coefficient)from the experimentalI–V curves.It should be noted that all the structures studied show quite high values of the ideality coefficient(between4.5and8.05),which might be explained by the presence of a large concentration of interface states,specially in the BTO300structures,which contain ahigher OH group density.These structures also showhigh values of the series resistance,which vary accord-ingly with the evolution of the sol–gel condensationand densification mechanisms.The BTO300structurepresents the highest series resistance due to a lack ofdensification and the presence of organic species.DenserBTO500structures undergo a reduction in the seriesresistance due to Ba mobility in a non-developed2qnetwork.The crystallisation of the BTO800film inducesBa fixation in lattice sites and lead to a stable coatingwith high resistivity values.The spin coating process combined with an adequate sol preparation has demonstrated the feasibility of depos-iting BaTiO coatings with ferroelectric structure.Fur-3ther experiments must determine their applicability for the development of opto-electronic devices. AcknowledgementsWe acknowledge Direccion General de Investigacion´´of CAM for the financial support of a research grant to M.Manso.Referencesw1x R.A.McKee,F.J.Walker,J.R.Conner, E.D.Specht,Appl.Phys.Lett.59(1991)782.w2x T.Zhao,H.Lu,F.Chen,G.Yang,Z.Chen,J.Appl.Phys.87 (2000)7448.w3x N.M.Abuhabda,C.R.Aita,J.Appl.Phys.71(1992)3045.w4x B.Bihari,J.Kumar,G.T.Stauf,P.C.Van Buskirk,C.S.Hwang, J.Appl.Phys.76(1994)1169.w5x Y.Liu,Z.Chen,C.Li,D.Cui,Y.Zhou,G.Yang,J.Appl.Phys.81(1997)6328.w6x N.M.Kamalasanan,N.D.Kumar,S.Chandra,J.Appl.Phys.76(1994)4603.w7x G.W.Scherrer,C.J.Brinker,Sol–Gel Science:The Physics and Chemistry of Sol–Gel,Academic Press,Boston,MA, 1990.w8x M.Burgos,nglet,J.Sol–Gel Sci.Technol.17(1999) 267.w9x J.Rodrıgez-Carvajal,Physica B192(1993)55.´w10x M.C.Blanco-Lopez,G.Fourlaris,B.Rand,F.L.Riley,J.Am.´Ceram.Soc.82(1999)1777.w11x M.Burgos,nglet,Thin Solid Films349(1999)19.w12x R.A.Nyquist,R.O.Kagel,Infrared Spectra of Inorganic Com-pounds,Academic Press,San Diego,CA,1997,pp.98–99.w13x J.Perez-Rigueiro,C.Jimenez,R.Perez-Casero,J.M.Martınez-´´´´Duart,J.Vac.Sci.Technol.B14(1996)2623.。
BaTiO
Bi蒸汽提升PTCR同时，抑制晶粒 长大
常规掺杂
Bi2O3蒸汽掺杂
B 2 O i3 2 B B • iaV B '' a3 O O
Bi蒸汽掺杂抑制晶粒长大
Mn和Bi蒸汽的协同作用
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Resistance /
107 106 105
104
103
102
101
100 0
Y-BST Y-BST+Bi O vapor
利用移峰效应可将铁电陶瓷在居里温度处出现的介电常数的峰值移到室温附近， 这有利于制造大容量、小体积的陶瓷电容器。
也可利用移峰效应改善陶瓷材料的电容温度系数：为了在工作情况下（室温附近 ）材料的介电常数与温度关系曲线尽可能平缓，即要求居里点远离室温。如掺入 PbTiO3可使BaTiO3居里点升高。
压峰效应
其他蒸汽掺杂
108
J. Q. Qi, Z. L. Gui, L. T. Li and Y. J. Wu, “Positive
107
temperature coefficient resuistance effect in Ba1-
xSrxTiO3 ceramics modified with Bi2O3 and PbO by
0.10
Temperature / 0C
104 0.12
Qi Jianquan, Chen Wanping, Zhang Zhongtai and Tang Zilong, Acceptor compensation in (Sb,Y)-doped semiconducting Ba 1-x SrxTiO3, J. Mater. Sci., Vol.32, p713-717, 1997
BaTiO3PPT精选文档
32
33
2、除铁
制备粉料时往往混入铁质，常采用电 磁除铁器除去其中的铁质。电磁除铁 器有干式和湿式。
3、过滤
一般采用压滤的方法除去湿法细磨 浆料中的水。
困料
困料是为使压滤出来的泥饼水分分 布均匀，常把泥饼放在避光、空气 不流通的室内或密闭容器内，保持 10一20天。
练泥
经过压滤、困料的坯料组织疏松、不均匀，含有大量的气 泡，可塑性降低，挤压成型性差。采用练泥的方法可除去 坯料中的气泡，提高其均匀性和致密度。
5
钛酸钡材料的特点
BaTiO3的很高。a轴方向与c轴方向的巨大差异表明：在 电场作用下， BaTiO3中的离子沿a轴方向具有更大的可动 性。
相变温度附近， 均具有峰值，在TC下的最高。说明相 变温度附近，离子具有较大可动性，在电场作用下易于 使晶体中的电畴沿电场方向取向。
与相变的热滞现象相对应， 随温度变化时也存在热滞现 象，在四方 斜方相变温度及斜方 菱形相变温度附 近表现得很明显。
0.10
Temperature / 0C
104 0.12
Qi Jianquan, Chen Wanping, Zhang Zhongtai and Tang Zilong, Acceptor compensation in (Sb,Y)-doped semiconducting Ba 1-x SrxTiO3, 2J0. Mater. Sci., Vol.32, p713-717, 1997
压峰效应
压峰效应是为了降低居里点处的介电常数的峰值，即 降低ε-T非线性，也使工作状态相应于ε-T平缓区。 例如在BaTiO3中掺入CaTiO3可使居里峰值下降。 常用的压峰剂（或称展宽剂）为非铁电体。如在 BaTiO3掺入Bi2/3SnO3，其居里点几乎完全消失，显示 出直线性的温度特性，可认为是加入非铁电体后，破 坏了原来的内电场，使自发极化减弱，即铁电性减小。
以偏钛酸为钛源的化学共沉淀法制备钛酸钡粉体
Ａｂｓｔｒａｃｔ
Ｂａｒｉｕｍ ｔｉｔａｎａｔｅ ｐｏｗｄｅｒ ｗａｓ ｆａｂｒｉｃａｔｅｄ ｂｙ ｔｈｅ ｗａｙ ｏｆ ｃｈｅｍｉｃａｌ ｃｏｐｒｅｃｉｐｉｔａｔｉｏｎ，ｉｎ ｗｈｉｃｈ Ｂ＿ＴｉＤ
（ＯＨ）２ ｗａｓ ｕｓｅｄ ａｓ ｔｈｅ ｒｅｓｏｕｒｃｅ ｏｆ ｅｌｅｍｅｎｔ ｔｉｔａｎｉｕｍ ａｎｄ，ｆｉｒｓｔ ｏｆ ａｌｌ，Ｔｉ４＋ｗａｓ ｍａｄｅ ｉｎ［ＴｉＯ（Ｃ２０４）２］卜ａｎｄ ｔｈｅｎ ＣＯ—
（２）实验表明，沉淀反应过程的ｐＨ值对钡钛摩尔比起着关 键的作用，ｐＨ值控制在２．５左右比较理想。
（３）在ｐＨ值为２．５时进行共沉淀反应制备草酸氧钛钡，并 在９５０。Ｃ恒温３．５ｈ进行热分解，用化学共沉淀的方法制备出平 均粒度为０．３７”ｍ、粒度分布均匀、形貌为球形、钡钛比达到 ０．９９９、具有四方相的晶体结构、主含量为９９．８１％、杂质含量少 的ＢａＴｉ０３粉体。
品的粒度和晶型有利。在该工艺条件下制备的样品平均粒度为０．３７“ｍ、粒度分布均匀，形貌为球形，钡钛比达到
０．９９９，具有四方相的晶体结构、主含量为９９．８１％，杂质含量符合行业标准要求。
关键词 钛酸钡化学共沉淀法 电子陶瓷正温度系数偏钛酸
Ｐｒｅｐａｒａｔｉｏｎ ｏｆ Ｂａｒｉｕｍ Ｔｉｔａｎａｔｅ Ｐｏｗｄｅｒｓ ｂｙ Ｃｈｅｍｉｃａｌ Ｃｏｐｒｅｃｉｐｉｔａｔｉｏｎ ｗｉｔｈ
钛酸钡粉体的制备方法主要有固相法、水热合成法、化学共 沉淀法、溶胶一凝胶法等［２’３］。从经济成本、生产工艺、粉体性能 等方面综合考虑，化学共沉淀法是斛备钛酸钡粉体的较好方 法［４］。我国虽然已有一些厂家采用草酸盐共沉淀法生产 ＢａＴｉ０３粉体，但钛元素来源均是价格较高、有效钛含量较低、有 大量液体和气体污染物排放的四氯化钛，并存在着微观均匀性 不好、产品性能不稳定、粒度较大及粒度分布不均匀等问题，生 产的ＢａＴｉＯｓ粉体只能用于性能要求不高的元件，致使我国电 子元件制造企业必须大量进口高纯超细钛酸钡来满足高性能元 件制造。
BaTiO3织构PMN-32PT单晶电畴偏光显微镜论文
外场对取向性铁电材料电畴和性能的影响【摘要】铁电材料作为一类重要的功能材料被广泛地应用于微电子学和微机械系统等领域,成为当前国际上新型功能材料研究领域中的一个热点。
铁电材料的铁电性受到晶化、取向等因素的影响,其中,取向对改善电学性能起着非常重要的作用,因而为了更优化地裁剪铁电材料的性能,取向性铁电材料的组织与性能研究越来越引人注目。
一直以来,弛豫铁电单晶和织构陶瓷是人们研究铁电材料的两大结构体系。
本文首先采用模板晶粒生长法(TGG)制备了钛酸钡织构陶瓷,并研究了外场对其性能的影响规律；其次利用偏光显微镜研究了不同切向的PMN-32PT单晶电畴组态在外场作用下的变化规律,以及外场对PMN-32PT单晶性能的影响规律。
研究结果如下：1)采用模板晶粒生长(TGG)技术制备了钛酸钡织构陶瓷；发现随烧结温度的升高(1180℃-1250℃),织构度也随之升高(30.60%～34.62%)；高织构钛酸钡陶瓷获得高的压电常数129 pC/N。
2)不同频率下的介电温谱表明,钛酸钡陶瓷表现出明显的弥散相变特性；相对于商用原料钛酸钡陶瓷,等轴钛酸钡陶瓷和织构钛酸钡陶瓷表现出更小的弥散相变；与等轴钛酸钡陶瓷相比,织构钛酸钡陶瓷的居里温度稍有降低,然而介... 更多还原【Abstract】 As an important functional material,ferroelectric materials have been widely used in microelectronics and micro-mechanical systems. Recently, theresearch on ferroelectric materials has become a hot. The ferroelectric properties are affected by crystallization, orientation etc. The orientation plays an important role in improving electrical properties. To further improving ferroelectric properties, the oriented ferroelectric materials are drawing more and more attention. Relaxation single cryst... 更多还原【关键词】BaTiO3；织构；PMN-32PT单晶；电畴；偏光显微镜；【Key words】BaTiO3；Texture；PMN-32PT single crystal；domain；polarizing microscope；摘要3-5Abstract 5-6主要符号表9-101 绪论10-221.1 引言10-111.2 无铅压电陶瓷概述11-171.2.1 无铅压电陶瓷的研究进展11-141.2.2 织构化陶瓷14-171.3 弛豫铁电单晶概述17-201.3.1 弛豫铁电单晶的研究进展17-181.3.2 弛豫铁电单晶的制备方法18-201.3.3 弛豫铁电单晶的电畴研究201.4 本文的研究目的和意义20-221.4.1 选题的背景和意义20-211.4.2 研究内容21-222 实验方法和过程22-312.1 钛酸钡织构陶瓷的实验研究过程22-272.1.1 前驱体制备22-232.1.2 流延成型23-242.1.3 素坯排塑242.1.4 烧结24-252.1.5 组织、结构分析和性能测试25-272.2 PMN-PT弛豫铁电单晶的实验研究过程27-312.2.1 试样的制备272.2.2 晶体的定向和切割27-282.2.3 PMN-32PT样品的研磨、抛光282.2.4 样品的清洗282.2.5 样品的退火28-292.2.6 试样的极化工艺292.2.7 压电性能的测试292.2.8 介电性能的测试292.2.9 电畴组态观察29-313 模板晶粒生长法制备的钛酸钡织构化陶瓷的组织与性能31-363.1 钛酸钡模板313.2 预烧粉体XRD分析31-323.3 流延膜XRD分析32-333.4 陶瓷样品的体积密度33-343.5 钛酸钡织构陶瓷样品的组织34-353.6 本章小结35-364 外场对钛酸钡陶瓷性能的影响36-424.1 样品制备364.2 实验结果及分析36-404.2.1 无偏置电压下钛酸钡陶瓷的介电性能36-394.2.2 偏置电场下BT陶瓷的介电性能39-404.3 钛酸钡陶瓷的介电非线性机理40-414.4 本章小结41-425 外场对PMN-32PT单晶的电畴形貌的影响42-545.1 PMN-32PT单晶样品的结构425.2 [100]_(cub)切型PMN-32PT单晶电畴在温度场下的行为42-495.2.1 [100]_(cub)切型PMN-32PT单晶电畴在升温过程中的行为42-445.2.2 [100]_(cub)切型PMN-32PT单晶电畴在降温过程中的行为44-455.2.3 [100]_(cub)切型PMN-32PT单晶电畴在不同温度区段的消光行为45-495.3 [100]_(cub)切型PMN-32PT单晶电畴在直流电场下的行为49-515.3.1 [100]_(cub)切型PMN-32PT单晶电畴在正向电场下的行为49-505.3.2 [100]_(cub)切型PMN-32PT单晶电畴在反向电场下的行为50-515.4 [100]_(cub)切型PMN-32PT单晶电畴在交流电场下的行为51-535.5 本章小结53-546 外场对PMN-32PT单晶性能的影响54-586.1 电场对[100]_(cub)切型PMN-32PT单晶介电性能的影响54-566.2 力场对[100]_(cub)切型PMN-32PT单晶的性能影响56-576.3 本章小结57-587 结论58-60参考文献。
水热法合成钛酸钡
Hydrothermal Synthesis of Barium Titanate:Effect of Titania Precursor and Calcination Temperature on Phase TransitionNatarajan Sasirekha,Baskaran Rajesh,and Yu-Wen Chen*Department of Chemical Engineering,Nanocatalysis Research Center,National Central Uni V ersity,Chung-Li320,Taiwan,Republic of ChinaNanosized barium titanate powders were synthesized by a hydrothermal method.The effect of titania precursorson the phase transition of BaTiO3with respect to Ba/Ti ratio,reaction temperature,reaction time,and calcinationtemperature was investigated.The synthesized materials were characterized by X-ray diffraction,scanningelectron microscopy,and transmission electron microscopy.BaTiO3in pure cubic phase with sphericalmorphology was observed with a lower calcination temperature,Ba/Ti ratio,reaction temperature,and time.Increase in the tetragonal phase was ascertained in treatments at higher reaction temperature with a longerreaction time.The lattice hydroxyl release is believed to be the reason for tetragonality at high reaction andcalcination temperatures.To prepare tetragonal BaTiO3using HClO4-TiO2,the optimum synthesis conditionsviz.,Ba/Ti ratio,reaction temperature,and reaction time,are1.2,160°C,and3h,respectively,at a calcinationtemperature of1150°C.The reaction time and reaction temperature for the cubic-tetragonal phasetransformation of BaTiO3shifted toward shorter reaction time and lower reaction temperature when TiO2was synthesized by hydrolysis using HClO4as the acid catalyst.1.IntroductionBarium titanate(BaTiO3),one of the most well-known ferroelectrics,has played an important part in the modern ceramic industry since the discovery of ferroelectric properties in the tetragonal phase of BaTiO3during the1940s.1It has been broadly used as a dielectric material in multilayer ceramic capacitors(MLCCs),2-4printed circuit boards(PCBs),5,6dy-namic random access memory(DRAM),positive temperature coefficient of resistance thermistors(PTCRs),piezoelectric sensors for ultrasonic and measuring devices,pressure transduc-ers,infrared detectors,and electrooptic devices7-9due to its unique perovskite structure(ABO3)and exceptionally high dielectric[(2-5)×103]10and piezoelectric properties at room temperature.However,these outstanding behaviors mainly depend on the crystal structure,shape,size,stoichiometry, homogeneity,and surface properties of BaTiO3,which in turn depends on the synthesis method.Among the crystal structures of BaTiO3,the cubic phase exhibits paraelectric properties,while the tetragonal phase shows ferroelectric properties.The direct generation of tetragonal BaTiO3is of considerable interest. The conventional synthesis of barium titanate compounds typically involves high-temperature(∼1200°C)calcinations of a BaCO3and TiO2powder mixture,which often results in low purity and polydispersity due to high reaction temperature and heterogeneous solid-phase reaction.11,12Nevertheless,fine Ba-TiO3ceramics can be prepared by wet-chemistry synthesis techniques,including the coprecipitation method,13coprecipi-tation in combination with the inverse microemulsion method,14,15 sol-gel processing,16,17the hydrothermal method,18-21spray pyrolysis,22the oxalate route,23the high-temperature ceramic route,the microwave hydrothermal method,24-27and the polymeric precursor method.28Hydrothermal synthesis for the preparation of crystalline BaTiO3has gained popularity recently.29-33It involves the chemical reactions of Ba(OH)2,TiO2,or gels of Ba-Ti acetate mixtures at a high temperature.It has the advantage of producing finer particles with more uniform size.In addition,the interac-tion between the solid and fluid phases determines the physical characteristics of BaTiO3,and hence,the synthesis allows one to control the particle size by adjusting the synthesis parameters, such as the reaction temperature,time,and pH values.The synergistic effect of solvent,temperature,and pressure on the ionic reaction equilibrium in the hydrothermal reaction medium can stabilize the formation of BaTiO3and retards the formation of impurities.Also,the precursors for the preparation of BaTiO3 by hydrothermal synthesis are readily available,inexpensive, and easy to handle,which makes the hydrothermal synthesis an easy and effective method to adopt for the synthesis of BaTiO3.Although the hydrothermal synthesis has the above-mentioned advantages,the formation and growth mechanisms in BaTiO3synthesis have not been well understood.There are a few articles about the study of phase transformation of BaTiO3 with preparation conditions.However,so far there is no report about the effect of titania precursor,which plays a vital role in the growth mechanism of BaTiO3synthesis,on the phase transformation of BaTiO3.In this paper,a systematic investiga-tion on the effect of titania precursor on phase transformation of BaTiO3with respect to Ba/Ti ratio,reaction temperature, reaction time,and calcination temperature is presented.2.Experimental Section2.1.Hydrothermal Synthesis of BaTiO3.Barium titanate powders were synthesized by hydrothermal method.All the reagents used were of analytical grade.Hydrothermally produced TiO2particles34prepared using hydrochloric acid and perchloric acid as acid precursors were used as the titania source.The detailed characteristics of these materials have been presented in our previous paper.34Ba(OH)2‚8H2O(Showa Chemical Co., Ltd.)and TiO2powder(rutile)34were mixed with a Ba/Ti ratio (1.2,1.4,1.6,1.8,2.0)in a50mL autoclave with45mL of deionized water.The autoclave was sealed,shaken,and placed in an oven at160°C for a variable reaction period ranging from*To whom correspondence should be addressed.Tel.:(886)3422751,ext34203.Fax:(886)34252296.E-mail:ywchen@cc.ncu.edu.tw.1868Ind.Eng.Chem.Res.2008,47,1868-187510.1021/ie070986m CCC:$40.75©2008American Chemical SocietyPublished on Web02/16/20083to 24h.After cooling naturally to room temperature,the contents of the autoclave were diluted in 90mL of 0.1M formic acid with stirring for 5min in an attempt to remove any BaCO 3formed and the addition of excess Ba 2+to the starting solution.The mixture was filtered and washed with distilled water (500mL)three times,and the residue was dried in oven at 100°C for 24h.2.2.Characterization.The crystalline phase of BaTiO 3was analyzed by powder X-ray diffraction (XRD)using a Siemens D500automatic powder diffractometer.Nickel-filtered Cu K R radiation (λ)0.15418nm)was used with a generator voltage of 40kV and a current of 29mA.Bragg -Brentano focusing geometry was employed with an automatic divergence slit (irradiated sample length was 12.5nm),a receiving slit of 0.1nm,a fixed slit of 4°,and a proportional counter as a detector.It was operated in the step scan mode,at scanning speeds of 0.1°2θ/s and 1s step time in the range 20-80°for barium titanate.Scherrer’s equation 35was used to calculate the crys-tallite size of barium titanate crystals from the full width at half-maximum of the XRD peak.The morphology of the particles was analyzed by scanning electron microscopy (SEM)and transmission electron micros-copy (TEM).SEM images were acquired with a Hitachi S-800field emission microscope using an acceleration voltage of 20kV.The samples were coated with Au prior to analysis and imaged directly.TEM images were obtained on a JEOL JEM-2000FX Πmicroscope using an electron beam generated by a tungsten filament and an accelerating voltage of 160kV,a beam diameter of approximately 1-2µm,and an objective lens aperture of 20µm.The sample grids were prepared via sonication of powdered sample in ethanol for 10min and evaporating 1drop of the suspension onto a carbon-coated,porous film supported on a 3mm,200-300mesh copper grid.TEM images were recorded at a magnification of 100000-400000×.The magnification was calibrated in pixels per nanometer on the camera.3.Results and Discussion3.1.Effect of Ba/Ti Ratio.BaTiO 3powder,prepared with Ba/Ti ratio in the range of 1.2-2.0,was characterized by XRD,SEM,and TEM techniques to study the cubic -tetragonal phase transformations and surface modifications of BaTiO 3.The ratio of Ba/Ti was chosen to be greater than 1to avoid contamination of BaTiO 3(s)with excess TiO 2(s)under the conditions of hydrothermal synthesis.24Furthermore,Ba/Ti >1increases the pH of the solution,which is an important thermodynamicvariable for the synthesis of perovskite materials,and helps to avoid the addition of an alkaline mineralizer to facilitate the formation of BaTiO 3.According to the thermodynamic calcula-tions of stability diagrams for the hydrothermal Ba -Ti system,high pH and Ba/Ti >1are necessary for the synthesis of high-purity BaTiO 3crystals.28XRD patterns of the as-synthesized BaTiO 3,with Ba/Ti ratios of 1.2,1.4,1.6,1.8,and 2.0,showed the characteristic peaks of both cubic BaTiO 3(JCPDS File No.79-2263)and those of BaCO 3for both TiO 2precursors.Modest BaCO 3contamination was noted in almost all the samples due to the introduction of airborne CO 2,which would dissolve as CO 32-and reacts with Ba 2+to form BaCO 3during the posttreatment.26,36The formation of BaCO 3,observed in this work,is quite common in the hydrothermal processing as BaCO 3can precipitate at lower pH values than those needed to precipitate BaTiO 3.24AccordingtoFigure 1.XRD patterns of BaTiO 3synthesized at various Ba/Ti ratios using (a)HCl-TiO 2and (b)HClO 4-TiO 2precursors and washed with formic acid followed by calcination at 900°C (2h).Table 1.Effect of Ba/Ti Ratio and Calcination Temperature on the Crystalline Phase of BaTiO 3calcination temp (°C)Ti precursor Ba/Ti ratio synthesis temp (°C)synthesis time (h)9001150crystalline phase HCl-TiO 21.21606--cubic 1.41606--cubic 1.61606--cubic 1.81606--cubic2.01606--cubic HClO 4-TiO 21.21606--cubic 1.41606--cubic 1.61606--cubic 1.81606--cubic2.01606--cubic HCl-TiO 21.21606yes -cubic 1.41606yes -cubic 1.61606yes -cubic 1.81606yes -cubic2.01606yes -cubic HClO 4-TiO 21.21606yes -cubic 1.41606yes -cubic 1.61606yes -cubic 1.81606yes -cubic2.01606yes -cubic HCl-TiO 21.21606-yes cubic 1.41606-yes cubic 1.61606-yes cubic 1.81606-yes cubic2.01606-yes cubic HClO 4-TiO 21.21606-yes tetragonal 1.41606-yes tetragonal 1.61606-yes tetragonal 1.81606-yes tetragonal2.01606-yestetragonalInd.Eng.Chem.Res.,Vol.47,No.6,20081869the thermodynamic stability diagram of Ba -Ti systems,BaCO 3precipitates at lower pH values than those needed to precipitate BaTiO 3.28Moreover,the formation of BaCO 3is more predomi-nant in BaTiO 3prepared using HClO 4-TiO 2,which indicates its lower pH compared with those prepared using HCl-TiO 2.The relative pH of HClO 4is lower than that of HCl as the p K a values of HClO 4and HCl are -10and -7,respectively.37The XRD results of the as-synthesized BaTiO 3illustrate the absence of apparent peak splitting at 2θ)45°,which corre-sponds to the tetragonal phase of BaTiO 3(JCPDS Card No.5-0626)and hence confirmed the cubic structure with symmetry Pm 3m .In order to remove BaCO 3,the powders were washed with formic acid and calcined at 900°C for 2h.Figure 1shows XRD patterns of BaTiO 3chemically treated with formic acid and calcined at 900°C,which confirms the removal of BaCO 3from BaTiO 3prepared using HCl-TiO 2and HClO 4-TiO 2.Moreover,it indicates that there is no change in the cubic phase of BaTiO 3upon treatment with formic acid.The effect of Ba/Ti ratio and calcination temperature on the crystal structure of BaTiO 3is summarized in Table 1.It can be observed that,even at a Ba/Ti molar ratio of 2without heat treatment,the tetragonal splitting of the diffraction peaks corresponding to the (200)and (002)planes of the perovskite BaTiO 3could not be distinguished,indicating the presence of pure cubic crystalline phase.Shi et al.36observed tetragonal BaTiO 3crystallites when the precursor with high Ba/Ti molar ratio of 3was used,which reduces the probability of forming barium vacancies and stabilized tetragonal phase.The influence of calcination temperature at various Ba/Ti ratios on the cubic -tetragonal phase transition of BaTiO 3can be observed by comparing Figures 1and 2.Figure 2shows the XRD patterns of BaTiO 3prepared using HCl-TiO 2and HClO 4-TiO 2,subjected to calcination at 1150°C.For BaTiO 3samples prepared using HCl-TiO 2,the XRD results show that the crystalline phase is metastable cubic phase for all the samples at 900and 1150°C.The peaks are very sharp,indicating that the crystalline structure is well developed.Generally,the tetragonality of BaTiO 3is deduced from the plane spacing of (002)over that of (200);the corresponding peak appears near 45°in the XRD patterns.Because the peak splitting at 45°is a predominant one to confirm the formation of tetragonal phase,tetragonal splitting of the peaks corresponding to (200)and (002)planes have been chosen to verify the formation of tetragonal phase.Normally,BaTiO 3cubic -tetragonal phase changes begin at 900and 1150°C,but for BaTiO 3prepared using HCl-TiO 2the phase transformation does not proceed to completion even at 1150°C.In the case of BaTiO 3prepared using HClO 4-TiO 2and calcined at 900°C,the XRD patterns are almost identical irrespective of Ba/Ti ratio and cubic phase is observed.However,tetragonal phase was observed at 2θ)45°for the powders synthesized at Ba/Ti ratios of 1.2,1.4,1.6,1.8,and 2.0and calcined at 1150°C for 2h,as shown in Figure 2b.A slight increase in the intensity of tetragonal phase with Ba/Ti ratio may be due to the removal of barium vacancies (charge compensator of OH -defect)by excess barium content,which stabilizes tetragonality.Figures 3and 4show the SEM pictures of BaTiO 3samples prepared using HCl-TiO 2and HClO 4-TiO 2,respectively,at various Ba/Ti ratios of 1.2,1.4,1.6,and 1.8.The particles agglomerated in a spherical shape with ca.0.05-0.15µm (ca.50-150nm)diameters.A possible mechanism for the formation of BaTiO 3by hydrothermal synthesis is the dissolution-Figure 2.XRD patterns of BaTiO 3synthesized at various Ba/Ti ratios using (a)HCl-TiO 2and (b)HClO 4-TiO 2precursors and calcined at 1150°C (2h).Figure 3.SEM micrographs of as-synthesized BaTiO 3prepared using HCl-TiO 2at various Ba/Ti ratios:(a)1.2,(b)1.4,(c)1.6,and (d)1.8.1870Ind.Eng.Chem.Res.,Vol.47,No.6,2008precipitation method,38in which there is a chemical equilibrium between TiO 2and [Ti(OH)x ]4-x .[Ti(OH)x ]4-x ,which is a highly active species,can combine with Ba 2+to form a new nucleus,and hence,with an increase in Ba/Ti ratio the chance for the formation of a new nucleus by [Ti(OH)x ]4-x increases and leads to a decrease in the particle size of BaTiO 3.Figure 3shows the decrease in the particle size of BaTiO 3with an increase in Ba/Ti ratio,which is in accord with the dissolution -precipitation mechanism.TEM results show that the primary particles of the sample prepared at Ba/Ti ratios of 1.2and 1.4using HCl-TiO 2as precursor are spherical in shape with 50-60nm diameters,as shown in Figure 5a,b,whereas the particle sizes of BaTiO 3prepared using HClO 4-TiO 2are in the range of 40-50nm diameters (Figure 5c,d).Moreover,it can be ascertained from TEM images that the Ba/Ti ratio increases the cluster size of primary particles.Besides,it should be noted that the stability of cubic and tetragonal phases depends on the critical size of BaTiO 3particles and the critical size was reported to be ∼50nm.39The crystallite size of BaTiO 3is a principal factor controlling tetragonality because the surface defects of nanoc-rystallites are predominant over the bulk ones below a certain critical size of BaTiO 3.The surface defects can prevent the completion of phase transformation,leading to high strains within the crystal.Increase in the cluster size of primary particles reduces the strain within the cubic structure for distortion.It can be concluded that the phase transition of cubic BaTiO 3occurs at 1150°C irrespective of Ba/Ti ratio when HClO 4-TiO 2was used as the TiO 2precursor.Moreover,the primaryparticleFigure 4.SEM micrographs of as-synthesized BaTiO 3prepared using HClO 4-TiO 2at various Ba/Ti ratios:(a)1.2,(b)1.4,(c)1.6,and (d)1.8.Figure 5.TEM micrographs of BaTiO 3prepared using (a)HCl-TiO 2,Ba/Ti )1.2;(b)HCl-TiO 2,Ba/Ti )1.4;(c)HClO 4-TiO 2,Ba/Ti )1.2;and (d)HClO 4-TiO 2,Ba/Ti )1.4.Figure 6.XRD patterns of BaTiO 3synthesized at various temperatures using (a)HCl-TiO 2and (b)HClO 4-TiO 2precursors calcined at 900°C (2h).Ind.Eng.Chem.Res.,Vol.47,No.6,20081871size of BaTiO 3prepared using HClO 4-TiO 2is smaller than that prepared using HCl-TiO 2,which can be ascribed to the smaller particle size of TiO 2prepared by HClO 4.34The agglomeration of BaTiO 3nanoparticles at higher calcination temperature promotes the stability of the tetragonal phase.3.2.Effect of Synthesis Temperature.To study the influence of synthesis temperature on the phase of BaTiO 3and particle morphology,BaTiO 3was prepared at 80,120,160,180,and 200°C while keeping the rest of the process parameters as Ba/Ti )1.2and the synthesis time as 6h.The XRD patterns of BaTiO 3obtained at different reaction temperatures and calcined at 900°C are given in Figure 6.The XRD results illustrated well-developed cubic crystalline phase and the intensity of the peaks increased with reaction temperature.There is a possibility of decrease in the unit-cell volume and decrease in density with an increase in the reaction temperature due to the release of lattice hydroxyls.20The influence of calcination temperature withsynthesis temperature is compiled in Table 2.Figure 7depicts the XRD patterns of BaTiO 3calcined at 1150°C.With increased synthesis temperature,the diffraction peaks related to the (200)and (002)planes of the tetragonal BaTiO 3separated and the c /a ratio of the lattice increased,confirming the cubic -tetragonal phase transition with synthesis temperature.The splitting of the (200)reflection is apparent for the samples synthesized at 180and 200°C,suggesting the tetragonal phase.The intensity ratio of 45°peaks significantly increased as the synthesis temperature increased.The presence of shoulders at 80and 120°C represents the formation of partially tetragonal phase in cubic BaTiO 3.The stabilization of tetragonal phase with increase in temperature may be due to the removal of hydroxyl groups in the BaTiO 3lattice.Figure 8depicts the SEM micrographs of BaTiO 3prepared using HCl-TiO 2at 80,120,180,and 200°C for 24h.The particles agglomerated into a spherical shape,and the particle sizes estimated from the SEM micrographs are within 0.05-0.15µm in diameter.When the synthesis temperature increased from 80to 200°C and with reacting for 24h,the particle size of BaTiO 3increased to 0.09-0.15µm.The shape of the particles was observed to be spherical independent of treated tempera-tures.The increase in the synthesis temperature leads to an increase in the particle size,which may explain the stronger agglomeration at higher temperature.However,for BaTiO 3prepared using HClO 4-TiO 2(Figure 9),the secondary particle size increased from 0.05to 0.10µm with an increase in the reaction temperature from 80to 200°C for 24h,confirming that the particle size of BaTiO 3was dependent on synthesis temperature.The overall shape of the agglomerated secondary particle size was estimated to be 0.05-0.10µm in diameter.As shown in Figure 10,the particles consist of near-monodis-perse spherical nanoparticles of BaTiO 3.The agglomeration of primary particles with an increase in the reaction temperature can be identified from TEM images,as shown in Figure 10.The clusters of primary particles observed for HClO 4-TiO 2precursor are more than those of HCl-TiO 2.This result suggested that the crystallite size is one of the vital factors that control tetragonality.At a higher reaction temperature,the phase transition has occurred from cubic to tetragonal.3.3.Effect of Calcination Temperature.The stability of the cubic phase in BaTiO 3prepared by hydrothermal synthesis at room temperature may be accounted for by the presence of weakly bound water molecules absorbed onto the surface of the particles and the more strongly bonded structural water as lattice OH -ions.The content of barium vacancy as well as OH -defects in the cubic crystallites is higher than that inTable 2.Effect of Synthesis Temperature and Calcination Temperature on the Crystalline Phase of BaTiO 3calcination temp (°C)Ti precursor Ba/Ti ratio synthesis temp (°C)synthesis time (h)9001150crystalline phase HCl-TiO 21.2806--cubic 1.21206--cubic 1.21606--cubic 1.21806--cubic 1.22006--cubic HClO 4-TiO 21.2806--cubic 1.21206--cubic 1.21606--cubic 1.21806--cubic 1.22006--cubic HCl-TiO 21.2806yes -cubic 1.21206yes -cubic 1.21606yes -cubic 1.21806yes -cubic 1.22006yes -cubic HClO 4-TiO 21.2806yes -cubic 1.21206yes -cubic 1.21606yes -cubic 1.21806yes -cubic 1.22006yes -cubic HCl-TiO 21.2806-yes cubic 1.21206-yes cubic 1.21606-yes cubic 1.21806-yes tetragonal 1.22006-yes tetragonal HClO 4-TiO 21.2806-yes cubic 1.21206-yes cubic 1.21606-yes tetragonal 1.21806-yes tetragonal 1.22006-yestetragonalFigure 7.XRD patterns of BaTiO 3synthesized (Ba/Ti )1.2)at various temperatures using (a)HCl-TiO 2and (b)HClO 4-TiO 2precursors and calcined at 1150°C (2h).1872Ind.Eng.Chem.Res.,Vol.47,No.6,2008tetragonal phase.The cubic -tetragonal phase transformation at higher reaction temperature and calcination temperature is due to the elimination of OH -vacancies from the lattice with heat treatment,which leads to the tetragonal stability.At the Curie point,where BaTiO 3undergoes a phase transition,the relative displacement of cation sublattice to O 2-sublattice causes the phase transition of BaTiO 3from cubic to tetragonal.The oxygen vacancies have significant mobility above 650°C,whereas the cation vacancies acquire measurable mobility only above 1050°C.24Moreover,the decrease in the lattice parameter of the crystallites with temperature led to the conclusion that the removal of OH -defects caused the enlargement of the unit cell and released the lattice strain to form the tetragonal phase.3.4.Effect of Synthesis Time.The effect of synthesis time on the formation of crystalline BaTiO 3was also studied by performing the experiments at different reaction times ranging from 3to 24h at 160°C with Ba/Ti )1.2.The crystalline form at shorter periods of time,viz.,3and 6h,is primarily the metastable cubic form.SEM micrographs indicated no signifi-cant difference in the morphology.The cluster size was larger by extending the processing time,but the particle size has no significant difference.BaTiO 3powders prepared by using HClO 4-TiO 2as the precursor resulted similarly to those prepared by using HCl-TiO 2as the precursor.The XRD patterns as shown in Figure 11confirm the cubic phase of BaTiO 3prepared using HCl-TiO 2and HClO 4-TiO 2calcined at 900°C.Figure 12a shows the typical phase transformation of BaTiO 3prepared using HCl-TiO 2.When BaTiO 3prepared using HCl-TiO 2was treated for 12h,a noticeable peak splitting appeared.With increasing reaction time from 12to 24h,the intensity and sharpness of the tetragonal peak splitting increased,indicating an increase in the crystallinity of the tetragonal phase along with anincreaseFigure 8.SEM micrographs of BaTiO 3synthesized using HCl-TiO 2at various temperatures:(a)80,(b)120,(c)180,and (d)200°C.Figure 9.SEM micrographs of BaTiO 3synthesized using HClO 4-TiO 2at various temperatures:(a)80,(b)120,(c)180,and (d)200°C.Figure 10.TEM micrographs of BaTiO 3prepared using (a)HCl-TiO 2,120°C;(b)HCl-TiO 2,180°C;(c)HClO 4-TiO 2,120°C;and (d)HClO 4-TiO 2,180°C.Ind.Eng.Chem.Res.,Vol.47,No.6,20081873in the particle size of BaTiO 3.The results are summarized in Table 3.The tetragonal BaTiO 3can be synthesized from HCl-TiO 2at reaction times of 12and 24h using Ba/Ti ratio of 1.2,synthesis temperature of 160°C,and calcination temperature at 1150°C for 2h.However,even at 3h reaction time,BaTiO 3prepared using HClO 4-TiO 2achieved the cubic -tetragonal phase transformation (Figure 12b).The reason can be ascribed to the acidic nature of HClO 4,which leads to the formation of BaTiO 3with fewer defects so as to stabilize the tetragonal phase apart from the possible influence of Cl -ions present in the reaction mixture.The presence of chloride ions is speculated to influence the diffusion of Ba 2+ions and retard the crystal growth process,thereby stabilizing the tetragonal phase by forming larger crystals.40In the early stage of reaction,chloride ions produce more nuclei and form smaller particles,which grow larger at prolonged time.Sun and Li 41reported that BaTiO 3particles synthesized in the presence of chloride ions are slightly larger than the particles synthesized in the absence of chloride ions,however,with an enhanced tetragonality compared to the latter.The physicochemical properties of TiO 2prepared from HCl and HClO 4make the difference in the properties of BaTiO 3.34Therefore,it is concluded that the BaTiO 3tetragonal phase can be successfully synthesized using HClO 4-TiO 2as the precursor at a [H +]/[Ti 4+]ratio of 1.2,synthesis temperature of 160°C,and calcination temperature of 1150°C (2h).344.ConclusionIn the present study,the morphology and phase transformation of BaTiO 3prepared using HCl-TiO 2and HClO 4-TiO 2with respect to reaction temperature,reaction time,Ba/Ti ratio,and calcination time were investigated.Increase in Ba/Ti ratio,temperature,and reaction time increases the possibility of cubic phase transformations.Well-crystallized tetragonal BaTiO 3powders of high purity were obtained using HCl-TiO 2as the precursor at optimum conditions of Ba/Ti ratio )1.2,temper-Figure 11.XRD patterns of BaTiO 3synthesized at various reaction times using (a)HCl-TiO 2and (b)HClO 4-TiO 2precursors calcined at 900°C (2h).Figure 12.XRD patterns of BaTiO 3synthesized at various reaction times using (a)HCl-TiO 2and (b)HClO 4-TiO 2precursors.Conditions:reaction temperature )160°C;Ba/Ti )1.2;calcination temperature )1150°C (2h).Table 3.Effect of Synthesis Time and Calcination Temperature on the Crystalline Phase of BaTiO 3calcination temp (°C)Ti precursor Ba/Ti ratio synthesis temp (°C)synthesis time (h)9001150crystalline phase HCl-TiO 21.21603--cubic 1.21606--cubic 1.216012--cubic 1.216024--cubic HClO 4-TiO 21.21603--cubic 1.21606--cubic 1.216012--cubic 1.216024--cubic HCl-TiO 21.21603yes -cubic 1.21606yes -cubic 1.216012yes -cubic 1.216024yes -cubic HClO 4-TiO 21.21603yes -cubic 1.21606yes -cubic 1.216012yes -cubic 1.216024yes -cubic HCl-TiO 21.21603-yes cubic 1.21606-yes cubic 1.216012-yes tetragonal 1.216024-yes tetragonal HClO 4-TiO 21.21603-yes tetragonal 1.21606-yes tetragonal 1.216012-yes tetragonal 1.216024-yestetragonal1874Ind.Eng.Chem.Res.,Vol.47,No.6,2008ature)180°C,synthesis time)6h,and calcination at 1150°C for2h.However,the phase transformation of BaTiO3 prepared using HClO4-TiO2occurred at lower reaction temper-ature(160°C)and synthesis time(3h)than BaTiO3particles prepared using HCl-TiO2.BaTiO3particles are agglomerated to a spherical shape with ca.80-90nm and BaTiO3particles synthesized with HClO4-TiO2were smaller than those prepared by HCl-TiO2.The stabilization of cubic BaTiO3is caused by defects including OH-defects and barium vacancies.The formation of tetragonal BaTiO3is promoted by the use of high reaction temperature,calcination temperature,and reaction time, which reduces the probability of forming OH-vacancies.In summary,the precursor has a strong influence on the size and morphology of BaTiO3.BaTiO3prepared from HClO4-TiO2 indeed increases the transformation of cubic-to-tetragonal phase at lower reaction conditions without significant particle growth. 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