CUSP磁场对8英寸半导体级硅晶体生长的影响(FEMAG晶体生长仿真软件)

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FEMAG整体介绍

FEMAG整体介绍
熔体内氧浓度分布情况
仿 真 智 领 创 新
Simulating inspires innovation
FEMAG/CZ软件的主要功能
• 与时间相关的动态模拟与缺陷和空位浓度控制
全局动态模型预测每个阶段的缺陷浓度(Ci-Cv)分布
仿 真 智 领 创 新
Simulating inspires innovation
Simulating inspires innovation
FEMAG公司简介
FEMAG公司成立于2003年,总部位于比利时
FEMAG致力于为全球用户开发专业的晶体生长多场耦合仿 真分析工具,为全球晶体设备供应商、晶体材料科研机构 提供设计、优化其生长工艺过程的帮助
• • • • • 比利时新鲁汶大学教授 Franç ois Dupret FEMAG公司创始人和首席科学家 国际晶体生长模型构建与仿真的奠基人 第二届晶体生长模型国际研讨会主席 曾任Journal of Crystal Growth主编
质量控制工程
成本控制工程
仿 真 智 领 创 新
Simulating inspires innovation
晶体生长仿真难点
晶体生长是一个复杂的过程,涉及:
多物理场耦合(传热、流体、磁场、热应力、传质) 多尺度耦合(多时间尺度、多空间尺度) 高度非线性过程
相变,材料参数高度温度依赖 熔体流动,考虑多种对流效应
FEMAG 产品线
仿 真 智 领 创 新
Simulating inspires innovation
FEMAG/CZ软件
• 模拟提拉生长工艺(直拉法, Czochralski 法, Cz 法,柴氏法)
• 适用于半导体单晶 Si 、 Ge 、太 阳能光伏单晶 Si 、 YAG 、小尺 寸蓝宝石等晶体提拉生长工艺 过程的2D/3D全局数值模拟 • FEMAG/CZ软件包括 CZ基本模 块与CZ/TMF模块

全球半导体晶体生长仿真著名商业软件FEMAG-建模策略

全球半导体晶体生长仿真著名商业软件FEMAG-建模策略

2
Introduction
How to improve the growth process in terms of:
- crystal quality ? - process yield ? - energy consumption ? - production rate ?
FEMAGSoft © 2013
FEMAGSoft © 2013
1. Numerical strategy (cont’d)
Inverse QS and TD simulation of the Global temperature growth of a 300 field mm silicon crystal
t 1800 1740 1680 1620 1560 1500 1440 1380 1320 1260 1200 1140 1080 1020 960 900 840 780 720 660 600 540 480 420 360 300
• To resort to appropriate and up-to-date numerical simulation techniques to couple and solve these models
→ quasi-steady and dynamic models
FEMAGSoft © 2013
Numerical strategy (cont’d)
Stream function Temperature field
FEMAGSoft © 2013
Numerical strategy (cont’d)
FEMAG-1 timedependent simulation of Czochralski Ge growth

碳化硅同质外延质量影响因素的分析与综述

碳化硅同质外延质量影响因素的分析与综述

第53卷第2期2024年2月人㊀工㊀晶㊀体㊀学㊀报JOURNAL OF SYNTHETIC CRYSTALSVol.53㊀No.2February,2024碳化硅同质外延质量影响因素的分析与综述郭㊀钰1,2,刘春俊1,张新河2,沈鹏远1,张㊀博1,娄艳芳1,彭同华1,杨㊀建1(1.北京天科合达半导体股份有限公司,北京㊀102600;2.深圳市重投天科半导体有限公司,深圳㊀518108)摘要:碳化硅(SiC)外延质量会直接影响器件的性能和使用寿命,在SiC器件应用中起到关键作用㊂SiC外延质量一方面受衬底质量的影响,例如衬底的堆垛层错(SF)会贯穿到外延层中形成条状层错(BSF),螺位错(TSD)会贯穿到外延层中形成坑点或Frank型层错(Frank SF)等㊂另一方面受到外延工艺的影响,如在外延过程中衬底的基平面位错(BPD)受应力等条件作用会滑移形成Σ形基平面位错(Σ-BPD),衬底的TSD或刃位错(TED)会衍生为腐蚀坑(Pits),以及新产生SF和硅滴等㊂因此,获得高质量的SiC外延晶片需要从优选SiC衬底和优化外延工艺两方面入手㊂本文对外延生长过程中晶体缺陷如何转化并影响器件性能进行了系统分析和综述,并基于北京天科合达半导体股份有限公司量产的高质量6英寸SiC衬底,探讨了常见缺陷,如BPD㊁层错㊁硅滴和Pits等的形成机理及其控制技术,并对Σ-BPD的产生机理和消除方法进行研究,最终获得了片内厚度和浓度均匀性良好㊁缺陷密度低的外延产品,完成了650和1200V外延片产品的开发和产业化工作㊂关键词:碳化硅;同质外延;外延生长;缺陷;位错;小坑中图分类号:O78;O484;O47㊀㊀文献标志码:A㊀㊀文章编号:1000-985X(2024)02-0210-08 Analysis and Review of Influencing Factors of SiCHomo-Epitaxial Wafers QualityGUO Yu1,2,LIU Chunjun1,ZHANG Xinhe2,SHEN Pengyuan1,ZHANG Bo1,LOU Yanfang1,PENG Tonghua1,YANG Jian1(1.Beijing TankeBlue Semiconductor Co.,Ltd.,Beijing102600,China;2.Shenzhen MITK Semiconductor Co.,Ltd.,Shenzhen518108,China)Abstract:The performance and lifetime of silicon carbide(SiC)devices are directly affected by the quality of SiC epitaxial films.On the one hand,the quality of SiC epitaxial films is affected by the quality of substrates.For examples,the stacking faults(SF)in substrates penetrate into the epitaxial layer,forming bar-shaped stacking faults(BSF),and the threading screw dislocation(TSD)penetrate into the epitaxial layer to form pits or Frank-type stacking faults(Frank SF).On the other hand, the quality of SiC epitaxial films is also influenced by the epitaxial growing process.For examples,basal plane dislocation (BPD)in the substrate formΣ-basal plane dislocation(Σ-BPD)in the epitaxial layer under thermal stress or other unstable conditions,the TSD and threading edge dislocation(TED)in the substrate may be etched and derived into pits,and SF and silicon droplets may also be produced.Therefore,high quality SiC substrates and optimized epitaxial growing process are both crucial for obtaining high-quality silicon carbide epitaxial wafers.In this article,based on the SiC epitaxial films grown on 6inch SiC substrates batch-produced by TankeBlue Company,the defects reproducing process in substrates during epitaxial growing were analyzed,and the formation mechanism and controlling technology of common defects such as BPD,SF,silicon droplets and pits were overviewed.The generation mechanism ofΣ-BPD and its eliminating methods were also explored. Finally,we obtained the mass-production technologies of SiC epitaxial films with good thickness and concentration uniformity, and low defect density,which are qualified for making650and1200V SiC-based MOSFETs.Key words:SiC;homo-epitaxial;epitaxial growth;defect;dislocation;pit㊀㊀收稿日期:2023-05-29㊀㊀基金项目:北京市科协卓越工程师培养计划㊀㊀作者简介:郭㊀钰(1983 ),女,辽宁省人,博士,教授级高工㊂E-mail:guoyu03201@㊀㊀通信作者:刘春俊,博士,研究员㊂E-mail:liuchunjun@㊀第2期郭㊀钰等:碳化硅同质外延质量影响因素的分析与综述211㊀0㊀引㊀㊀言SiC作为目前被广泛关注的第三代半导体材料,具有高击穿电压㊁高电子迁移率㊁高热导率等特性,由其制备的半导体器件相比传统的硅(Si)基半导体器件拥有体积小㊁开关损耗低㊁功率密度更高等优势㊂随着绿色能源革命对电力电子器件耐高压㊁低功耗需求的日益迫切,以及电动汽车㊁充电桩等新兴应用的蓬勃发展, SiC器件在智能电网㊁电动汽车㊁轨道交通㊁新能源并网㊁开关电源㊁工业电机和白色家电等领域展现出良好的发展前景和巨大的市场潜力㊂与传统硅功率器件制作工艺不同,SiC功率器件不能直接制作在SiC单晶材料上,必须在导通型SiC单晶衬底上使用外延技术生长出高质量的外延材料,然后在外延层上制作各类器件㊂之所以不直接在SiC衬底上制作SiC器件,一方面是由于衬底的杂质含量较高,且电学性能不够好㊂另一方面是掺杂难度大,即使采用离子注入的方式,也需要后续的高温退火,远不如在外延层上的掺杂效果好㊂因此,制造出外延层的掺杂浓度和厚度符合设计要求的SiC器件至关重要㊂常见的SiC外延技术有化学气相沉积(chemical vapor deposition,CVD)㊁液相外延生长(liquid phase epitaxy,LPE)㊁分子束外延生长(molecular beam epitaxy,MBE)等,目前CVD是主流技术,具备较高生长速率㊁能够实现可控掺杂调控等优点㊂CVD外延生长通常使用硅烷和碳氢化合物作为反应气体,氢气作为载气,氯化氢作为辅助气体,或使用三氯氢硅(TCS)作为硅源代替硅烷和氯化氢,在约1600ħ的温度条件下,反应气体分解并在SiC衬底表面外延生长SiC薄膜㊂目前国内外SiC外延技术已经取得较大进展,产业界也已成功实现6英寸(1英寸=2.54cm)SiC外延批量生产㊂国外产业化公司主要有美国Wolfspeed公司㊁II-VI公司,日本的Showa Ddenko公司等,国内有厦门瀚天天成电子科技有限公司㊁东莞天域半导体有限公司㊁河北普兴电子科技股份有限公司㊁三安集成等㊂2022年美国Wolfspeed公司已成功实现8英寸SiC外延产品的量产㊂市场上主流的量产产品主要是650㊁1200㊁1700V金氧半场效晶体管(metal-oxide-semiconductor field-effect transistor,MOSFET)器件用6英寸外延产品㊂本研究团队基于十多年在SiC衬底材料制备技术研究和产业推广经验的积累,2022年开始启动SiC外延技术研发,重点针对1200V车规级MOSFET器件用SiC外延材料进行研发和产业化工作㊂本文首先介绍了SiC外延的研究历史,然后结合本团队SiC外延产品相关研发工作综述了SiC外延掺杂浓度控制和缺陷控制方面的研究进展,最后对国产SiC外延的发展进行了总结和展望㊂1㊀研发历史SiC同质外延技术研究需要基于SiC衬底开展,因此研发时间晚于SiC衬底,最早开始于20世纪60年代㊂研究人员主要采用了液相外延法[1-3]和CVD法进行SiC同质外延[4-9]㊂但由于SiC存在200多种晶体结构,外延生长时存在严重的多型夹杂问题,因此早期获得的外延材料质量都很差,这也制约了SiC器件的发展㊂第一个突破性的里程碑是在1987年,Kuroda等[10]和美国Kong等[11]各自相继提出了台阶流外延生长模型,在6H-SiC衬底上进行完美多型体复制,并给出了最优偏离晶向和偏角㊂具体来说,代表SiC晶型的堆垛顺序信息主要在SiC衬底表面台阶的侧向,通过SiC衬底表面偏角度的控制,使得同质外延在衬底表面原子台阶处侧向生长,从而继承衬底的堆垛次序,通过台阶流生长实现晶型的完美复制㊂这项技术同样适用其他晶型,如4H-SiC㊁15R-SiC的同质外延生长㊂4H-SiC同质外延的成功促进了SiC基肖特基二极管的研发,带动了4H-SiC在功率器件应用领域独特的发展㊂第二个标志性里程碑是热壁(温壁)CVD反应室设计,传统冷壁CVD反应腔室[12-13]结构较为简单,但存在一些缺点,如晶片表面法线方向的温度梯度非常大,导致SiC晶片翘曲比较严重[14];另外冷壁CVD加热效率比较低,热辐射损耗严重㊂通过热壁CVD反应室设计,腔室内温度梯度得到显著降低,容易实现良好的温度均匀性,这对于产业化生产至关重要㊂第三个里程碑是氯基快速外延生长技术,传统SiC的CVD生长技术通常使用硅烷和碳氢化合物作为反应气体,氢气作为载气,气相中Si团簇容易形成Si滴,导致外延生长工艺窗口相对较窄,同时也限制了外延生长的速率㊂通过引入氯基化学成分(通常有TCS,或者HCl)可以极大地抑制Si团簇,目前已成功应用于212㊀综合评述人工晶体学报㊀㊀㊀㊀㊀㊀第53卷SiC 快速外延生长中[15]㊂近年来,SiC 外延技术逐渐成熟,产业研究重点关注外延材料掺杂浓度控制和缺陷控制两个方面㊂2㊀SiC 外延层的掺杂浓度控制SiC 是优秀的宽禁带半导体材料,其优点是可以相对容易地在一个宽的范围内控制n 型和p 型掺杂㊂氮(N)或磷(P)用于n 型掺杂,而铝(Al)常用于p 型掺杂㊂硼(B)也曾用作p 型掺杂,但其电离能较大(约350MeV)[16],现在已经不是p 型掺杂的首选㊂Larkin 等[17-19]发现的竞位效应是实现SiC 掺杂控制的关键㊂N 原子替位C 原子位置,而P㊁Al 和B 替位Si 原子位置㊂因此,低C /Si 比有利于提高N 掺杂效率,高C /Si 比不利于N 的掺入;对于Al 和B 则刚好相反㊂目前大部分SiC 器件是基于n 型外延材料制作,氮气也是普遍采用的掺杂气体,N 掺杂与氮气流量㊁生长温度和压力㊁C /Si 比㊁生长速率等参数的依赖关系已有详细的研究[20-22],可以实现N 掺杂浓度大范围的调控(1ˑ1014~2ˑ1019cm -3)㊂对于大尺寸SiC 外延材料,SiC 外延层掺杂浓度的均匀性(δ/mean)是研究及产业界目前关注的另一重点㊂2011年Burk 等采用热壁气相外延(vapour phase epitaxy,VPE)炉制作出了厚度均匀性和浓度均匀性分别低于1.6%和12%的6英寸SiC 外延片[23],2014年Thomas 等在2800W 设备上获得了厚度和浓度均匀性分别低于1.5%和8%㊁良品率97.5%的外延片[24-25]㊂8英寸外延片方面,Mattia 和Danilo 等各自在PE1O8设备上获得厚度和浓度均匀性均低于2%的外延片[26]㊂在水平式外延生长中,气体高速流入生长腔室,中心流速高,两侧接近生长腔室边界的地方流速降低;同时在气体流动的方向上,随着反应气体的消耗,反应气体的浓度降低,这些现象会引起SiC 外延层厚度和浓度的不均匀,进而影响器件的性能㊂解决上述问题的方法是设计适当的反应腔室结构,从进气端到尾气端的反应腔室逐步变窄,使得气体的速度沿着流动方向增加,同时反应气体向衬底的扩散距离减小,抵消气体消耗和边界流速降低带来的影响㊂另外,通过调整SiC 衬底的旋转速度,使用适当比例的氩气和氢气的混合气体作为旋转气体源,调整反应气体中的C /Si 比例,调整中路和旁路的反应气体和掺杂气体的流量,都可以获得更加均匀的载流子浓度和厚度分布[27]㊂图1㊀量产外延片的载流子浓度均匀性(a)和厚度均匀性(b)分布统计Fig.1㊀Uniformity of doping (a)and thickness (b)of epi-wafers 本团队采用水平式外延生长方法,三氯氢硅和乙烯作为反应气源,氮气作为掺杂气体,氢气作为载气,氢气和氩气作为驱动托盘旋转的气源,生长厚度适用于1200V 的SiC 基MOSFET 用SiC 外延层㊂通过调整掺杂氮气在中心和边缘分布比例㊁托盘旋转的速度以及旋转气体中氩气与氢气的比例,优化外延工艺的C /Si 比等生长参数,实现SiC 外延层掺杂浓度及均匀性的有效控制,图1是量产1000片的厚度和浓度均匀性统计数据,C /Si 比在1.0~1.2㊁温度在1600~1650ħ和压力在100mbar 的工艺条件下,统计的外延产品100%达到厚度均匀性小于3%㊁浓度均匀性小于6%㊂3㊀SiC 外延层的缺陷控制研究根据晶体缺陷理论,SiC 外延材料的主要缺陷可归纳为4大类:点缺陷㊁位错(属于线缺陷)㊁层错(属于面缺陷)和表面缺陷(属于体缺陷)㊂3.1㊀点缺陷SiC 外延材料的点缺陷主要有硅空位㊁碳空位㊁硅碳双空位等缺陷[28-30],它们在禁带中产生深能级中心,影响材料的载流子寿命㊂在轻掺杂的SiC 外延层中,点缺陷产生的深能级中心浓度通常在5ˑ1012~2ˑ1013cm -3,与外延生长条件特别是C /Si 比和生长温度相关㊂3.2㊀位㊀错SiC 材料的位错包括螺位错(threading screw dislocation,TSD)㊁刃位错(threading edge dislocation,TED)㊀第2期郭㊀钰等:碳化硅同质外延质量影响因素的分析与综述213㊀和基平面位错(basal plane dislocation,BPD)㊂微管是伯氏矢量较大的螺位错形成的中空管道,可认为是一种超螺位错㊂SiC外延层的位错缺陷基本都和衬底相关,图2是SiC外延层中观察到的典型位错演变图[31-32]㊂大部分微管和螺位错会复制到外延层中,在合适的工艺条件下,部分微管分解为单独的螺位错,形成微管闭合[33],只有一小部分TSD(<2%)转为Frank型层错[34-35]㊂衬底TED基本都会复制到外延层中㊂图2㊀4H-SiC外延层中位错演变图Fig.2㊀Schematic illustration of dislocation evolution process in4H-SiC epitaxial layerBPD位错主要源于衬底中BPD向外延层的贯穿,通常偏4ʎ4H-SiC衬底中大部分BPD位错(>99%)在外延过程中会转化为TED位错,只有少于1%左右的BPD会贯穿到外延层中并达到外延层表面㊂在后续器件制造中,BPD主要影响双极型器件的稳定性,如出现双极型退化现象[36-40]㊂在正向导通电流的作用下, BPD可能会延伸至外延层演变成堆垒层错(SF),造成器件正向导通电压漂移㊂由于刃位错对器件性能的影响要小得多,所以提高SiC外延生长过程中BPD转化为TED的比例,阻止衬底中的BPD向外延层中延伸对提高器件的性能十分重要㊂对于BPD向TED的转化技术已经有比较多的研究报道,例如,外延生长前的KOH刻蚀或氢气刻蚀优化表面[41]㊁外延生长间断[42],或者提高生长速率,结合这些技术,转化率已经提升到99.8%,甚至达到100%[43]㊂此外生长过程中,在应力等条件作用下,BPD很容易在衬底和外延层界面上沿着台阶流法线方向发生滑移,形成界面位错(interfacial dislocations)[44-45],滑移方向取决于BPD的伯氏矢量及应力方向㊂特定条件下,成对BPD同时发生滑移,会形成Σ-BPD㊂在本团队研发过程中也观察到过该缺陷,其典型形貌如图3所示,光致发光检测BPD形貌如图3(a)所示,对外延片进行KOH腐蚀后形貌如图3(b)所示,可以看到一个Σ-BPD包含两条界面位错,其长度可以达到毫米级,在其尾部存在两个BPD㊂Σ-BPD形成机理示意图如图3(c)所示[46-47],其起源于衬底的BPD对,其伯氏矢量方向刚好相反,滑移过程中形成两条界面位错和2个半环位错(half-loop arrays,HLAs)㊂半环位错的长度不一,决定于其驱动力大小,影响滑移的驱动力主要是温场的不均匀性㊂图3㊀Σ-BPD的形貌图(a)㊁氢氧化钾腐蚀坑图(b)和形成机理示意图(c)Fig.3㊀Morphology(a),etched image by KOH(b)and schematic illustration of formation mechanism(c)ofΣ-BPD针对外延BPD,本文在快速外延生长的基础上优化外延层缓冲层工艺窗口,目前可以实现BPD密度小于0.1cm-3的外延层批量制备,如图4所示㊂3.3㊀层错缺陷SiC外延层中的层错包括两大类:一类来源于衬底的层错和位错缺陷,衬底的层错会导致外延层形成214㊀综合评述人工晶体学报㊀㊀㊀㊀㊀㊀第53卷图4㊀外延片的BPD 分布(a)及其统计(b)Fig.4㊀Distribution of BPD (a)and its statistics (b)of epi-wafers Bar-shaped SFs [48-49],衬底的部分TSD 如3.2所述会形成Frank SFs;另一类层错为生长层错(in-grown SFs),是外延生长过程中产生的,与衬底质量没有关系㊂目前,大多数外延层错属于第二类,这些层错中绝大部分为Shockley SFs,是通过在基平面中的滑移产生的[50-51]㊂这些层错缺陷都会对器件性能产生不利影响,例如漏电流的增加㊂降低外延生长速率㊁原位氢气刻蚀优化㊁增加生长温度㊁改善衬底质量都可以有效降低层错数量,本研究团队已经可以提供Shockley SFs 密度小于0.15cm -2的6英寸SiC 衬底㊂3.4㊀表面缺陷SiC 外延层表面缺陷尺度比较大,一般通过光学显微镜可以直接观察到,包括掉落物[52]㊁三角形缺陷[53-54]㊁ 胡萝卜 缺陷[55-56]㊁彗星缺陷[57]㊁硅滴[58]和浅坑[59-60]㊂掉落物主要由反应室的部件上形成的SiC 颗粒脱落形成,通过定期清理或更换反应室部件能够有效控制㊂其他几种表面缺陷的形成机制目前已经有了较多研究,虽然不能形成统一的模型,但是大部分与衬底表面状态(包括划痕/损伤层㊁颗粒沾污㊁凹坑)㊁衬底位错(特别是TSD)等缺陷存在一定的关联性㊂由于台阶流生长模式的放大作用和位错转化的综合效应,导致缺陷形成各种宏观表面形貌特征㊂表面缺陷与器件性能的影响目前也已经有了较多的研究报道,除浅坑缺陷外,其他表面缺陷基本都会对器件的性能产生一定的不利影响,导致器件击穿电压降低或者反向漏流增加[61]㊂浅坑(Pits)是4H-SiC 外延层表面出现在TSD 位错顶端的小凹陷或小坑状的形貌缺陷,其宽度尺度小于10μm㊂TED 在外延层表面引起的小坑尺寸远小于TSD 诱发的小坑尺寸,很难被观察到㊂图5是本团队在外延生长中观察到的典型浅坑AFM 形貌,在台阶流动方向的上游端,小坑缺陷有陡峭的倾斜侧面,在下游端,侧面相对平缓,通过AFM 可以看到Pits 宽度为2μm,深度为4nm,深宽比约为0.002㊂Ohtani㊁Noboru 等则利用TUNA 技术研究了Pits 和Large Pits 的产生机理,认为宽度在几微米㊁深度在14nm 左右的Large pits 是由TSD 产生,而宽度在1μm㊁深度在3~4nm 的Pits 由TED 产生[62-63]㊂近年来,有研究表明:当存在浅坑时,由于几何效应会导致局部电场集中,对于二极管特性基本不存在负面影响㊂Kudou 等[64]研究了Pits 缺陷对SiC 器件的影响,认为Pits 密度不会影响SBD 的漏电流和MOSFET 的TDDB 栅氧可靠性㊂同时指出深宽比较小(小于0.02)的Pits对SBD 和MOSFET 的影响较小㊂图5㊀外延表面宽度和深度分别为2μm 和4nm 的浅坑的AFM 照片Fig.5㊀AFM image of a pit with 2μm width and 4nm depth降低Pits 的主要途径包括:优选TSD 数量较少的优质衬底㊁降低碳硅比和降低外延生长速率㊂目前市场上主要的商业化衬底中TSD 的密度小于1000cm -2㊂本研究团队已经可以提供TSD 密度小于300cm -2的6英寸SiC 衬底㊂通过采用优质衬底,调整外延工艺,可以将Pits 数量从103降低到50以内㊂综合来看,SiC 外延层缺陷一方面取决于衬底结晶质量以及表面加工质量,另一方面受制于外延生长工艺窗口的优化,需要综合考虑各种缺陷的调整方案,例如提高外延生长速率会导致BPD 向TED 转化率的提㊀第2期郭㊀钰等:碳化硅同质外延质量影响因素的分析与综述215㊀高,但会导致层错密度的增加㊂基于本研究团队量产的高质量6英寸SiC衬底,本团队通过大量的实验研究,可以有效控制住SiC外延的各种缺陷,完成650和1200V外延片产品开发和产业化工作㊂图6是典型的650和1200V外延片产品缺陷mapping图,3mmˑ3mm良品率分别为98.9%和97.3%㊂图6㊀650和1200V外延片产品缺陷mapping图Fig.6㊀Mapping diagram of defects in650and1200V epi-wafers4㊀结语与展望SiC外延在产业链中起着承上启下的重要作用,通过不断积累对SiC材料的性能认知和改良,以及器件的不断迭代验证,最终提升外延品质,推动SiC器件的应用㊂本文采用天科合达自有的商业化6英寸衬底,在4H-SiC同质外延过程中,研究了外延层中BPD㊁层错㊁硅滴和Pits缺陷的控制,并对Σ-BPD的产生机理和消除进行研究,最终获得厚度均匀性小于3%㊁浓度均匀性小于6%㊁表面粗糙度小于0.2nm㊁良品率大于96%㊁BPD密度小于0.1cm-2的外延产品㊂目前从本团队的研发进度来看,通过对工艺温度㊁C/Si比和生长速率等参数优化使得浓度和厚度均匀性分别控制在3%和2%以内,BPD的密度可以控制在0.075cm-2以内,但仍需要大量的外延数据进行工艺稳定性验证㊂参考文献[1]㊀BRANDER R W,SUTTON R P.Solution grown SiC p-n junctions[J].Journal of Physics D:Applied Physics,1969,2(3):309-318.[2]㊀IKEDA M,HAYAKAWA T,YAMAGIWA S,et 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N,et al.Long term operation of4.5kV PiN and2.5kV JBS diodes[J].Materials ScienceForum,2001,353/354/355/356:727-730.[38]㊀BERGMAN P,LENDENMANN H,NILSSON PÅ,et al.Crystal defects as source of anomalous forward voltage increase of4H-SiC diodes[J].Materials Science Forum,2001,353/354/355/356:299-302.[39]㊀LENDENMANN H,BERGMAN P,DAHLQUIST F,et al.Degradation in SiC bipolar devices:sources and consequences of electrically active㊀第2期郭㊀钰等:碳化硅同质外延质量影响因素的分析与综述217㊀dislocations in SiC[J].Materials Science Forum,2003,433/434/435/436:901-906.[40]㊀MUZYKOV P G,KENNEDY R M,ZHANG Q,et al.Physical phenomena affecting performance and reliability of4H-SiC bipolar junctiontransistors[J].Microelectronics Reliability,2009,49(1):32-37.[41]㊀SKOWRONSKI M,HA S.Degradation of hexagonal silicon-carbide-based bipolar devices[J].Journal of Applied Physics,2006,99(1):011101.[42]㊀YANG L,ZHAO L X,WU H W,et al.Characterization and reduction of defects in4H-SiC substrate and homo-epitaxial wafer[J].MaterialsScience 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al.Influence of epi-layer growth pits on SiC device characteristics[J].Materials Science Forum,2015,821/822/823:177-180.。

CUSP磁场对8英寸半导体级硅晶体生长的影响(晶体生长建模软件FEMAG)

CUSP磁场对8英寸半导体级硅晶体生长的影响(晶体生长建模软件FEMAG)

Ox yge n ppm a
5 4 3 2 1 0 C U SP N o C U SP
0,12
0
0,02
0,04
Ra dius (m )
0,06
0,08
0,1
0,12
Radial Oxygen Concent rat ion H=1290mm
8 7 6
Radial Oxygen Concent rat ion H=1550 mm
FEMAGSoft © 2012
CRYSTAL RADIAL OXYGEN MEASUREMENT POSITIONS
FEMAGSoft © 2012
PREDICTION OF RADIAL OXYGEN CONCENTRATION
Radial Oxygen Concent rat ion H=90mm
PARAMETERS OBSERVED

The concentration of the Oxygen in the melt and crystal. The concentration of the Phosphorous in the melt and crystal The concentration of the Carbon in the melt and crystal
9 8 7 6
Ox yge n ppm a
5 4 3 2 1 0 0,2 0,4 0,6 0,8 1 1,2 1,4 1,6 C U SP N o C U SP
H e ight (m )
FEMAGSoft © 2012
PREDICTION OF CARBON CONCENTRATION ALONG THE HEIGHT OF THE CRYSTAL

CUSP磁场对8英寸半导体级硅晶体生长的影响(FEMAG晶体生长专业模拟软件)32页PPT

CUSP磁场对8英寸半导体级硅晶体生长的影响(FEMAG晶体生长专业模拟软件)32页PPT
FEMAGSoft © 2019
CUSP vs NOCUSP - CARBON CONCENTRATION - CRYSTAL
FEMAGSoft © 2019
CUSP vs NOCUSP - PHOSPHOROUS CONCENTRATION-CRYSTAL
FEMAGSoft © 2019
CUSP vs NO CUSP - PHOSPHORUS RESISTIVITY - CRYSTAL
FEMAGSoft © 2019
PREDICTION OF PHOSPHOROUS CONCENTRATION ALONG THE HEIGHT OF THE CRYSTAL
FEMAGSoft © 2019
CUSP vs NOCUSP - OXYGEN CONCENTRATION - CRYSTAL
the neutral-axis located on the melt free surface.
SIMULATION PROCESS
The growth process is modeled as Time dependent simulation taking in to account the transient effects taking place during the growth of the crystal
Radius (m)
Radial Oxygen Concent rat ion H=430mm
9 8 7 6 5
Radial Oxygen Concent rat ion H=640mm
9 8 7 6 5
Oxygen ppma
4
4
CUSP
CUSP
3
No CUSP

8英寸半导电型GaAs单晶衬底的制备与性能表征

8英寸半导电型GaAs单晶衬底的制备与性能表征

8英寸半导电型GaAs单晶衬底的制备与性能表征
任殿胜;王志珍;张舒惠;王元立
【期刊名称】《人工晶体学报》
【年(卷),期】2024(53)3
【摘要】本文使用垂直梯度凝固(VGF)法制备了直径超过200 mm的Si掺杂GaAs单晶。

通过多线切割、磨边、研磨、化学机械抛光和湿法化学清洗等加工工序制备出8英寸半导电型GaAs单晶衬底。

使用X射线衍射、位错密度检测、霍尔测试、非接触式表面电阻率测试、光致发光测试和晶圆表面缺陷检测等对8英寸GaAs衬底的晶体质量、位错、电学性能和表面质量等特性进行了测试分析。

结果表明:衬底(400)衍射峰半峰全宽低于0.009°;平均位错密度低于30 cm^(-2),其中,晶体头部平均位错密度为1.7 cm^(-2),且有98.87%的面积位错密度为0;衬底面内电阻率标准差小于6%,面内光致发光强度标准差小于4%,≥0.2μm的表面光点缺陷(LPD)个数小于10。

上述结果表明,所制备的8英寸GaAs衬底质量优异,满足外延器件对高质量衬底的要求。

【总页数】10页(P487-496)
【作者】任殿胜;王志珍;张舒惠;王元立
【作者单位】北京通美晶体技术股份有限公司
【正文语种】中文
【中图分类】O785
【相关文献】
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2.LPCVD法在制绒单晶硅片衬底上制备ZnO∶B透明导电薄膜及其性能的研究
3.Si衬底上热壁外延制备GaAs单晶薄膜材料
4.8英寸导电型4H-SiC单晶衬底制备与表征
5.低位错密度8英寸导电型碳化硅单晶衬底制备
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晶体生长计算软件FEMAG系列之晶体生长方法介绍

晶体生长计算软件FEMAG系列之晶体生长方法介绍

可扩展性
软件具有开放性和可扩展性, 用户可以根据需要添加新的材 料属性和边界条件。
图形界面
提供友好的图形界面,方便用 户进行模型建立、参数设置和 结果分析。
软件应用领域
半导体晶体生长
用于研究半导体晶体生长过程中的物理和化学行 为,优化晶体质量和性能。
光学晶体生长
用于研究光学晶体的生长过程,优化晶体光学性 能和加工工艺。
增强可视化功能
为了更好地帮助用户理解和分析计算结果,FEMag软件将 增加更强大的可视化功能,如3D图形界面、实时渲染等, 使用户能够更直观地查看和操作计算结果。
拓展应用领域和范围
扩大应用领域
随着晶体生长研究的不断发展,FEMag软件的应用领域将不 断扩大。未来,FEMag软件将不仅应用于传统的晶体生长研 究,还将拓展到其他相关领域,如材料科学、化学、生物学 等。
该软件通过建立数学模型,模拟晶体生长过程中各 种因素对晶体形态、结构和性能的影响。
FEMag软件提供了丰富的材料属性和边界条件设置 ,支持多种晶体结构和生长条件。
软件特点
01
02
03
04
高效计算
采用有限元方法进行数值计算 ,能够快速求解大规模的晶体 生长问题。
精确模拟
能够模拟晶体生长过程中的温 度场、浓度场、应力场等物理 场,以及化学反应过程。
专业和深入。
与实验结果的比较
FEMag与实验的一致性
FEMag软件在模拟晶体生长方面取得了与实验结果高度一致的结果。通过对比实验和模拟数据,可以验证 FEMag软件的准确性和可靠性,进一步推动其在晶体生长研究中的应用。
实验验证的局限性
尽管FEMag软件与实验结果具有较好的一致性,但实验验证仍然存在局限性。实验条件和参数的微小变化可能 会对结果产生显著影响,而模拟结果可能无法完全反映这些细微差异。因此,将实验和模拟结果相结合,进行综 合分析是更为可靠的方法。

晶体生长建模软件FEMAG-横向磁场直拉硅晶体生长的全局模拟

晶体生长建模软件FEMAG-横向磁场直拉硅晶体生长的全局模拟

Top: field
velocity
Bottom: temperature field
Growth of a 300 mm crystal under a 500 mT TMF
FEMAGSoft © 2013
Cz Si growth under a TMF (cont’d)
Left: melt surface Right: meridional cross-sections parallel and perpendicular to the magnetic field
This hypothesis is satisfactory because:
- generally 3D components are rotating with respect to the 2D environment
- 3D components mostly view 2D components because of the presence of heat shields
Main modeling hypothesis:
- the viewed and hidden parts are calculated as axisymmetric
- or, equivalently, each surface of the enclosure is viewed as axisymmetric from the other surfaces
FEMAGSoft © 2013
Cz Si growth under a TMF (cont’d)
Flow and global heat transfer in a silicon Cz puller under the effect of a TMF

前驱体转化法制备超高温陶瓷粉体研究进展 

前驱体转化法制备超高温陶瓷粉体研究进展 

第42卷第8期2023年8月硅㊀酸㊀盐㊀通㊀报BULLETIN OF THE CHINESE CERAMIC SOCIETY Vol.42㊀No.8August,2023前驱体转化法制备超高温陶瓷粉体研究进展孙楚函,王洪磊,周新贵(国防科技大学空天科学学院,新型陶瓷纤维及其复合材料重点实验室,长沙㊀410073)摘要:超高温陶瓷(UHTC)在航空航天的热防护领域具有重要作用,高质量的UHTC 粉体是制备高性能UHTC 的重要原料㊂在制备UHTC 粉体的工艺中,前驱体转化法制备的粉体纯度高㊁粒径小㊁各组分分布均匀,具有广阔的应用前景㊂本文根据前驱体合成机理将UHTC 前驱体转化法分为金属醇盐配合物合成法㊁基于格氏反应合成法以及引入支链合成法,综述了近年来通过三种方法制备UHTC 粉体的研究进展,分析总结了三种方法的优缺点,指出了UHTC 前驱体转化法目前存在的问题以及未来发展方向㊂关键词:前驱体转化法;超高温陶瓷粉体;反应机理;碳热还原;陶瓷产率;微观结构中图分类号:TH145㊀㊀文献标志码:A ㊀㊀文章编号:1001-1625(2023)08-2865-16Research Progress on Ultra-High Temperature Ceramics Powder Prepared by Precursor-Derived MethodSUN Chuhan ,WANG Honglei ,ZHOU Xingui(Science and Technology on Advanced Ceramic Fibers and Composites Laboratory,College of Aerospace Science and Engineering,National University of Defense Technology,Changsha 410073,China)Abstract :Ultra-high temperature ceramics (UHTC)plays an important role in the field of thermal protection in aerospace.High quality UHTC powder is important raw material for the preparation of high performance UHTC.In the process of preparing UHTC powder,the powder prepared by precursor-derived method has high purity,small particle size and uniform distribution of component,so it has broad application prospects.According to the synthesis mechanism of precursor,the precursor-derived methods of UHTC were divided into metal alkoxides complex synthesis method,synthesis based on Grignard reaction method and synthesis by introducing branch chains method.The research progress of preparation of UHTCby three methods in recent years was reviewed.The advantages and disadvantages of three methods were analyzed and summarized.The existing problems and future development direction of the UHTC powder prepared by precursor-derived method were pointed out.Key words :precursor-derived method;ultra-high temperature ceramics powder;reaction mechanism;carbothermic reduction;ceramic yield;microstructure 收稿日期:2023-04-12;修订日期:2023-05-30作者简介:孙楚函(2001 ),男,硕士研究生㊂主要从事超高温陶瓷的研究㊂E-mail:151****6953@通信作者:王洪磊,博士,副教授㊂E-mail:honglei.wang@0㊀引㊀言近年来,航空航天技术快速发展,先进飞行器正朝着高机动㊁轻质化㊁低成本和可重复使用等方向发展[1],其发动机热端㊁鼻锥和机翼前缘等部件往往要承受2000ħ甚至3000ħ以上的高温,同时还将处于高温氧化㊁热疲劳和高应力等恶劣服役条件下[2-5],传统的难熔合金材料难以满足使用要求,而超高温陶瓷(ultra-high temperature ceramics,UHTC)因其优良的性能已成为该领域的研究重点[6-8]㊂超高温陶瓷一般是指熔点超过3000ħ,且在高温㊁高载荷等极端环境下仍能保持物理及化学性能稳定的过渡金属化合物,主要包括第IVB 族和第VB 族的钛(Ti)㊁锆(Zr)㊁铪(Hf)和钽(Ta)的硼化物㊁氮化物和碳化物[9-10]㊂表1列出了常见UHTC 的物理及力学性能[10-29](HCP 为密排六方结构,FCC 为面心立方结构)㊂2866㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷表1㊀常见超高温陶瓷的物理及力学性能Table1㊀Physical and mechanical properties of common ultra-high temperature ceramicsMaterial Crystalstructure Meltingpoint/ħDensity/(g㊃cm-3)CTE/(10-6㊃K-1)Thermalconductivity/(W㊃m-1㊃K-1)Elasticmodulus/GPaHardness/GPa ReferenceTaB2HCP304012.58.54155126[12-15] TiB2HCP3225 4.58.66556025[11-13,16] ZrB2HCP3245 6.1 6.26048923[12-13,17-18] HfB2HCP338011.2 6.610448028[12-13,17-18] TiC FCC3100 4.97.41740026[13,16,19-20] ZrC FCC3530 6.6 6.72036933[17,19-23] TaC FCC388014.5 6.32250322[17,24-25] HfC FCC389012.7 6.62245229[17,19-23] TaN FCC308713.4 3.2849010[10,26-27] TiN FCC2950 5.49.32946021[10,13,19,26,28] ZrN FCC29507.37.22039016[10,13,19,26,28-29] HfN FCC338513.9 6.92138516[10,13,19,26,28]㊀㊀Note:CTE,coefficient of thermal expansion.高质量UHTC粉体是制备高性能UHTC的关键,UHTC粉体的传统合成工艺是利用相应的金属氧化物粉体经碳热还原反应实现的㊂但原料颗粒的尺寸较大㊁反应物无法充分接触以及可能存在杂质等因素,导致反应温度较高㊁产物晶粒尺寸过大㊁纯度不高等问题,使其应用存在较大的局限性㊂近年来被广泛研究的前驱体转化法是通过化学手段在溶液体系中合成一类包括合成陶瓷时所需元素的金属有机聚合物,再将前驱体在一定温度范围进行交联㊁热解,最终得到陶瓷粉体产物的方法㊂前驱体转化法可对前驱体分子结构进行设计,且在制备过程中具有很好的加工性,可应用于制备陶瓷粉体㊁纤维㊁涂层和复合材料等[30]㊂由于原料组分可以在分子层面均匀混合,缩短元素间的扩散距离,进而降低热解温度,这避免了晶粒粗大的问题,且使产物的相组成分布均匀㊂前驱体转化为陶瓷粉体主要包含两个过程:1)在100~400ħ低温条件下的交联过程中,前驱体分子将交联形成不熔的网状结构;2)在600~1400ħ高温条件下的热解过程中,在600~1000ħ时交联的前驱体发生有机-无机转变,生成非晶陶瓷,继续升高热解温度则会发生相分离与结晶化过程,最终得到多晶陶瓷㊂含氧前驱体会额外发生碳热还原反应,将氧化物陶瓷转化为碳化物陶瓷[31]㊂目前合成UHTC前驱体的工艺按照反应机理可大致分为三类:一是采用金属醇盐配合物经水解缩合形成聚合物前驱体;二是以格氏反应为核心合成单体,再经缩合反应得到聚合物前驱体;三是将有机金属化合物单体作为支链引入聚合物,从而得到目标前驱体㊂1㊀金属醇盐配合物前驱体制备UHTC粉体在制备金属醇盐配合物前驱体的过程中,主要采用过渡金属氯化物作为金属源,通过与醇的取代反应得到金属醇盐㊂由于金属醇盐水解剧烈,利用乙酰丙酮等配体与金属醇盐反应形成配合物以实现可控水解缩合,得到聚合物前驱体㊂同时为保证后续碳热还原反应充分,往往还需向前驱体溶液中加入碳源㊂该方法既可以利用单种金属醇盐配合物制备单相高纯UHTC粉体,也可以通过引入多种金属醇盐配合物制备UHTC 固溶体粉体,或引入含Si聚合物制备复相UHTC粉体㊂1.1㊀金属醇盐配合物前驱体制备单相UHTC粉体TaC具有高熔点㊁高硬度和高强度等诸多优点,是超高温碳化物陶瓷的研究热点之一㊂Jiang等[32]以TaCl5为钽源,酚醛树脂为碳源,乙醇和乙酰丙酮为溶剂,混合得到TaC的前驱体溶液㊂随后在80ħ下固化, 200ħ下保温2h除去溶剂,在1000ħ时开始发生碳热还原反应,1200ħ时反应完全,得到的TaC陶瓷粉体元素分布均匀,平均晶粒尺寸为40nm,但陶瓷产率为57%(质量分数),仍有提升空间㊂图1为前驱体合成和热解过程中可能发生的反应(Hacac为乙酰丙酮;acac为失去一个H原子的乙酰丙酮根)㊂第8期孙楚函等:前驱体转化法制备超高温陶瓷粉体研究进展2867㊀图1㊀TaC 前驱体制备可能发生的反应机理[32]Fig.1㊀Possible reaction mechanism for preparation of TaC precursor [32]常规的前驱体碳热还原法包括前驱体合成㊁固化㊁惰性气氛热解以及最终的碳热还原处理等多个步骤,存在反应时间长㊁生产效率低的问题㊂为优化生产工艺,Cheng 等[33]通过高温喷雾热解(high temperature spray pyrolysis,HTSP)工艺,低成本㊁单步合成了纳米TaC 粉体㊂TaC 前驱体溶液由TaCl 5和酚醛树脂溶解在乙醇和1-戊醇中得到,然后通过喷雾器将其破碎成细小的液滴,液滴处在Ar 气氛的高温管式炉中,再经过溶剂一次性去除㊁热解和1650ħ的快速原位碳热还原,在几分钟内即可制得纳米TaC 粉体㊂但由于采用的是医用雾化器,难以产生足够细小的液滴,且部分产物附着在管式炉内壁上,所以生成的TaC 颗粒存在团聚现象,产率较低,工艺流程需要继续改进㊂图2为高温喷雾热解示意图(CTR 为碳热还原反应)㊂图2㊀高温喷雾热解示意图[33]Fig.2㊀Schematic diagram of high temperature spray pyrolysis [33]单相UHTC 的高温抗氧化能力较弱,尤其是过渡金属碳化物表面被氧化后,无法生成致密氧化膜来阻止内部被进一步氧化㊂例如,当HfC 暴露在空气中时,400ħ以上就开始氧化[34],TaC 在850ħ时即会被完全氧化[35]㊂在实际应用过程中,使用单相UHTC 的情况较少㊂1.2㊀金属醇盐配合物前驱体制备UHTC 固溶体粉体为改善TaC 和HfC 的抗氧化性能,Zhang 等[36]系统地研究了Ta-Hf-C 三元陶瓷在1400~1600ħ等温条件下各种成分的氧化机理,研究表明氧化过程取决于成分㊂与单相TaC 和HfC 陶瓷相比,1TaC-1HfC 和1TaC-3HfC 的抗氧化性显著提高,这是因为氧化生成的三维共晶Hf 6Ta 2O 17-Ta 2O 5结构和致密纯Hf 6Ta 2O 17层都能够抑制O 2扩散,改善抗氧化性㊂因此,与单相UHTC 相比,使用钽醇盐配合物与铪醇盐配合物混合得到前驱体所制备的UHTC 固溶体具有更好的抗高温氧化能力[37]㊂在碳热还原过程中,多相氧化物由于各相反应活化能不同,往往会发生某相优先析出㊁碳化物之间固溶不充分和碳源过剩等问题㊂为解决以上问题,蒋进明[38]以Ta㊁Hf㊁Zr 的氯化物为金属源,乙酰丙酮多齿配体为螯合剂,酚醛树脂为碳源,经200ħ溶剂热处理12h,合成出具有多层核壳结构的前驱体㊂前驱体中心区富含Ta㊁次外层富含Hf(Zr),外壳由树脂包覆㊂该结构的前驱体在热解过程中可以实现外层碳原子向内核逐层扩散,使元素分布均匀,得到粒径为200~300nm 的Ta-Hf(Zr)-C 三元陶瓷纳米粉体㊂图3为Ta-Hf(Zr)-C 碳热还原转化机理示意图㊂2868㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷图3㊀Ta-Hf(Zr)-C 碳热还原转化机理示意图[38]Fig.3㊀Schematic diagram for carbothermal reduction synthesis of Ta-Hf(Zr)-C [38]TaC 和HfC 晶体结构相同(均为NaCl 结构)且晶格常数相近(分别为0.445和0.464nm),可以形成不同比例的固溶体,其中Ta 4HfC 5具有目前已知物质中的最高熔点4215ħ[39],是一种极具发展前景的耐超高温陶瓷㊂Cheng 等[40]等以酚醛树脂作为碳源,与摩尔比为4ʒ1的TaCl 5和HfCl 4溶解在乙醇和乙酰丙酮的混合溶剂中,经过磁力搅拌得到Ta 4HfC 5前驱体溶液,随后在Ar 气氛中200ħ油浴交联固化2h,再通过真空蒸馏除去剩余溶剂,接下来在Ar 气氛中进行热解,Ta 2O 5的碳热还原在1000ħ左右开始,1200~1400ħ时,Hf 6Ta 2O 17的碳热还原以及TaC 和HfC 之间的固溶反应同时发生,最后HfC 和TaC 在1800ħ下固溶充分反应,得到粒度为200~300nm㊁元素分布均匀的Ta 4HfC 5粉体㊂高温下生成的熔融Hf 6Ta 2O 17层可作为氧扩散屏障,使得陶瓷具有优秀的高温抗烧蚀性能㊂但1800ħ的固溶温度过高,不利于得到晶粒细小的高质量粉体㊂图4㊀Ta 4HfC 5粉体TEM 照片[42]Fig.4㊀TEM image of Ta 4HfC 5powder [42]改进前驱体合成工艺可以降低HfC 和TaC 发生固溶反应的温度㊂Lu 等[41]利用摩尔比4ʒ1的TaCl 5和HfCl 4与三乙胺㊁甲基叔丁基醚和乙酰丙酮反应后缩聚,得到聚钽铪氧烷(polytantahafnoxane,PTHO),再将其与含烯丙基的树脂混合即得到Ta 4HfC 5前驱体,固化后在1600ħ下热解制备得到了Ta 4HfC 5粉体㊂孙娅楠等[42]则将含烯丙基的树脂替换为酚醛树脂,与PTHO 混合后得到了Ta 4HfC 5前驱体,将前驱体在250ħ下保温2h 以固化,随后在Ar 气氛中1350~1450ħ热解1.5~3.0h,得到粒径为100~200nm㊁晶粒尺寸为25~50nm 的Ta 4HfC 5粉体㊂图4为Ta 4HfC 5粉体的TEM 照片㊂综合以上研究发现,固溶反应发生的温度普遍高于碳热还原反应㊂与Cheng 等[40]和Lu 等[41]相比,孙娅楠等[42]将固溶反应完成温度从1800ħ降至1450ħ,且所得陶瓷粉体粒径更小㊂通过金属醇盐配合物前驱体制备的超高温陶瓷粉体多为碳化物,也可以通过向前驱体溶液中加入硼酸以制备硼化物复相陶瓷粉体㊂IVB 族硼化物陶瓷ZrB 2和HfB 2在高于1200ħ的氧化环境中,表面的B 2O 3保护层将蒸发,因此主要依赖于ZrO 2或HfO 2层作为抗氧化屏障[43-44]㊂在向ZrB 2和HfB 2中添加高价阳离子Ta 5+后,氧化生成的Ta 2O 5可以填充氧晶格空位以减缓O 2传输速率,并与ZrO 2或HfO 2形成中间相,从而增强相稳定性[45]㊂Xie 等[46]采用乙酰丙酮与Zr(OPr)4通过回流生成Zr(OPr)4-x (acac)x ,得到ZrO 2前驱体㊂类似地,使用Ta(OC 2H 5)4作为Ta 源合成Ta 2O 5前驱体,然后在溶液中分别加入酚醛树脂和硼酸,将溶液浓缩㊁干燥获得前驱体粉末后,在800~1800ħ的Ar 气氛中热解,热解过程中金属氧化物优先进行碳热还原生成金属碳化物,在硼源过量的情况下会继续反应生成金属二硼化物㊂图5为ZrB 2-TaB 2在1300ħ热第8期孙楚函等:前驱体转化法制备超高温陶瓷粉体研究进展2869㊀图5㊀ZrB 2-TaB 2在1300ħ热处理2h 的SEM 照片[46]Fig.5㊀SEM image of ZrB 2-TaB 2after heat treated at 1300ħfor 2h [46]处理2h 的SEM 照片㊂ZrB 2和TaB 2之间的固溶反应从1400ħ开始,1800ħ时TaB 2相完全消失㊂与由ZrB 2和TaB 2两相混合的陶瓷粉体相比,固溶体陶瓷粉体在性能上具有哪些差异值得继续研究㊂1.3㊀金属醇盐配合物前驱体制备复相UHTC 粉体另一种提高UHTC 抗氧化性能的方法则是引入SiC,高温下SiC 氧化生成的玻璃相SiO 2可提高多孔结构的金属氧化物致密度,具有良好的抗高温氧化和抗烧蚀性[47]㊂同时两种成分在结晶过程中的相互抑制效应可以起到细化晶粒的作用㊂聚碳硅烷(polycarbosilane,PCS)是一种以Si 和C 交替排列作为聚合物骨架的有机硅化合物,常被用来作为制备SiC 的前驱体[48]㊂Lu 等[49]以三乙胺为共沉淀剂,用TaCl 5㊁正丁醇和乙酰丙酮反应制备得到Ta 2O 5前驱体溶液,将其与PCS 混合后蒸馏得到TaC-SiC 前驱体溶液,前驱体充分交联固化后,在1600ħ的Ar 气氛中热解2h,得到了平均晶粒尺寸50nm 的TaC-SiC 陶瓷粉体㊂图6为1800ħ热解的TaC-SiC 陶瓷粉体的HR-TEM 照片(标尺101/nm 为10个1/nm,下文图17㊁18中标尺含义类似)㊂由图6可知,TaC 和SiC 晶粒以接近球形的形态均匀分散,同时还有少量无定形碳嵌在晶界位置㊂该前驱体合成方法同样适用于IVB 族UHTC,可推广用于制备ZrC-SiC 和HfC-SiC㊂图6㊀1800ħ热解的TaC-SiC 陶瓷粉体的HR-TEM 照片[49]Fig.6㊀HR-TEM images of TaC-SiC ceramics powder pyrolyzed at 1800ħ[49]PCS 的交联主要依靠硅氢化反应,通过向前驱体中加入如二乙烯基苯(divinylbenzene,DVB)等含不饱和C C 键的物质可以进一步提升前驱体的交联程度㊂Cai 等[50]利用该原理,以HfCl 4与异丙醇和乙酰丙酮反应得到铪醇盐配合物,再通过水解得到HfO 2前驱体(polyhafnoxane,PHO),随后将PHO 与PCS 和DVB 混合,控制n (Hf)/n (Si)摩尔比为1ʒ1,交联后在1600ħ下碳热还原得到了元素分布均匀㊁结晶质量高㊁粒径分布窄的HfC-SiC 复相陶瓷粉末㊂图7为HfC-SiC 复相陶瓷的TEM 照片,可以观察到分别属于HfC 和SiC 的晶格条纹㊂由于PHO 的弱极性,其与PCS 和DVB 具有良好的相容性,可以大范围改变n (Hf)/n (Si)摩尔比来调控陶瓷粉体成分㊂合成前驱体的单体中交联位点越多,前驱体越易形成高度交联的三维网状结构㊂每个四乙氧基硅烷(tetraethoxysilane,TEOS)分子中含有四个Si O C 键可供交联,是另一种理想的制备含Si 前驱体的原料㊂Patra 等[51]采用TEOS 与HfCl 4㊁乙酰丙酮㊁对苯二酚反应合成HfC-SiC 前驱体㊂经过回流和固化后,在1500ħ的Ar 气氛中发生碳热还原反应,生成HfC-SiC 陶瓷粉体㊂图8为1500ħ热解的HfC-SiC 前驱体亮场TEM 照片和平均粒径㊂由图8可知,碳热还原所生成的球形HfC 和SiC 颗粒平均尺寸为25~50nm㊂由于对苯二酚和四乙氧基硅烷具有较高的C㊁Si 含量,因此前驱体在热解过程中质量损失较少,1600ħ时陶瓷产率高达65%,具有很好的应用前景㊂PCS 作为SiC 前驱体的缺陷在于其常温下为固态,需要利用二甲苯等有机溶剂将其配制成溶液使用,增2870㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷加了前驱体合成的复杂程度㊂Wang 等[52]采用常温下为液态的低分子量SiC 前驱体(LPVCS)与HfCl 4㊁乙酰丙酮和1,4-丁二醇反应合成了HfC-SiC 前驱体(PHCS)㊂HfO 2和SiO 2的碳热还原主要发生在1400~1600ħ,生成的HfC-SiC 复相陶瓷粉体的SEM 照片和EDS 分析如图9所示㊂与PCS 相比,LPVCS 结构中引入的V4分子具有 CH CH 2基团,可在较低温度下实现自交联,有利于陶瓷产率的提升[53]㊂同时LPVCS 中较高的碳含量可以补偿PHCO 热解产物中碳含量的不足,制备出不含HfO 2和微量游离碳的高性能HfC-SiC 陶瓷㊂图7㊀1600ħ热解的HfC-SiC 粉末TEM 照片[50]Fig.7㊀TEM images of HfC-SiC powder pyrolyzed at 1600ħ[50]图8㊀1500ħ热解的HfC-SiC 前驱体亮场TEM 照片和平均粒径[51]Fig.8㊀Bright-field TEM image and average particle size of HfC-SiC precursor pyrolyzed at 1500ħ[51]第8期孙楚函等:前驱体转化法制备超高温陶瓷粉体研究进展2871㊀图9㊀HfC-SiC 粉末的SEM 照片和EDS 分析[52]Fig.9㊀SEM images and EDS analysis of HfC-SiC powder [52]㊀㊀综上可见,合成金属醇盐配合物前驱体所需的原料结构简单,反应时间较短㊂但由于前驱体中存在氧元素,有可能会导致生成的UHTC 粉体中有氧残留,使陶瓷高温抗氧化性能和机械性能下降㊂另外为防止金属醇盐水解,该反应需全程在惰性气氛中进行,对设备要求较高㊂2㊀基于格氏反应的前驱体制备UHTC 粉体基于格氏反应的前驱体制备工艺主要采用茂金属化合物和含不饱和键的格氏试剂合成单体,再通过与非金属源分子的聚合反应得到前驱体㊂金属醇盐配合物前驱体的各目标元素由不同种聚合物提供,多数通过机械搅拌的方法实现分子间的混合㊂不同的是,基于格氏反应的前驱体中金属源与非金属源在同种聚合物分子中,实现了分子内的混合㊂所合成的聚合物分子包括线型聚合物与网状聚合物㊂2.1㊀线型聚合物前驱体制备UHTC 粉体合成线型聚合物前驱体的原料通常依靠分子两端的基团发生缩聚反应,交联程度相较于网状聚合物更低,可以通过在主链上插入交联位点来减少热解过程中的质量损失㊂Cheng 等[54]在四氢呋喃(tetrahydrofuran,THF)溶剂中利用反-1,4-二溴-2-丁烯与镁反应制备格氏试剂,再与Cp 2HfCl 2和氯甲基三甲基硅烷通过缩聚合成了主链包含Hf C㊁Si C 和 CH CH 基团的线性PHCS 前驱体聚合物,图10为前驱体合成过程中可能发生的化学反应㊂前驱体在经过1600ħ热解后得到了元素分布均匀的HfC-SiC 纳米复合陶瓷粉体㊂前驱体主链中的不饱和 CH 2CH CHCH 2 基团提供了潜在的交联位点或反应位点,可用于后续固化或改性㊂图10㊀PHCS 前驱体合成过程中可能发生的反应[54]Fig.10㊀Reactions that may occur during synthesis of PHCS precursors [54]基于格氏反应的MC-SiC(M =Zr,Hf)前驱体分子结构中往往含有M C Si 键,普遍认为该键是由格氏反应所致㊂Gao 等[55]提出了一种新的前驱体合成机制,该机制基于㊃MgCl 辅助下的活性物质Cp 2Zr(II)的自由基聚合,合成过程如图11所示,首先将二氯二茂锆Cp 2ZrCl 2与Mg 和四氢呋喃在60ħ下搅拌混合2872㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷4h 后冷却,得到活性物质双环戊二烯基锆Cp 2Zr (II),再将Cp 2Zr (II)分别与CH 3Si (CH CH 2)Cl 2和(CH 3)2Si(CH 2Cl)2在110ħ下反应16h,经过冷却过滤并真空浓缩得到了含有[ Zr C Si ]n 主链结构的单源聚合物前驱体聚锆碳硅烷(PZCS-1,PZCS-2)㊂随后将前驱体在N 2气氛中进行热解,SiO 2和ZrO 2相在1000ħ时析出,随着温度继续升高转化为SiC 和ZrC 相,且均匀分布在自由碳基体中,形成ZrC /SiC /C 复合陶瓷㊂由于该前驱体为线型聚合物且不含可作为交联位点的不饱和基团,热解过程中质量损失较为严重,900ħ时陶瓷产率仅有43.9%㊂图11㊀PZCS-2前驱体合成过程[55]Fig.11㊀Synthesis process of the PZCS-2precursor [55]2.2㊀网状聚合物前驱体制备UHTC 粉体与线型聚合物前驱体相比,合成网状聚合物前驱体的原料多含有三个以上的交联位点,前驱体交联程度高,质量损失较少,有利于陶瓷产率的提高㊂Wang 等[56]通过格氏反应将Cp 2ZrCl 2和CH 2 CHMgCl 制成Cp 2Zr(CH CH 2)2,然后将其与B 源H 3B㊃SMe 2混合,利用氢化反应得到网状结构的大分子前驱体聚锆碳硼烷(polyzirconcarborane,PZCB),合成机理如图12所示㊂随后将前驱体放置于Ar 气氛的石墨管式炉中进行热解,1600ħ时碳热还原完全,得到充分结晶且分布均匀的ZrC-ZrB 2陶瓷粉体,继续加热至2200ħ,产物失重仅为2.5%,说明该复相陶瓷粉体具有良好的耐热性㊂在该合成过程中,利用了硼烷分子具有三个反应位点的特征,以其作为骨架合成了网状大分子,使得前驱体充分交联㊂SiBNC 非晶陶瓷在2000ħ仍具有很好的高温稳定性,而引入过渡金属元素可以进一步抑制其在高温下的结晶与氧化[57]㊂龙鑫[58]将锆源(Cp 2ZrCl 2)与格氏试剂(CH 2 CHCH 2MgCl)反应制备得到双官能度的活性单体(PZC),然后引入低分子量聚硼硅氮烷(LPBSZ),PZC 中的C C 键与LPBSZ 中的Si H 发生硅氢化反应,ZrC /SiBNC 前驱体合成机理如图13所示(Me 3Si 为三甲基亚砜)㊂未参与反应的C C 键则为后续交联提供活性位点,最终形成网状结构的ZrC /SiBNC 前驱体㊂随后将前驱体置于Ar 气氛中经过1200ħ热解生成ZrC /SiBNC 陶瓷粉体,其中ZrC 纳米颗粒均匀分散在无定形SiBNC 基体中㊂ZrC 相提高了SiBNC 的第8期孙楚函等:前驱体转化法制备超高温陶瓷粉体研究进展2873㊀热稳定性,经过1800ħ以上高温处理后,ZrC /SiBNC 仍能够保持均匀细小的纳米晶结构,同时SiBNC 也改善了ZrC 的耐高温氧化性能㊂但该前驱体的不足之处在于碳含量过高导致陶瓷粉体产物中含有过量的碳,影响UHTC 的高温抗氧化性能㊂图12㊀PZCB 前驱体合成机理[56]Fig.12㊀Synthesis mechanism of PZCB precursor[56]图13㊀ZrC /SiBNC 前驱体合成机理[58]Fig.13㊀Synthesis mechanism of ZrC /SiBNC precursor [58]基于格氏反应的前驱体制备工艺实现了各目标元素在聚合物分子内的混合,比金属醇盐配合物前驱体混合更加充分,能更好地避免陶瓷产物中元素偏析现象的发生㊂同时原料中不含氧元素,热解过程中不会发生碳热还原反应,能降低热解温度㊂但该工艺的原料结构较为复杂,反应时间较长,为避免合成过程中引入空气中的氧等杂质,反应必须在保护气氛中进行,对设备要求较高㊂3㊀引入支链的前驱体制备UHTC 粉体在制备引入支链的前驱体过程中,需以一种聚合物分子作为主链,再将其他含目标元素的小分子通过反应作为支链连接到主链上㊂常见的作为主链的大分子有聚碳硅烷和聚硅氮烷等,其分子结构中包含大量可与含目标元素的小分子发生交联反应的基团,同时自身足够大的分子量可避免在热处理过程中分解挥发㊂3.1㊀以聚碳硅烷作主链制备UHTC 粉体聚碳硅烷的主链由Si 和C 交替组成,Si 和C 上连接有H 或 CH 2 CH CH 2等基团作为交联位点[48],通过向主链上引入UHTC 组分,热解后可原位生成含SiC 的UHTC 粉体㊂Amorós 等[59]系统性地研究2874㊀陶㊀瓷硅酸盐通报㊀㊀㊀㊀㊀㊀第42卷图14㊀1350ħ热解的SiC-TiC-C 陶瓷粉体的SEM 照片[59]Fig.14㊀SEM image of SiC-TiC-C ceramics powder pyrolyzed at 1350ħ[59]了采用聚二甲基硅烷(polydimethylsiloxane,PDMS)和PCS 与Cp 2MCl(M =Ti,Zr,Hf)反应制备SiC-MC-C 陶瓷粉体的机理和工艺流程㊂与PDMS 相比,PCS 中的Si H 键促进了前驱体的交联,提高了陶瓷产率,金属配合物则通过取代反应连接在前驱体的网状结构中㊂经过900ħ热解后,前驱体转变为非晶态陶瓷,结晶化在1350ħ下基本完成,生成由β-SiC㊁MC 以及自由碳组成的复相陶瓷粉体,但仍有部分非晶态物质存在㊂图14是1350ħ热解所得的SiC-TiC-C 陶瓷粉体的SEM 照片㊂该研究采用同种前驱体转化工艺成功制备出了含IVB 族三种元素碳化物的复相UHTC 粉体,但对热解过程的探究不够深入,1350ħ时结晶尚未完成㊂通过对PCS 进行改性,可以进一步提高前驱体交联程度㊂Yu 等[60]以含烯丙基的聚碳硅烷AHPCS(商品名SMP10)为SiC 源,与TaCl 5的CHCl 3溶液混合后,在真空中加热至160ħ脱除溶剂得到前驱体,前驱体合成过程如图15所示,随后将前驱体在Ar 气氛下的管式炉中进行热解,得到SiC-TaC-C 陶瓷粉体㊂研究发现,随着热解温度升高,前驱体由于发生脱氢耦合反应而失重,在900ħ时聚合物完全转化为非晶陶瓷粉末,1200ħ时TaC 相开始析出,并被非晶态碳薄壳所包裹,形成核壳结构的TaC@C 纳米颗粒,而β-SiC 相则在1400ħ下结晶㊂所得的β-SiC 和TaC 的晶粒尺寸均小于30nm㊂前驱体热解后的游离碳需要通过生成TaC 来消耗,由于没有额外添加碳源,所以需要准确掌握TaCl 5和AHPCS 的比例以保证陶瓷产物中有少量包裹在TaC 晶粒表面的游离碳㊂图15㊀SiC-TaC-C 前驱体合成过程[60]Fig.15㊀Synthesis of SiC-TaC-C precursor [60]在利用引入支链的前驱体制备含N 原子的超高温陶瓷粉体时,Wen 等[61]以AHPCS 为SiC 源,四(二甲氨基)铪(TDMAH)为Hf 源和N 源合成SiHfCN 陶瓷前驱体㊂AHPCS 中的Si H 键可与TDMAH 中的N CH 3键反应生成Si N Hf 键,使Hf 连接到大分子上㊂Si H 键还会与AHPCS 侧链上的烯丙基发生硅氢化反应以增加前驱体交联程度,可能发生的化学反应如图16所示㊂热解后所得UHTC 组分为HfC 0.87N 0.13,其被碳层包裹镶嵌在SiC 基体中,两相的晶粒尺寸均小于100nm㊂2~4nm 厚的碳层可作为扩散屏障,有效。

最全的材料晶体生长工艺汇总

最全的材料晶体生长工艺汇总

最全的材料晶体生长工艺汇总提拉法提拉法又称直拉法,丘克拉斯基(Czochralski)法,简称CZ法。

它是一种直接从熔体中拉制出晶体的生长技术。

用提拉法能够生长无色蓝宝石、红宝石、钇铝榴石、钆镓榴石、变石和尖晶石等多种重要的人工宝石晶体。

提拉法的原理:首先将待生长的晶体的原料放在耐高温的坩埚中加热熔化,调整炉内温度场,使熔体上部处于过冷状态;然后在籽晶杆上安放一粒籽晶,让籽晶下降至接触熔体表面,待籽晶表面稍熔后,提拉并转动籽晶杆,使熔体处于过冷状态而结晶于籽晶上,并在不断提拉和旋转过程中,最终生长出圆柱状的大块单晶体。

提拉法的工艺步骤可以分为原料熔化、引晶、颈缩、放肩、等径生长、收尾等几个阶段。

具体过程如示意图。

提拉法晶体生长工艺有两大应用难点:一是温度场的设置和优化;二是熔体的流动和缺陷分析。

下图为提拉法基本的温度场设置以及五种基本的熔体对流模式。

在复杂的工艺条件下,实际生产需要调整的参数很多,例如坩埚和晶体的旋转速率,提拉速率等。

因此实际中熔体的温度场和流动模式也更复杂。

下图是不同的坩埚和晶体旋转速率下产生的复杂流动示意图。

这两大应用难点对晶体生长的质量和效率都有很大影响,是应用和科研领域中最关心的两个问题。

通常情况下为了减弱熔体对流,人为地引入外部磁场是一种有效办法,利用导电流体在磁场中感生的洛伦兹力可以抑制熔体的对流。

常用的磁场有横向磁场、尖端磁场等。

下图是几种不同的引入磁场类型示意图。

引入磁场可以在一定程度上减弱对流,但同时磁场的引入也加大了仿真模拟的难度,使得生长质量预测变的更难,因此需要专业的晶体生长软件才能提供可靠的仿真数据。

晶体提拉法有以下优点:(1)在晶体生长过程中可以直接进行测试与观察,有利于控制生长条件;(2)使用优质定向籽晶和“缩颈”技术,可减少晶体缺陷,获得优质取向的单晶;(3)晶体生长速度较快;(4)晶体光学均一性高。

晶体提拉法的不足之处在于:(1)坩埚材料对晶体可能产生污染;(2)熔体的液流作用、传动装置的振动和温度的波动都会对晶体的质量产生影响。

全球半导体晶体生长仿真著名商业软件FEMAG--Optimization of Silicon Ingot Quality

全球半导体晶体生长仿真著名商业软件FEMAG--Optimization of Silicon Ingot Quality

全球半导体晶体生长仿真著名商业软件FEMAGOptimization of Silicon I n go t Q u al it yby the Numerical P r e d i c ti on of Bulk Crystal D e f ec t sF. Loix2*, F. D upr e t1,2*, A. de Potter2, R. R ol i n s ky2, W. L i a ng2, N. Van den B oga e rt21CESAME R e s e a r ch C e nt e r, Uni v e r s i téc at ho li que de L ouva i n,Bât i m e nt EULER, 4 av. Georges L e m aît r e,B-1348 Louva i n-l a-N e uv e,B e l g i um2FEMAGSoft S.A. Company, 7 Rue André Dumont, A x i s Parc, B-1435 Mont-Saint-Guibert, B e l g i umThe growth of S ili c on(Si) i ngot s by the Cz oc hra l s ki(Cz) technique for e l e c t roni c (IC) a ppl i c a t i ons has been governed for more than 50 years by two, somewhat c ont ra di c t ory, t e c hnol ogi c a l obj e c t i ve s.First, the c rys t a l diameter has to be nearly constant and the l a rge s t pos s i bl e according to current market requirements. Second, the product quality has to be pe rfe c t l y c ont rol l e d in terms of c rys t a l defects and composition. On the other hand, grow i ng Cz S i c rys t a l s for phot o-volt a i c(PV) a ppl i c a t i ons requires to m i ni m i z e both the energy c ons umpt i on and the growth dura t i on without, however, gene ra t i ng a too l a rge content of m i c ro-voi ds in the c rys t a l.A c hi e vi ng these goa l s i s by no means easy s i nc e i nc re a s i ng the c rys t a l diameter re qui re s a l a rge r m e l t volume and hence results in a much more complex m e l t flow regime with c om pl i c a t e d heat, momentum and s pe c i e s transport effects, a very s e ns i t i ve s ol i d-li qui d i nt e rfa c e shape (with a l e ss uniform t he rm a l gradient), and in gene ra l an enhanced dynamic s ys t e m behavior. A s i m il a r enhancement of the systemdyna m i c behavior can result from the use of a high pull rate to i nc re a s e the growth speed. Therefore, in gene ra l,de s i gni ng the furnace hot zone will require to i nt roduc e appropriate heat s hi e l ds in order to w e ll-c ont rol the ra di a t i on he a t transfer whi l e a s a t i s fa c t ory m e l t flow pattern can only be obt a i ne d for l a rge diameter c rys t a l s by the a c t i on of transverse or configured m a gne t i c fi e l ds.Moreover the s e l e c t i on of opt i m a l proc e ss parameters (heater power, c rys t a l pulling rate, c rys t a l and c ruc i bl e rot a t i on rates, m a gne t i c fi e l d i nt e ns i ty if any, ambient gas flow rate, etc.) becomes much more difficult in vi e w of the i nc re a s e d system nonl i ne a ri ty and t i m e-depe nde nc y,e s pe c i a ll y during the c ri t i c a l process stages (necking, shouldering, t a il-e nd stage, c rys t a l detachment, …).N one t he l e ss,compared to the high difficulty to address these different t e c hnol ogi c a l i ss ue s,it i s worth observing that huge progress has been achieved in the l a s t decades in s e ve ra l s c i e nt i fi c dom a i ns.F i r s t, the phys i c s of ra di a t i on and c onve c t i on in Cz furnaces, and of defect form a t i on and transport in growing S i c rys t a l s,i s much better known, and hence the m a t he m a t i c a l m ode l s governing Cz S i c ry s t a l growth are better and better e s t a bl i s he d.In spite of the i mport a nt i m prove m e n t s that re m a i n necessary in the m ode li ng of turbulence in the m e l t a nd the ambient gas (i nc l udi ng the m ode li ng of m e l t turbulence under the effect of a m a gne t i c fi e l d) and of the s t ill i ns uffi c i e nt kno w l e dge of the m a t e ri a l parameters governing point-and micro-defect e vol ut i on in S i s i ngl e c rys t a l s,an a l m os t complete picture of the phy s i c s of S i growth today i s a va il a bl e.Secondly, num e ri c a l methods and computers have a l s o quickly progre ss e d s i nc e the de ve l opm e n t ofthe first m ode l s of Cz growth achieved in the 1980’s. Nowadays the qua s i-s t ea dy or t i m e-depe nde nt s i m ul a t i on of the Cz process has become po ss i bl e in an a cc e pt a b l e c om put i ng t i m e,with s uffi c i e nt l y refined meshes to resolve the key de t a il s of the problem, and with appropriate num e ri c a l techniques to handle the system de form i ng ge om e try (which comprises s e ve ra l moving components together with free boundaries such as the m e l t-c rys t a l and m e l t-ga s i nt e rfa c e s).Therefore, having at one’s di s pos a l the appropriate phys i c a l m ode l s, num e ri c a l tools a nd computer hardware, the route i s directly opened to process opt i m i z a t i on by means of num e ri c a l s i m ul a t i on.The obj e c t i ve of the present paper i s to ill us t ra t e how this strategy can be a ppl i e d by use of the FEMAG-CZ software as today co-de ve l ope d by FEMAG Soft S.A. Company and the CESAME research center of the Uni ve rs i téde L ouva i n (Belgium).We will here focus on the S i ingot quality pre di c t i on and i t s opt i m i z a t i on.We present a fully t i m e-depe nde nt m ode l devoted to predict the gl oba l heat transfer in the furnace, the s ol i d-liquid i nt e rfa c e shape, and the re s ult i ng di s t ri but i ons of point-and m i c ro-de fe c t s as c a l c ul a t e d from the S i nno-D ornbe rge r (S-D) model together with an e xt e ns i on of the l um pe d model of Voronkov and Kulkarni. All the t ra ns i e nt s are c ons i de re d including the effects of c rys t a l a nd c ruc i bl e lift, of the heat c a pa c i t i e s of the furnace c ons t i t ue nt s, of the t he rm a l i ne rt i a of the s ol i di fi c a t i on front, and of the dyna m i c defect governing l a w s.We hence show that dynamic effects deeply affect the defect di s t ri but i on inthe c rys t a l(fig 1.). In a ddi t i on to the c l a ss i c a l point-defect e vol ut i on mechanisms, a new l um pe d m ode l i s de ve l ope d to c a l c ul a t e the form a t i on a nd growth of m i c ro-de fe c t s in order to predict their dens i t i e s and s i z e di s t ri but i ons anywhere in the c rys t a l.Another key i ss ue in Cz S i growth i s to control the dens i ty of oxygen and any other s pe c i e s(i nc l udi ng dopants and i m puri t i e s) i ns i de the c rys t a l.M ode li ng i ss ue s will be here a ga i n de t a il e d.Finally, off-line process control pri nc i pl e s will be addressed. Results will ill us t ra t e how this tool can he l p in opt i m i z i ng c rys t a l shape and quality.P r e di c t e d defect de l t a C i-C v di s t r i bu t i on(C i,C v be i ng the c onc e nt r at i on ofi n t e r s t i t i al s , v acan c i e re s p e c t i v e l y) with aquas i-s t e ady (a) and a t i m e-d e pendent (b)simulation. The OSF ring is located at theposition where delta~= 0. This picturehighlights the strong impact on the pointdefect of the transient effects in the growing crystal.。

呼和浩特市科技局赴“科技兴蒙”重点专项实施单位中科汇通(内蒙古)投资控股有限公司调研

呼和浩特市科技局赴“科技兴蒙”重点专项实施单位中科汇通(内蒙古)投资控股有限公司调研

科技动态呼和浩特市科技局组织召开首次内蒙古生物发酵产业技术创新中心建设运行推进会为深入实施首府“两区两中心五基地”科技创 新工程,加快推进生物发酵特色产业基地建设,1月13日,呼和浩特市科技局组织召幵了首次生物发酵 产业技术创新中心建设运行推进会,市科技局副局 长安伟、托县工信与科技局局长蔡艳茹、内蒙古工 业大学教授刘占英、内蒙古蒙戈利智业投资咨询 (集团)有限公司董事长赵景志、市科技局高新科 工作人员参加了会议。

会议深入研究了内蒙古生物发酵产业基础和创 新能力,自治区生物发酵产业技术创新中心组建的 意义、任务、实施计划以及攻关生物发酵工艺及菌 种选育等行业“卡脖子”技术的必要性和可行性。

会议认为,创建生物发酵产业技术创新中心,呼和浩特市具备深厚的产业基础和创新资源,是实现 生物发酵产业高质量发展的关键核心。

下一步将主 要开展以下工作:一是深人研究《内蒙古自治区技 术创新中心建设实施方案(暂行)》,高标准编 制《内蒙古生物发酵产业技术创新中心建设与运行 方案》,要面向整个行业集聚资源组建创新中心,并实现在行业内开放共享;二是依托内蒙古工业大 学等高校科研院所专家学者,立足行业发展实际,进一步对关键核心技术进行聚焦和凝练,以市场为 导向,着重解决行业急需的关键核心问题;三是加 强同自治区科技厅的对接沟通,争取上级部门的支 持;四是成立生物发酵产业技术创新中心筹备领导 小组,建立常态化的工作推进机制,.(成博)呼和浩特市科技局赴“科技兴蒙”重点专项实施单位 中科汇通(内蒙古)投资控股有限公司调研1月26曰,呼和浩特市科技局副局长安伟带领相 关科室工作人员赴“科技兴蒙”重点专项实施单位 中科汇通(内蒙古)投资控股有限公司调研,了解 “科技兴蒙”重点专项的实施进展情况和存在的问 题,并就下一步工作进行探讨。

中科汇通(内蒙古)投资控股有限公司项目组 人员、中科院力学所陈启生研究员介绍了“第三代 半导体材料(碳化硅)晶体生长工艺及制备设备研 发项冃”的进展情况。

材料仿真软件用于晶硅生长实训教学的探索

材料仿真软件用于晶硅生长实训教学的探索

科技风2019年10月科教论坛D0I:10.19392/kg1671-7341.201929042材料仿真软件用于晶硅生长实训教学的探索徐立波常州工程职业技术学院江苏常州213017摘要:通过使用模拟软件,解决了实训中学生操作不规范性,参数设置不恰当,设备选型盲目,性能检测麻烦等问题,软件以其零成本,可操作性强,界面友好,功能强大的优点,在实训过程中可以激发学生的求知欲和兴趣,起到了很好的辅助教学作用,提升了实训效果。

关键词:晶硅生长;FEMAG;网格剖分;温度场分布晶硅生长实训中遇到的困难晶硅生长是一个复杂的过程,其形核生长是微观的过程,不易观察,单晶生长是非均匀形核生长过程,原子在固液界面堆积,在二维表面成核,再侧向层状生长,涉及传热,传质,磁场,流体,热应力等多物理场,以及时间空间的多尺度,还有相变,材料参数变动,熔体流动等掺在一起的非线性过程,实训的目的是从直观感性的角度来验证和理解理论,以及很好的操控整个长晶过程,学生除了动手装料,安装设备,制定工艺参数,分段观察检测,最终目的要对晶硅的质量负责,由于此过程周期长,耗能多,生长过程不直观,凭借经验不能保证长晶过程完善,一旦失败,成本太高。

2仿真软件的特点和软件作用选用FEMAG软件,软件通过多物理场耦合全局有限元分析,并使用高效的流体力学求解器来处理非线性过程,能够对 晶硅的最终质量指标化,从而不需要真实生长单晶,即可从设 定条件模拟出产品,并对界面,温度梯度,缺陷分布,速度场等给出全面结果,通过渲染软件,使其三维可见,可以零成本对晶硅质量进行预判,并对热屏形状材料位置进行设计,对保温层和多加热器进行设计,能够有效控制热应力,氧碳含量,掺杂物浓度,缺陷,成本,寿命,对能耗,气耗,物耗进行计算和控制,在win7下32位即可运行,内存占用少,准确率高达95%以上。

3设置完整合理参数利用软件出结果的做法软件分为前处理,计算,后处理三个模块,其重点是前处理中按照实际生长情形进行参数设置输入,所以影响热场,影响固液界面温度梯度的因素都要考虑到,分几个部分分别设置,包括几何条件、物理条件、操作条件。

基于Filmstar软件的Yb:YAG激光器谐振腔镀膜设计

基于Filmstar软件的Yb:YAG激光器谐振腔镀膜设计

基于Filmstar软件的Yb:YAG激光器谐振腔镀膜设计廖川;贾红宝;郑孟莹;许崇【摘要】从光学谐振腔阈值条件的推导出发,利用光学薄膜设计软件Filmstar为谐振腔的左右腔镜各设计了两种光学薄膜,使1030.5 nm~1055.5 nm波段连续可调谐的Yb:YAG激光器实现了工作状态的最优化.针对左右腔镜,对设计方案所得结果进行了比较,并分析了方案的可行性.结果表明,左腔镜两设计方案与优化目标的差距均小于1.15%;右腔镜两设计方案与优化目标的差距均小于0.05%,本文设计为激光器谐振腔的腔镜镀膜提供了可靠的理论依据.【期刊名称】《大学物理实验》【年(卷),期】2017(030)005【总页数】4页(P96-99)【关键词】Filmstar软件;光学薄膜;谐振腔镀膜;Yb:YAG激光器【作者】廖川;贾红宝;郑孟莹;许崇【作者单位】辽宁科技大学,辽宁鞍山 114051;辽宁科技大学,辽宁鞍山 114051;辽宁科技大学,辽宁鞍山 114051;辽宁科技大学,辽宁鞍山 114051【正文语种】中文【中图分类】O4-33*通讯联系人随着InGaAs激光二极管(发射波长为0.9~1.1 μm)的发展,新型掺Yb3+激光材料引起了人们浓厚的兴趣[1]。

由于Yb:YAG晶体具有晶场分裂能大、光学和热力学性能优异、材料热负荷低、可进行较高浓度掺杂等特点,已经成为最具有发展潜力的掺 Yb3+激光材料之一[2]。

Yb:YAG激光器自20世纪70年代早期开始研究[3],1994年,德国Giesen等[4]设计出具有多通泵浦结构的Yb:YAG薄片激光器;2000年,C.Stewen等[5]利用Yb:YAG 薄片激光器获得了647 W 的连续激光输出;此后,张丽哲等[6]人研制出在1 030.5~1 055.5 nm范围内连续可调谐的Yb:YAG激光器。

众所周知,激光的输出与谐振腔镜的透射率有紧密联系。

要在该范围内对激光输出进行调节,就必须保证谐振腔在该波段的出射率保持不变,这就对在谐振腔上所镀的用于增透或增反的光学薄膜提出了要求。

纳米晶软磁合金磁损耗的研究

纳米晶软磁合金磁损耗的研究

天津理工大学本科毕业论文纳米晶Fe-Co-Cu-Nb-Si-B软磁合金磁损耗的研究学生姓名:赵承鹏学号:20072130所学专业:集成电路设计及集成系统班级:07级1班指导老师:丁燕红2011年6月摘要与传统的金属软磁材料相比,纳米晶软磁合金由纳米尺度的晶粒以及大量非晶界面组成,晶粒间的铁磁交换耦合作用使其内部的局域磁各向异性显著降低,因而具有优异的软磁性能,有望获得高磁导率。

交变磁场作用下,磁性材料一方面被磁化,另一方面会产生能量损耗,导致铁芯发热。

所谓磁损耗是指磁性材料在交变磁场作用下产生的各种能量损耗的统称。

磁性材料在交变磁场中,磁感应强度B在时间上滞后于交变磁场H的变化,使B和ACH之间存在一个相位差 。

通过测量合金交变磁场下的磁导率,来计算磁滞损耗AC系数a,涡流损耗系数和剩余损耗系数c,对于研究软磁合金的交流磁性有重要意义。

关键词:Fe-Co-Cu-Nb-Si-B合金、纳米晶、软磁材料、初始磁导率、磁损耗。

ABSTRACTCompared to soft magnetic materials of conventional metallic, nanocrystallinesoft magnetic alloy grains by the nano-scale and a large of amorphous interface, theferrom-agnetic exchange of grains coupling between the local magnetic anisotropy within itssignificantly lower. And thus excellent soft magnetic properties is expected to behigh permeability.In the Alternating magnetic field, the magnetic material is magnetized on the one hand, on the other hand the lossing energy will lead to core heat. The so-called Magnetic loss is the general designation of the magnetic materials in alternating magnetic field that produced by a variety of energy loss collectively. In alternating magnetic field,B(magnetic induction) is lagged to the change ofH(the alternating magneticACfield)on the time, There is a phase difference between the B andH. By measuringACthe magnetic permeability in the alloy alternating magnetic field, to calculate the coefficient of hysteresis loss, eddy current loss coefficient and residual loss coefficient is very important for researching the exchange magnetic of the soft magnetic alloy.Key Words: The alloy of Fe-Co-Cu-Nb-Si-B , Nanocrystal , Soft Magnetic Material , Initial permeability , Magnetic loss .目录第一章绪论 (1)引言 (1)1.1纳米晶软磁材料应用 (2)1.2 纳米晶软磁合金的制备 (3)1.3 Fe-Cu-Nb-Si-B纳米晶合金的形成机理 (4)1.4 纳米晶软磁材料的化学成分、元素作用及性能 (6)1.4.1 化学成分 (6)1.4.2 元素作用 (6)1.4.3 纳米晶软磁合金的磁特性 (7)1.5 纳米晶软磁材料的研究现状 (7)1.6 本文的主要研究内容 (8)1.6.1 本课题的研究内容 (8)1.6.2 课题的意义 (8)第二章实验方法 (9)2.1 样品的制备 (9)2.1.1 非晶晶化法制备纳米晶合金 (9)2.1.2 甩带与制环 (9)2.1.3 纳米晶合金样品的制备 (10)2.2 测量装置与测量方法 (10)2.2.1 HP4194A阻抗分析仪 (10)2.2.2 实验技术 (14)第三章Fe-Co-Cu-Nb-Si-B磁损耗的研究 (15)3.1 合金软磁材料的损耗分类 (15)3.1.1 涡流损耗 (15)3.1.2 磁滞损耗 (15)3.1.3 剩余损耗 (15)3.2 约旦分离法确定纳米晶合金的损耗系数 (15)3.2.1 (Fe0.5Co0.5)73.5Cu1Nb3Si13.5B9材料的损耗系数 (16)3.2.2 (Fe0.5Co0.5)76.9Cu0.6Nb2.5Si11B9材料的损耗系数 (20)第四章结论 (25)参考文献资料 (26)第一章 绪论引言一直以来,材料都是人类社会发展的物质先导和基础,新材料的发明和运用则是人类社会进步的里程碑。

FEMAG晶体生长计算软件之Czochralski (CZ) Process (FEMAG-CZ)

FEMAG晶体生长计算软件之Czochralski (CZ) Process (FEMAG-CZ)

FEMAG晶体生长计算软件Czochralski (CZ) Process(FEMAG-CZ)FEMAG直拉法模拟软件(FEMAG-CZ)用于模拟直拉法工艺(包括Cz, MCz, VCz,泡生法)。

FEMAG-CZ直拉法模拟软件用于新的热场设计,并研发新的方法以满足新的商业需求点,比如:✓大直径晶锭生长✓无缺陷硅晶锭生长✓提高成品率✓氧含量控制✓降低碳含量✓晶锭半径和沿轴向的电阻率差异减小✓CCZ工艺仿真✓磁场设计✓蓝宝石生长工艺设计FEMAG-CZ模拟软件通过降低试验成本而节省了R&D消耗。

大直径晶锭生长以期不进行大量昂贵的可行性试验生长大尺寸晶体看起来是不太现实的。

FEMAG-CZ软件提供这种可能性。

为了生产450 mm及以上的大尺寸无缺陷硅晶体,晶体生长工程师通过使用FEMAG-CZ来定义关键的工艺参数,而无需任何材料和能源的消耗。

FEMAG-CZ能够设计新的热场并研发新的工艺技术,在FEMAG直拉法模拟软件的帮助下,晶体生长工程师能够在一个有效的虚拟环境中优化每一个关键参数,比如旋转速率,提拉速度,气体流速,压强和功率消耗等。

FEMAG直拉法模拟还能进一步为晶体生长工程师给出在某一工艺配置下产出的最终成品的质量和成本信息,比如晶体中的温度梯度,氧/碳/掺杂物/微缺陷分布等。

通过软件能够获得硅/锗/蓝宝石晶体质量和产品成本信息,这一模拟过程无需任何材料和能量的消耗。

FEMAG 3D 熔体流动模拟结果FEMAG动态模拟无缺陷硅晶锭生长无缺陷晶体硅生长是世界上最大的难点之一。

FEMAG模拟软件能够帮助工程师运用自己创新的技术生长出无缺陷晶体。

运用FEMAG软件缺陷工程模块可以预测晶体炉或者其他指定直拉法工艺环境中生长的晶体成品质量。

缺陷工程模块能够洞悉硅、锗生长过程中填隙原子,空位和微孔演变过程。

FEMAG-CZ能够成为你的测试平台,试验在不同的操作条件下对晶体生长质量的影响,如✓热场设计✓加热器功率✓晶体和坩埚的旋转速率✓晶体提拉速度,坩埚的位置✓气体流率和压强一旦掌握了晶体生长工艺中的动态规律,就可以找到最优的配置以增加成品率和投资回报。

在稳态-动态磁场作用下大尺寸晶体直拉生长的硅熔体流动模拟(使用FEMAG)

在稳态-动态磁场作用下大尺寸晶体直拉生长的硅熔体流动模拟(使用FEMAG)

Journal of Crystal Growth 230(2001)92–99Numerical investigation of silicon melt flow in large diameter CZ-crystal growth under the influence of steady and dynamicmagnetic fieldsJ.Virbulis a,*,Th.Wetzel b ,A.Muiznieks b ,B.Hanna a ,E.Dornberger a ,E.Tomzig a ,A.M .uhlbauer b,W.v.Ammon aaWacker Siltronic AG,P.O.Box 1140,Bur g hausen,D-84479,GermanybInstitute for Electroheat,Uni v ersity of Hanno v er,Wilhelm-Busch str.4,D-30167,Hanno v er,GermanyAbstractTurbulent silicon melt flows are studied in large diameter Czochralski crucibles under the influence of alternating,steady and combined magnetic fields.The investigations are based on the experimentally verified two-dimensional axisymmetric mathematical models.The influence of steady,alternating and combined magnetic fields on the flow pattern andtemperature fieldis investigated .Global heat transfer andmelt flow calculations are coupledandthe influence of melt convection on the interface shape is studied and compared with experimental data.r 2001Elsevier Science B.V.All rights reserved.PACS:47.27.i;47.65.+aKeywords:A1.Fluid flows;A1.Heat transfer;A1.Magnetic fields;A2.Czochralski method;B2.Semiconducting silicon1.IntroductionThe conversion to large silicon (Si)single crystals of 300mm diameter requires larger batch sizes which generates turbulent melt convection with Grashof number up to 1010.Several efforts of the crystal growth industry are dedicated to control the melt flow andthe temperature d is-tribution.In particular,its focus is on the correct prediction of the interface shape and the related point defect distribution in the crystal,of theoxygen transport,of the crucible overheating and of the dislocation free growth.Besides conven-tional means,electromagnetic steady (DC)and dynamic (AC)fields offer new possibilities to meet the continuously increasing demands for crystal quality,yieldimprovement andcost red uction.A numerical simulation helps in investigating the wide field of possible process conditions,to save a lot of experimental costs for large diameter crystal growth andto red uce the time to market.Global heat transfer calculations [1–3]provide goodagreement with temperature measurements in crystal andinsulation [4]andare established as a standard tool for Czochralski (CZ)process development.*Corresponding author.Tel.:+49-8677-83-4227;fax:+49-8677-83-7303.E-mail address:janis.virbulis@ (J.Virbulis).0022-0248/01/$-see front matter r 2001Elsevier Science B.V.All rights reserved.PII:S 0022-0248(01)01321-5Several numerical studies are devoted to melt flow in CZ crucibles.Most of them are made for two-dimensional(2D)idealized cylindrical geo-metries with small crucible sizes andsimplified thermal boundary conditions.In some cases,mag-neticfield effects have been studied,e.g.Ref.[5]. Three-dimensional(3D)direct numerical simula-tion without turbulence models in idealized geometries show non-symmetric time-dependent flow behavior in small[6]andmed ium(Gr=108) [7]crucibles.Global heat transfer coupledwith2D meltflow in large crucibles using low Re number k2e models has been calculated[8,9]without the influence of magneticfields.In this work,we present numerical investiga-tions of turbulent Si-meltflows in large diameter CZ crucibles under the influence of AC,DC and combinedmagneticfield s.The investigations are basedon the experimentally verifiedsystem of axi-symmetric mathematical models,where the melt convection andthe global heat transfer are coupled.The influence of melt convection on the interface shape andthermal grad ients in the crystal is studied and compared with experimental data.2.Numerical models2.1.Couplin g between g lobal heat transferand melt con v ectionGlobal heat transfer in the whole axisymmetric CZ system is simulatedwith thefinite element cod e FEMAG,described in detail in Ref.[1].Heat transfer is modeled using quadraticfinite elements in solids and view factors in radiative enclosures. Conductive and convective heat transfer in argon gas is neglectedd ue to low pressure.The growth process is assumedto be quasi-stead y andthe release of latent heat of fusion at the crystallization interface is proportional to the imposedgrowth rate.Melt convection is simulatedwith the program package CFD-ACE[10].The coupling between global thermal andmeltflow simulations is realizedby exchange of the heatfluxes at the melt boundaries.Thefirst global simulation is carried out using the effective heat transport coefficient including conductive and convective contributions in the melt with FEMAG.The calculatedheat fluxes along the crucible andmelt free surfaces are provided as thermal boundary conditions for the meltflow simulation with CFD-ACE.Meltflow simulation uses the interface shape calculatedin FEMAG.Afterflow simulation,temperatures on melt boundaries in CFD-ACE and FEMAG are comparedandnormal heatfluxes on the melt boundaries in FEMAG are adjusted to reach the same crucible temperature as in CFD-ACE.At the crystallization interface,heatfluxes from the melt side in FEMAG are adjusted to get the same distribution as in CFD-ACE.New heatfluxes along the crucible andmelt free surfaces as well as the new interface shape are usedto perform the next meltflow calculation.Typically,five global iterations are requiredto reach thefinal solution.2.2.Melt con v ection and direct dampin g of turbulence by a steady ma g neticfield Buoyancy,centrifugal,Marangoni andLorentz forces are considered in the convection simulation with CFD-ACE.A low Reynolds number k2e turbulence model is used because of thin velocity andconcentration bound ary layers.AC andDC magneticfieldd istributions are calculatedwith FZH DEM[11,12].Since the effects of AC andDC fields are different,two separate Lorentz force terms are used.For ACfields,mean values of the Lorentz force are applied,which are calculated as a product of alternatingfield and in the melt induced alternating current.These forces depend on the frequency andon the geometry of coils and melt,but are independent of the melt velocities. For DCfields,Lorentz force depends onflow velocity.An additional equation of electric poten-tial,which is causedby interaction of stead y magneticfieldandmelt velocity,is solvedtogether with the hydrodynamic equations.2D axisym-metric models are used for the meltflow in a CZ crucible because they are a goodcompromise between efficiency andaccuracy for engineering applications.In the standard model,the Lorentz force suppresses the mean velocity and,hence less turbu-lence is generated.However,velocityfluctuations,J.Virbulis et al./Journal of Crystal Growth230(2001)92–9993modeled by using the turbulent energy k,are also directly damped by the magneticfield.Hence,we improve the standard model by adding a negative term for turbulence generation P B to the k equationP B¼@c B k kB2;ð1Þwhere k is the electric conductivity.A useful and precise value of c B has to be verifiedby compar-isons to experimental data.Models of melt convection and direct damping of turbulence by steady magneticfield are reported in more detail in Ref.[12].2.3.Additional turbulence g enerationOne of the main problems for2D axisymmetric model is the artificial symmetry line.Below the interface,the meltflow is directed towards the center and then downwards along the symmetry axis.Due to the mass conservation,theflow velocity near the axis is very high(see Fig.3a).In reality,the symmetry line does not exist and the position of this jet canfluctuate about the axis, generating an additional mixing in this region.In metallurgical applications,similar problems are solved by applying additional turbulence in the areas between two vortices where high velocities exist[13].We propose a model taking into account this effect in CZ crucibles.An additional low frequency turbulence will be generatedin the presence of a jet on the center line according to the equationk LF¼u2=2:ð2ÞTherefore,we introduce an additional generation term in k equation proportional to the square of velocity in a meridional planeP LF¼c LFðu2r þu2zÞ:ð3ÞThe value of the constant c LF can be foundby comparing the simulation results with experimen-tal d ata.This mod el is verifiedandd iscussedin Section3.2.3.Verification of models3.1.Couplin g between g lobal heat transferand melt con v ectionFig.1shows the temperatures andFig.2shows the heatfluxes along the crucible,free melt surface andcrystallization interface,calculatedfor a 200mm crystal anda24in crucible.The numerical gridwith12,300cells,usedfor the meltflow calcu-lations,isfine enough to reach a grid independent solution.The solidlines are the results without convection,the boldsolidlines are thefinal results with convection andthe d ashedlines result after thefirst convection calculation.The crucible temperature is strongly influencedbyconvection, Fig.2.Heatflux at the crucible wall,the free melt surface and the crystallization interface.Solidline:without convection,bold solidline:with convection,d ashedline:afterfirst couplingiteration.Fig.1.Temperature at the crucible wall andthe free melt surface.Solidline:without convection,boldsolidline:with convection,dashed line:afterfirst coupling iteration.J.Virbulis et al./Journal of Crystal Growth230(2001)92–99 94of course,depending on the initial guess of the effective heat transfer coefficient in the melt,used in global simulation.But the heat flux in our test case is changedat the crucible by less than 10%andat the free surface by less than 5%,if convection is considered.The heat flux at the crystallization interface is changedsignificantly,and,as a result,the interface shape is changed,as well.After the first convection simulation,the tem-perature differs from the final solution by less than 2K.The flow pattern is also nearly unchanged.Therefore,for calculations of crucible temperature andflow pattern only one iteration without back-coupling to global heat distribution is normally sufficient.A full coupling is necessary for the detailed calculation of the interface shape and temperature gradients at the interface.3.2.Turbulence models3.2.1.Additional low frequency turbulence g enerationConvection in a CZ system with a 14in crucible anda 4in crystal is shown as streamlines and vectors in a meridional plane calculated with the standard model in Fig.3a.High velocities exist in the central jet under the crystal.Although Si has a low Prandtl number of 0.013,the strong jet influences the temperature fieldandthe isolines show a distinct deflection downward near the axis (solidlines,Fig.4a).Using the mod el with an additional low frequency turbulence generation,the velocity of this jet is lower (Fig.3b)andthe isolines have less deflection (solid lines,Fig.4b).The dashed lines in Fig.4are the experimentaltemperature distribution,measured by Gr .abner et al.[14].As heat fluxes are usedfor thermal boundary conditions at the crucible wall,tempera-ture differences between the interface and crucible wall (at the reference point usedin Fig.5)are a measure for the right description of the turbu-lent heat transfer in the melt.The temperature difference with the standard model is 43K,in the calculation with an additional low frequency turbulence generation @39K,andin the experi-ment @35K.As the temperature difference in the calculation without convection is 75K,we con-clude that the mean turbulent characteristics of the heat transfer are described accurately using 2D k 2e models.As the temperatures on melt bound-aries in the experiment andsimulation are not explicitly the same,an additional source of error occurs,because the buoyancy convection strongly depends on the temperature boundary conditions.However,our experience indicates that for the mentionedd ifferences the flow structure andthe turbulence level changes are not remarkable.Especially in the central melt region,the advanced model shows a better agreement with the experimental data than the standard model.However,the temperature in the region under the crystal andnear the free surface in both mod els is higher than in the experiment.Obviously,real 3D turbulent convection cannot be described with all details by 2Dmodels.Fig.3.Melt flow in a 14in crucible.Streamlines every 1g/s and velocity in meridional plane with maximum value of 3.3cm/s.Standard model (a)and advanced model with low frequency turbulence generation (b).J.Virbulis et al./Journal of Crystal Growth 230(2001)92–99953.2.2.Direct dampin g of turbulence by a steady ma g netic fieldFig.5shows a comparison of temperatures at a reference point which is locatedat 60mm d epth on the wall of a 14in crucible for the uniform axial fieldwith 64and128mT as well as for a CUSP fieldwith 40mT on the crucible wall (CUSP fieldis a non-uniform axisymmetric fieldwith a low field intensity near the crystal anda high fieldintensity near the crucible wall,generatedusing a system of two coils with opposite directions of current).The dashed lines are the temperatures calculated using the model without direct damping of turbulence,but with turbulence generation by the fluctuating jet under the crystal (see Eq.(3)).The solid lines are the temperatures calculatedusing the ad vanced model with the direct damping of turbulence.Thelarge circles are the experimental temperature measurements from Ref.[14].The experimentally observedincrease of the crucible temperature for CUSP fields is well reproduced by using the direct damping model.The agreement between the simulation andexperiment for the axial fieldis also better using the model with direct damping.The differences for the axial field of 64mT are probably causedby 3D effects.The models for alternating fields are verified with experimental data in Ref.[12].4.Results and discussion4.1.Influence of ma g netic fields on the melt con v ectionThe use of magnetic fields is particularly promising for large diameter crystal growth where the control of highly turbulent convection with conventional means is difficult.Therefore,after verification of models in a 14in system,we have studied the melt convection for 300mm crystal growth from a 28in crucible filledwith 150kg of molten silicon.The rotation rates of the crystal andthe crucible are @12and4rpm,respectively.Thermal boundary conditions are taken from the global simulation as described in the previous section.Only one iteration without back-coupling is performed using the advanced model with damping and generation of turbulence.Fig.6shows the isotherms (right)as well astheFig.4.Temperature distribution in a 14in crucible.Standard model (a F solidlines)andmod el with low frequency turbulence generation (b F solid lines).Experimental data (dashed lines)from Ref.[14].3.6K betweenisolines.Fig.5.Temperature in reference point at the 14in crucible wall at a depth of parison of standard model,model with direct turbulence damping and experiment.J.Virbulis et al./Journal of Crystal Growth 230(2001)92–9996streamlines andvectors of merid ional velocity (left).The following cases are simulated:without magnetic field(a),stead y axial fieldof 50mT (b),CUSP fieldof 50mT (c),AC field(d )anda combinedAC fieldandaxial fieldof 50mT (e).The value of the CUSP fieldis d efinedas the maximum fieldstrength in the melt.The alternat-ing fieldis generatedfrom a coil with 4000A windings and with a frequency of 50Hz.Without a magnetic field,a strong vortex exists under the crystal,and,a weak vortex in the outer part of the crucible.The intensity of the outer vortex becomes lower with increasing crucible rotation due to conservation of angular momen-tum.Applying an axial field,the flow velocity is suppressedanda one-vortex structure is formed .The CUSP fieldsuppresses the velocity in the outer part of the crucible where the fieldhas high intensity.The convection is weakly changednear the crystallization interface,where the fieldin-tensity is low.Suppressedvelocities red uce the probability for particles,generatedby cruciblecorrosion,to reach the crystal,reducing the prob-ability of dislocation generation.The alternating fieldcreates a structure of two strong vortices at the outer part of the crucible causedby Lorentz force directed at the center of the crucible.The combination of DC andAC field s show character-istics of both fields,damped convection in the region near the crystal andstrong two-vortex-structure in the outer part of the melt.With both,alternating and combined fields,the flow direction on the free surface is reversedandparticles cannot reach the crystal along the free surface.CUSP andaxial field s increase the maximum temperature difference between solidification tem-perature andcrucible hot spot by 35%and12%,respectively.The AC fieldd ecreases the tempera-ture difference by 30%and the combined field by 6%.The variation of the crucible temperature is causedmainly by the change of turbulent heatconductivity.Fig.6.Streamlines every 1g/s andvelocity vectors with a maximal value of 2.3cm/s in merid ional plane (left)andisotherms every 2K (right)in a 28in crucible.(a)without magnetic field,(b)with CUSP field of 50mT,(c)with axial field of 50mT,(d)with dynamic field,(e)with a combinedd ynamic fieldandaxial fieldof 50mT.J.Virbulis et al./Journal of Crystal Growth 230(2001)92–99974.2.Effect of the melt con v ection on the defect distribution in crystalGrown-in defects in substrate wafers consider-ably affect the yieldin d evice manufacturing [15].The type of defects are determined by the v =G criteria as described in Refs.[16,17]where v is the growth rate and G is the temperature gradient in the crystal at the growth interface.The gradient G is determined by the thermal ambient of the crystal (hot zone design,heat shield),the growth rate,and the heat flux from the melt on the interface depending on the convection.Therefore,melt convection is an important factor for the grown-in defect formation in silicon crystals.Interface shapes can be measuredas oxygen striations or lifetime maps on the axial cuts of crystals.The interfaces of fast grown (positive deflection)and slow grown (negative deflection)crystals are shown in Fig.7.The large dots represent the measuredvalues,the thin d ashed lines are the calculatedvalues without convection,the boldd ashedlines are the calculatedvalues with our stand ardmod el andthe boldsolidlines are the calculated values with the advanced model.Simulations with convection reproduce the experi-mental shape better than those without convec-tion.The advanced model results agree better forhigh growth rate.For low growth rate,both models similarly deviate from the experiment.In Fig.7on the right,the effect of axial and CUSP magnetic fields on the interface shape is shown.Each magnetic fieldhas a certain impact on the interface shape.However,these differences are smaller than differences between the models,therefore,for quantitative results further adjust-ment of models based on experimental data is necessary.5.ConclusionsCoupling of global heat transfer andmelt convection is necessary for the precise prediction of interface shape.Only one iteration is sufficient for the prediction of crucible temperature and flow pattern if heat flux boundary conditions from global heat transfer calculation are applied.Proposedmod ifications of 2D turbulence mod els show better agreement with experiments as stan-dard models.Steady and alternating magnetic fields can be used to radically change the flow pattern,the heat andmass transfer in the melt as well as the shape of crystallization interface.AcknowledgementsThe authors wouldlike to thank O.Gr .abner andG.M .uller for the experimental results whichwere obtainedin cooperation with Wacker Sil-tronic AG andFraunhofer-Institute,IIS-B,De-partment of Crystal Growth,Erlangen,Germany.The work was supportedby the German Fed eral Ministry of Education,Science,Research and Technology under Contract Nr.01M2973A.The authors alone are responsible for the content of this publication.References[1]F.Dupret,P.Nicodeme,Y.Ryckmans,P.Wouters,M.J.Crochet,Int.J.Heat Mass Transfer 33(1990)1849.[2]R.A.Brown,T.A.Kinney,P.Sackinger, D.Bornside,J.Crystal Growth 97(1989)99.Fig.7.Left:Interface shapes for fast grown (positive deflec-tion)andslow grown (negative d eflection)crystals:without convection (thin dashed lines),with standard turbulence model (bold dashed lines),with advanced model (bold solid lines),experiments (large dots).Right:effect of the magnetic field:without magnetic field,with steady axial magnetic field of 50mT andwith CUSP fieldwith 50mT.J.Virbulis et al./Journal of Crystal Growth 230(2001)92–9998[3]J.Baumgartl,A.Bune,K.Koai,G.M.uller, A.Seidl,Mater.Sci.Eng A173(1993)9.[4]E.Dornberger, E.Tomzig,H.-J.Leister,Ch.Schmitt,S.Schmitt,A.Seidl,G.M.uller,J.Crystal Growth180(1997)461.[5]P.Sabhapathy,M.E.Salcudean,J.Crystal Growth113(1991)164.[6]K.Yi,V.B.Booker,M.Eguchi,T.Shyo,K.Kakimoto,J.Crystal Growth156(1995)383.[7]C.Wagner,R.Fridrich,in:H.Koener,R.Hilbig(Eds.),New Results in Numerical andExperimental Fluid Mechanics,Vieweg Verlag,Deutschland,pp.367.[8]A.Lipchin,R.A.Brown,J.Crystal Growth216(2000)192.[9]Yu.E.Egorov,Yu.N.Makarov, E.A.Rudinsky, E.M.Smirnov,A.I.Zhmakin,Mat.Res.Soc.Symp.Proc.490(1998)181.[10]CFD-ACE Theory Manual,Version4.0CFD ResearchCoorporation,Huntsville,1998.[11]A.M.uhlbauer,A.Muiznieks,G.Raming,H.Riemann,A.L.udge,J.Crystal Growth198/199(1999)107.[12]Th.Wetzel, A.Muiznieks, A.M.uhlbauer,Y.Gelfgat,L.Gorbunov,J.Virbulis, E.Tomzig,W.v.Ammon, Proceedings of the Workshop of Modeling in Crystal Growth,Stony Brook,NY,2000.[13]A.M.uhlbauer,E.Baake,A.Jakovics,Energy Transfer inMHD Flows,Aussois,1994,Vol.1,pp.349.[14]O.Gr.a bner,G.M.uller,J.Virbulis, E.Tomzig,W.v.Ammon,E-MRS Spring Meeting2000,Strasbourg, France,Mater.Sci.Eng.B.,to be published.[15]J.Park,G.Rozgonyi.SolidState Phenom.47–48(1996)327.[16]V.V.Voronkov,J.Crystal Growth59(1982)625.[17]E.Dornberger,W.von Ammon,J.Electrochem.Soc.143(1996)1648.J.Virbulis et al./Journal of Crystal Growth230(2001)92–9999。

大尺寸氟化铽锂晶体生长与性能

大尺寸氟化铽锂晶体生长与性能
crystals.
Key words: LTF; magneto-optic crystal; Czochralski method; absorption coefficient; optical property; magneto-optical
property; large size
0 引 言
达 2 英寸(1 英寸 = 2. 54 cm) 的大尺寸 LTF 晶体,并测试了晶体的光学质量,研究了晶体的光学性能和磁光
性能,为后续进一步开展应用研究奠定了基础。
1 实 验
1. 1 晶体生长
利用炉膛内径为 800 mm 的 DJL-800A 型上称重自动控径单晶炉,采用电阻加热提拉法生长 LTF 晶体。
试晶体的透过光谱。 在晶坯转肩部位取样品研磨后,利用粉末 XRD 测试生长晶体物相组成。 利用金刚石单
线切割机和无形外圆磨床,从晶坯上加工出 2 支直径 10 mm、长度 35 mm 的 c 切 LTF 晶体棒,端面经光学精
密抛光后,参照国标 GB / T 11297. 1—2017《 激光棒波前畸变的测量方法》 、GB / T 27661—2011 《 激光棒单程
absorption coefficient were tested. The results show that the grown LTF crystal have good optical quality. The magneto-optical
performance of the LTF crystal was also researched, and its Verdet constant is about 98% of that of the commercial TGG
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The following observations can be drawn:

With a cusp magnetic field, the overall oxygen concentration in the crystal decreases strongly This is effect is more important in the lower part of the crystal, where oxygen concentration decreases from about 7 ppma down to 3 ppma This lower oxygen concentration is the result of: - a stronger turbulent effect (see the turbulent viscosity) below the solid/liquid interface when no cusp field is applied, thereby facilitating the transfer of oxygen from the bottom of the melt towards the solid-liquid interface;
8 7 6
Ox yge n ppm a
C U SP 4 3 2 1 N o C U SP
Ox yge n ppm a
5
5 C U SP 4 3 2 1 N o C U SP
0 0 0,02 0,04
Ra dius (m )
0,06
0,08
0,1
0,12
0 0 0,02 0,04
Ra dius (m )
FEMAGSoft © 2012
NOCUSP vs CUSP - OXYGEN IN THE MELT
Animation - Oxygen-melt.avi
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NOCUSP vs CUSP - PHOSPHOROUS IN THE MELT - 1 Animation - Phosphorous-melt.avi
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EFFECT OF A CUSP MAGNETIC FIELD ON OXYGEN CONCENTRATION (continued) This lower oxygen concentration is the result of:

- a different flow configuration: with a cusp field, the main flow cell (also called a Proudman cell) extends farther from the axis and closer to the melt/gas interface, thereby facilitating oxygen evaporation, while without cusp field the main flow cell remains located under the solid/liquid interface The radial segregation is not very different with or without cusp field, showing a strong radial segregation close to the crystal wall in both cases, as a consequence of the strong oxygen evaporation prevailing near the tri-junction
PARAMETERS OBSERVED

The concentration of the Oxygen in the melt and crystal. The concentration of the Phosphorous in the melt and crystal The concentration of the Carbon in the melt and crystal
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PREDICTION OF RADIAL PHOSPHORUS RESISTIVITY
FEMAGSoft © 2011
NOCUSP VS CUSP TEMPERATURE IN THE MELT
Animation – Temperature-melt.avi
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0,06
0,08
0,1
0,12
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PREDICTION OF OXYGEN CONCENTRATION ALONG THE HEIGHT OF THE CRYSTAL
Oxygen Concent rat ion variat ion in a Vert ical Direct ion
PREDICTION OF RADIAL OXYGEN CONCENTRATION
Radial Oxygen Concent rat io H=840mm
8 7 6
8 7 6
Radial Oxygen Concent rat ion H=1070
Ox yge n ppm a
5 4 3 2 1 0 0 0,02 0,04 0,06(m ) Ra dius 0,08 0,1 C U SP N o C U SP
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CRYSTAL RADIAL OXYGEN MEASUREMENT POSITIONS
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PREDICTION OF RADIAL OXYGEN CONCENTRATION
Radial Oxygen Concent rat ion H=90mm
FEMAG晶体生长仿真软件 --CUSP磁场对8英寸半导体级硅晶体生长的影响
Report Bruker ASC Feb, 9, 2012
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OBJECTIVE
PROCESS CONDITION

CZ furnace - 8 inch diameter crystal, 24 inch diameter crucible To predict the impact of a CUSP magnetic field operating at 140 A with the neutral-axis located on the melt free surface.
NOCUSP vs CUSP – MELT TURBULENT VISCOSITY
Animation – Turbulent viscosity-melt.avi
FEMAGSoft © 2012
NOCUSP vs CUSP-VELOCITY & STREAM FUNCTION Animation - psi-melt.avi


- a lower temperature along the bottom of the crucible when a cusp field is applied, leading to a lower oxygen dissolution in this region; this lower temperature appears as an additional consequence of the lower melt flow turbulence, the associated heat transfer reduction, and the resulting cooling down of the bottom of the melt
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PREDICTION OF PHOSPHOROUS CONCENTRATION ALONG THE HEIGHT OF THE CRYSTAL
FEMAGSoft © 2012
CUSP vs NOCUSP - OXYGEN CONCENTRATION - CRYSTAL
9 8 7 6
Ox yge n ppm a
5 4 3 2 1 0 0,2 0,4 0,6 0,8 1 1,2 1,4 1,6 C U SP N o C U SP
H e ight (m )
FEMAGSoft © 2012
PREDICTION OF CARBON CONCENTRATION ALONG THE HEIGHT OF THE CRYSTAL
8
9
Radial Oxygen Concent rat ion H=250 mm
8 7
7 6
Ox yge n ppm a
Ox yge n ppm a
5 4 3 2 1 0 0 0,02 0,04 0,06 0,08 0,1 C U SP N o C U SP
6 5 4 3 2 1 0 C U SP N o C U SP
Ox yge n ppm a
5 4 3 2 1 0 C U SP N o C U SP
0,12
0
0,02
0,04
Ra dius (m )
0,06
0,08
0,1
0,12
Radial Oxygen Concent rat ion H=1290mm
8 7 6
Radial Oxygen Concent rat ion H=1550 mm
FEMAGSoft © 2012
Animation – Phos-resistivity.avi
NOCUSP vs CUSP - CARBON IN THE CRYSTAL
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