Ultrathin epitaxial ferroelectric films grown on compressive substrates Competition between
旋涂法制备功能薄膜的研究进展
旋涂法制备功能薄膜的研究进展摘要：作为众多的薄膜制备方法之一，旋涂法具备薄膜厚度精确可控、高性价比、节能、低污染等优势，在微电子技术、纳米光子学、生物学、医学等领域中有着广阔的应用前景．功能薄膜是发展信息技术、生物技术、能源技术等领域和国防建设的重要表面材料和器件，关系到资源、环境及社会的可持续发展．旋涂法制备的薄膜厚度在30nm一2000nm之间精确可控，其设备结构简单且易于操作，具备优良的性价比．现已广泛应用于微电子行业的光刻图案化( Lithographic patterning process ) 、印刷电路(Printed circuit)和集成电路(Integrated circuit)的制造，以及光储存媒体介质( DVD- R、CD- R等)的感光胶( P h o t o r e s i s t ) 、染料( Dye ) 、粘合剂、物理保护层等聚合物薄膜的涂覆．旋涂法在其它许多新型领域也有一定的应用，如薄膜晶体管、光子晶体材料、光波导、有机发光二极管薄膜、光电转换薄膜电极以及生物/化学功能膜等薄膜类器件的制备．旋涂法涉及到许多物理化学过程，如流体流动、润湿、挥发、粘滞、分散、浓缩等．在研究这些过程时，流体力学传质、传热、传动的原理是非常重要的．其中，需要考虑的参数主要包括薄膜的结构、厚度、面积等性能参数以及转速、粘度、挥发速率等操作参数[1,2,3]。
1、旋涂法原理旋涂法因其所用流体粘度较大，呈胶体状，所以也被称为匀胶。
一个典型的旋涂过程主要分为滴胶、高速旋转和干燥( 溶剂挥发) 三个步骤．首先，滴胶是将旋涂液滴注到基片表面上，然后经高速旋转将其铺展到基片上形成均匀薄膜，再通过干燥除去剩余的溶剂，最后得到性能稳定的薄膜．对于各种粘度、润湿性不同的旋涂液，通常使用的滴胶方法有两种，即静态滴胶和动态滴胶[4]，旋涂法中的高速旋转和干燥是控制薄膜厚度、结构等性能的关键步骤，因此这两个阶段中工艺参数的影响成为研究的重点[5]．2、旋涂法的研究现状旋涂法符合光学、微电子学、纳米光子学、纳米电子学等许多薄膜类器件的制备要求．在微电子行业中，运用旋涂法制备出的新型薄膜晶体管器件具有电子迁移率高、厚度精确、均匀稳定等优良的性能，该薄膜可广泛用于制造LED显示屏、存储智能卡、微处理器等电子器件．例如IBM托马斯·J·沃森研究中心的科研工作者们[6,7]在高纯硅片上运用旋涂法制备出特定晶型结构、纳米级厚度的场效应电晶体薄膜，并大量减少了传统工艺中有毒溶剂的用量．在与电化学法、物理/化学气相沉积法等薄膜制备技术对比时，旋涂法具备工艺条件温和、操作控制简单等独特优势，所以在降低污染、节能、提高性价比等方面效果显著．近年来，旋涂法不断受到人们重视，其应用逐渐推广到物理学。
硒化铁超导薄片表面微结构的制备与性能表征
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自从 １ １ 年 Ｏ ｎ ｓ 发现单质汞在 ４２ ９１ ｎ ｅ［ ．Ｋ附近 电 阻变 为零 ，也就是超导现象 以来 ，超导物理这个凝 聚态物理 中最活跃的领域经历 了低 谷与高 潮 ，不断
探索着超 导现象的物理本质并 制备新型超 导材料 。
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了 ４种具有代表性的铁基超导体化合物 ：１ ）氟掺杂
的 Ｒ ＦＡ Ｏ ｅ ｓ四元化合物 （ ｌ１相，其中 Ｒ代表 Ｌ 、 １ｌ ａ ｃ 、Ｐ 、 、Ｓ ｅ ｒ Ｎｄ ｍ、Ｇ ｄ等稀土元素 ） ； ）碱金属掺 ２ 杂或 者不掺 杂的双 ＦＡｓ层的 Ａ ｅ ｓ 三元化合物 ｅ Ｆ２ ２ Ａ （２ １２相， 中Ａ为 ｃ ， ｒ ａ 其 ａ ｓ， ， Ｂ 等碱土金属）７］３ ［ ； ）单 － ８ Ｆ Ａ 层的 ＬＦＡｓ ｅｓ ｉｅ 三元化合物 （１ 相）［０ ４ １１ ９ ］ ）四方 －； １
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二硫化钛纳米片材料用于光声成像指导下的肿瘤光热治疗研究
二硫化钛纳米片材料用于光声成像指导下的肿瘤光热治疗研究随着肿瘤的发病率不断增加，对于肿瘤的治疗研究也越来越受到重视。
近年来，光声成像指导下的肿瘤光热治疗逐渐成为一种被广泛研究的方法。
在这种治疗方法中，纳米材料被用作光热转换剂，以实现对肿瘤组织的精确灭活。
二硫化钛纳米片作为一种具有优良光热转换性能的纳米材料，近年来在肿瘤光热治疗研究中受到了广泛的关注。
首先，二硫化钛纳米片具有优异的光热转换性能。
作为一种二维纳米材料，二硫化钛纳米片具有高比表面积和特殊的电子结构，这使得它具有良好的光热转换效率。
在光照作用下，二硫化钛纳米片能够吸收光能并迅速转化为热能，从而产生高温。
这种光热效应可以被应用于肿瘤光热治疗中，将其置于肿瘤组织中，然后通过外部激光的照射来激活其光热效应，从而实现对肿瘤组织的灭活。
其次，二硫化钛纳米片具有较好的可控性。
二硫化钛纳米片的光热效应能够通过调节外部激光的功率、波长和持续时间来实现精确控制。
即使在低功率激光照射下，二硫化钛纳米片也能够产生足够的热量，从而实现对肿瘤组织的灭活。
此外，二硫化钛纳米片还可以通过改变其形态和尺寸来实现对光热效应的调控，以满足不同肿瘤组织的治疗需求。
此外，二硫化钛纳米片还具有良好的生物相容性。
二硫化钛纳米片的合成方法多种多样，可以通过化学气相沉积法、热蒸发法和溶液法等多种方法来制备。
这些方法可以得到具有良好生物相容性的纳米片材料。
此外，二硫化钛纳米片在体内能够被肿瘤组织选择性富集，从而减少了对正常组织的伤害。
这使得二硫化钛纳米片成为一种理想的纳米材料，用于光声成像指导下的肿瘤光热治疗。
综上所述，二硫化钛纳米片作为一种具有优异光热转换性能、可控性和生物相容性的纳米材料，有望在光声成像指导下的肿瘤光热治疗中发挥重要的作用。
未来的研究可以进一步探索二硫化钛纳米片的制备方法和生物效应，以期实现其在临床肿瘤治疗中的应用。
EPITAXIAL GROWTH OF SILICON THIN FILMS FOR SOLAR CELLS
EPITAXIAL GROWTH OF SILICON THIN FILMS FOR SOLAR CELLSG. Andrä1, T. Gimpel1, A. Gawlik1, E. Ose1, A. Bochmann1, S. Christiansen1, G. Sáfrán2, J.L. Lábár2, F. Falk11Institute of Photonic TechnologyAlbert-Einstein-Str. 9, 07745 Jena, Phone (+49) 3641 206438, FAX (+49) 3641 206499, email fritz.falk@ipht-jena.de2Hungarian Academy of Sciences, Research Institute for Technical Physics and Materials Science, 1121 Budapest, Konkoly-Thege M út 29-33, HungaryABSTRACT: Crystalline silicon thin film solar cells on glass substrates are a low cost alternative to silicon wafer cells. As an alternative to a simple furnace annealing step in which a-Si is converted to c-Si with 1 µm grains, an epitaxial crystal growth process is presented here. First a seed layer is prepared on glass by diode laser crystallization of an a-Si layer on glass to result in 100 µm grains. Then a-Si is deposited on top of the seed which is converted to c-Si by epitaxial growth. A 1.1 µm thick c-Si layer with 100 µm grains was produced in this way. The paper presents details of the epitaxial growth process.Keywords: Si-Films, Epitaxy, Multicrystalline Silicon1 INTRODUCTIONCrystalline silicon thin film solar cells on glass substrates are a low cost alternative to silicon wafer cells. The challenge is to produce a suitable crystalline silicon thin film at temperatures up to 650°C endured by the glass. One way is employed by the company CSG (Thalheim, Germany) [1]. In the CSG-process first an amorphous silicon (a-Si) film with the desired doping profile is deposited by PECVD, which within 18 hours is crystallized by solid phase crystallization (SPC) in a furnace. SPC results in crystal grains in the 1 µm range. Solar cell efficiencies of up to 10.4% were achieved.Solar cells on glass with grains two orders of magni-tude larger can be produced by Layered Laser Crystalli-zation (LLC) [2]. In this process one starts from a diode laser crystallized a-Si layer. The resulting multi-crystalline layer has a grain size of about 100 µm and is used as a seed layer for the following epitaxial thickening process. This is done by continuously depositing a-Si on top of the seed and repeatedly applying excimer laser pulses during the deposition. The laser pulses melt the newly deposited a-Si for about 100 ns resulting in epi-taxial growth on top of the crystalline layer below. 2 µm thick cells showed an open circuit voltage of 514 mV and 20 mA/cm² short circuit current even without light trapping [3]. The LLC technology gives layers with the worldwide largest crystalline silicon grains on low temperature substrates. However, the repeated excimer laser irradiation of the growing layer is a technological challenge for industrial production.By combining both the mentioned methods, that is seed layer preparation by diode laser crystallization followed by a furnace annealing step for epitaxial growth of an a-Si layer system 100 µm large crystallites for a multicrystalline silicon thin film cell can be prepared as well. The advantage of this epitaxial solid phase crystalli-zation (ESPC) is the, compared to the LLC process, rather simple technology, which easily may be used in large scale production, and which gives crystal sizes two orders of magnitude larger than the CSG-SPC-process. Moreover, the required crystallization time is even lower.The temperature dependent nucleation and growth rates of c-Si in an a-Si matrix are described in [4,5]. For the ESPC process the problem is the competition of epitaxial growth from the seed with the nucleation of crystallites in the amorphous matrix, which leads to fine grained silicon. Beneficial for the ESPC process is the time lag of nucleation, which is the time needed for the evolution of a distribution of crystalline nuclei in a-Si exceeding the critical size [6]. Only after this time lag the stationary nucleation rate applies. At 600°C the time lag amounts to several hours whereas at 650°C it reduces to about 1 h. This time can be used for undisturbed epitaxial growth starting on the laser crystallized seed layer. A challenge is to find a temperature regime in which epi-taxy is fast enough so that no spontaneous nucleation occurs until epitaxy has converted the whole a-Si layer into c-Si. Another challenge is to get a clean enough interface between the seed and the a-Si so that epitaxy may occur.In the paper the kinetics of the epitaxial growth and the resulting crystal structure are investigated as depend-ing on deposition rate of a-Si. The properties of the films are discussed. The described methods will be used in the European project HIGH-EF to develop multicrystalline silicon thin film solar cells on glass.2 EXPERIMENTALThe preparation of the layer system for the multicrystalline silicon thin film solar cell is done in the following way. In a first step onto a borosilicate glass substrate a several 100 nm thick hydrogen free amor-phous silicon (a-Si) layer is deposited by electron beam evaporation. This layer is crystallized by scanning the 100 µm wide line focus beam of a diode laser (806 nm wavelength, about 10 kW/cm² power density) at a rate of 5 cm/s [2,7]. This procedure, which is performed in am-bient air, generates crystallites in the 100 µm range to act as a seed layer for further processing. After removing an oxide layer on top of the seed layer by etching in 2% HF and an excimer laser cleaning pulse, up to 1.1 µm of a-Si is deposited by electron beam evaporation at rates rang-ing from 10 to 230 nm/min. This layer is in the final step converted to a multicrystalline structure by epitaxial growth from the seed. The EPSC step is performed by a furnace anneal at 600°C or 650°C under Ar gas flow. For comparison, a-Si was crystallized by SPC without seed layer under the same conditions.To monitor the crystallization kinetics the transmission of the Si layers was recorded during the annealing. To this end the beam of a low power HeNelaser (632 nm wavelength) was introduced into the furnace. After passing the sample the transmitted intensity was recorded by a photodiode outside the furnace. The crystal structure after crystallization was investigated by transmission electron microscopy (TEM) and by EBSD (electron back scattering diffraction) which gives orientation maps of the crystal structure in a surface region of the sample.3 RESULTSThe typical fine grained crystal structure following from furnace SPC of a-Si on glass is shown in Fig. 1. The grain size is below 2 µm with no preferential crystal orientation. Fig. 2 shows the evolution of the trans-mission during the annealing which reflects the kinetics of the crystallization. As is obvious a remarkable in-crease of transmission due to an increasing amount of crystalline parts in the film starts after about 6 h. This is a hint for the time lag of nucleation. After the time lag a sigmoidal increase of transmission and of crystal content occurs due to nucleation and growth of crystallites. After about 10 h complete crystallization is achieved.Figure 1: EBSD map of a 1.1 µm thick a-Si layer on glass deposited at 100 nm/min and SPC crystallized by annealing at 600°C for 11 h.Figure 2: Evolution of transmission of an a-Si layer on glass deposited at 100 nm/min during annealing at 600°CIt turned out that the time lag of nucleation as well as the time needed for complete crystallization of a-Si deposited by electron beam evaporation depends on the deposition rate. This is shown in Fig. 3 for an annealing temperature of 600°C. Of particular interest is the time lag which varies between 6 and 25 h.Figure 3: Time lag of nucleation and time for complete crystallization of a 500 nm thick a-Si layer on glass for annealing at 600°C as depending on the deposition rate.In contrast to SPC the epitaxial growth experiments on top of a laser crystallized seed layer lead to much larger crystallites. Fig 4 shows the EBSD map of a seed layer which demonstrates that it consists of grains in the 100 µm range. There is no preferred crystal orientation.Figure 4: EBSD map of a diode laser crystallized seed layer on glass.Fig. 5 shows the evolution of the transmission during EPSC of a 500 nm thick a-Si layer at 600°C which was deposited at a rate of 100 nm/min on top of a laser crystallized seed layer on glass. There is no time lag in epitaxial growth. Instead from the beginning a nearly linear increase of transmission and therefore of crystal amount is observed. After 5.5 h the crystallization is complete. This is less than the time lag of nucleation atthe same temperature. Therefore one does not expect thatnuclei form in the a-Si matrix during the time needed for epitaxial growth.The complete transformation to large grains up to the surface is confirmed by the EBSD map of a 500 nm thick layer, which was taken after 5.5 h annealing at 600°C (Fig. 6). This Figure demonstrates that all grain orientat-ions present in the seed grew epitaxially to the surface. Fig. 7 shows a corresponding TEM cross section image of a 550 nm thick ESPC layer on a seed layer demonstrating that perfect epitaxy occurred. For this to occur a perfectly clean surface of the seed prior to a-Si deposition is crucial.Figure 5: Evolution of transmission of a 500 nm thick a-Si layer deposited at 90 nm/min on a seed layer during annealing at 600°C.Figure 6: EBSD map of a layer system of a 500 nm thick ESPC layer annealed for 5.5 h at 600°C on top of a laser crystallized seed.As for SPC also in ESPC the time needed for complete crystallization, i.e. the epitaxial growth speed, depends on the deposition rate of the a-Si (Fig. 8).Comparing the time needed for ESPC with the time lag (Fig. 3), both as depending on deposition rate, we learn that for 500 nm thick a-Si film at 600°C epitaxial growth is completed within the time lag of nucleation independent of the deposition rate. To completely epitaxially crystallize layers thicker than 500 nm one needs a-Si layers in which the time lag as well as the epitaxial growth speed are as large as possible. This requires a-Si films deposited at rates of about 200 nm/min. In this case the time lag of 14 h is larger than the time needed for epitaxial growth even for a 1µm thick film. Fig. 9 shows the EBSD image of a 1.1 µm thick film crystallized by ESPC at 600°C for 12 h. As for the thinner films large crystals extend up to the surface. Figure 7: TEM cross section image of a layer system consisting of ESPC layer (Si2) on seed (Si1) after annealing for only 7 h at 600°C.Figure 8: Time needed for complete ESPC of a 500 nm thick a-Si layer on top of a seed at an annealing temperature of 600°C.An appreciable reduction in the processing time can be achieved by increasing the annealing temperature. Fig.10 shows the evolution of transmission of a 500 nm thick a-Si layer on a crystalline seed during ESPC at 650°C. The change in transmission immediately starts as is typ-ical for epitaxial growth. After about 1 h the crystalliz-ation is complete. The EBSD map of Fig. 11, however, shows that not all grains of the seed continue to the surface. This is confirmed by the TEM cross section image of Fig 12, where on top of the seed (Si1) about 100 nm of the deposited a-Si epitaxially crystallized (Si2) but the remaining a-Si on top was concerted to fine grained material (Si3). Apparently a competing nucleat-ion process took place in a surface near region. As for the annealing process at 600°C the nucleation and growthkinetics depend on the deposition rate of a-Si.Figure 9: EBSD map of a 1 µm thick ESPC layer deposited at a rate of 180 nm/min and crystallized by annealing at 600°C for 12 h.Figure 10: Evolution of transmission of a 500 nm thick a-Si film on crystalline seed during annealing at 650°C. Figure 11: EBSD map of a 500 nm thick ESPC layer deposited at 47 nm/min and annealed at 650°C for 1.5 h.Figure 12: TEM cross section image of an ESPC layer on seed (Si1) annealed at 650°C for 6 h showing an epitaxially crystallized layer (Si2) and a fine grained layer (Si3).In further work we will investigate if we can deposit a-Si under conditions so that at temperatures between 600°C and 650°C perfect epitaxial growth is possible.4 SUMMARYa-Si layers deposited by electron beam evaporation on top of multicrystalline seed layers on glass were crystallized epitaxially at 600°C up to a thickness of 1.1 µm. In this way we succeeded in preparing a layer system as required for multicrystalline silicon thin film solar cells on glass with grains in the 100 µm range. The preparation process is simple and easily can be up-scaled. By optimizing the deposition and growth conditions we work on reducing the required annealing time and to increase the thickness of the epitaxial layers to well above 1 µm. First observations give hints, that doped a-Si grows faster so that thicker films can be crystallized in even shorter time. In the next steps we will prepare complete solar cells in ESPC silicon layers.5 ACKNOWLEDGEMENTThis paper describes work undertaken in the context of the HIGH-EF project, “Large grained, low stress multicrystalline silicon thin film solar cells on glass by a novel combined diode laser and solid phase epitaxy pro-cess”. HIGH-EF is a Small Scale Collaborative Project supported by the European 7th Framework Programme, contract number 213303.6 REFERENCES[1] M.J. Keevers, T.L.Young, U.Schubert, M.A.Green,Proceedings 22nd European Photovoltaic SolarEnergy Conference (2007), 1783.[2] G. Andrä, J. Plentz, A. Gawlik, E. Ose, F. Falk, K.Lauer, Proceedings 22nd European PhotovoltaicSolar Energy Conference (2007), 1967.[3] G. Andrä, C. Lehmann, J. Plentz, A. Gawlik, E. Ose,F. Falk, 33rd IEEE Photovoltaic Specialists Con-ference (2008), San Diego.[4] U. Köster, phys. stat. sol.(a) 48 (1978), 313.[5] G. Andrä, F. Falk, phys. stat. sol. (c) 5 (2008), 3221.[6] D. Kashchiev, Surf. Sci 14 (1969), 209.[7] G. Andrä, A. Bochmann, F. Falk, A. Gawlik, E. Ose,J. Plentz, 21st European Photovoltaic Solar EnergyConference (2006), 972.。
二硫化钛纳米片材料用于光声成像指导下的肿瘤光热治疗研究
二硫化钛纳米片材料用于光声成像指导下的肿瘤光热治疗研究中文摘要最近，有着不寻常物理化学性质的二维纳米材料在材料科学与工程领域被广泛的研究。
石墨烯是一种由碳原子组成呈蜂巢晶格的二维层状材料，显示出特殊的电子，光学，热力学，和机械性能。
作为一种石墨烯类似物，过渡族金属二维硫化物如MoS2, MoSe2, WS2, WSe2和Bi2Se3，都是由六边形的金属原子夹在两层硫族元素之间组成三明治结构。
近些年这些过渡金属二维硫化物已经发展成为材料科学领域的明星材料，并在包括纳米医药在内的诸多领域有着潜在的应用。
二硫化钛是一种典型的过渡族金属二维硫化物材料，其在电子器件领域或作为一种储氢材料在最近被广泛研究。
在此篇工作中，我们用自下而上的液相合成方法合成了二硫化钛纳米片，然后用聚乙二醇（PEG）修饰，在各种生理溶液中获得很好的稳定性，以与在体外没有明显的毒性。
由于在近红外区有很强的吸收，TiS2-PEG能够被用作光声成像的造影剂。
在对携带肿瘤的小鼠进行全身给药后发现材料能够在肿瘤部位实现很高的富集。
然后我们将TiS2-PEG纳米片用于体内肿瘤光热治疗，在尾静脉注射纳米材料以与外加近红外激光照射的情况下能够实现肿瘤的完全消融。
我们的工作表明，在经过良好的表面修饰后，TiS2纳米片能够作为一种新的光热制剂用于成像指导下的肿瘤治疗。
关键词：二硫化钛光声成像光热治疗Two-Dimensional TiS2 Nanosheets for in vivo Photoacoustic Imaging and Photothermal Cancer TherapyAbstractRecently, two-dimensional (2D) nanomaterials with unusual physical and chemical properties have been extensively explored in materials science and engineering. Graphene, a 2D single layer of carbon atoms of honeycomb lattice structure, as a typical example, has shown exceptional electronic, optical, thermal, and mechanical properties. As the analogues of graphene, transition-metal dichalcogenides (TMDCs) such as MoS2, MoSe2, WS2, WSe2and Bi2Se3, consisting of hexagonal layers of metal atoms sandwiched between two layers of chalcogen atoms, have also become a star in materials science in recent years, showing promising applications in many different areas including nanomedicine.Titanium dichalcogenides (TiS2) is a typical class of TMDC materials and has also been studied recently in electronic devices or as a hydrogen-storage material. In this study, TiS2 nanosheets, are synthesized by a bottom-up solution-phase method and then modified with polyethylene glycol (PEG), obtaining TiS2-PEG with high stability in physiological solutions and no appreciable in vitro toxicity. Due to their high absorbance in the near-infrared (NIR) region, TiS2-PEG nanosheets could offer strong contrast in photoacoustic imaging, which uncovers the high tumor uptake and retention of those nanosheets after systemic administration into tumor-bearing mice. We further apply TiS2-PEG nanosheets for in vivo photothermal therapy, which is able to completely eradicate the tumors on mice upon intravenous injection of TiS2-PEG and the followed NIR laser irradiation. Our work indicates that TiS2nanosheets with appropriate surface coating (e.g. PEGylation) would be promising new class ofphotothermal agent for imaging-guided cancer therapy.Keywords: TiS2photoacoustic imaging photothermal therapy第二章TiS2纳米薄片的制备和表面修饰2.1 引言有着不寻常物理化学性质的二维纳米材料在材料科学与工程领域被广泛的研究1。
太赫兹透镜材料
太赫兹透镜材料
- TPX透镜：近两年，维尔克斯光电陆续推出了不同型号的太赫兹透镜，其中标准化的透镜达23种，主要材质为TPX透镜。
- Ⅲ-Ⅴ族半导体材料：基于Ⅲ-Ⅴ族半导体材料锑化铟（InSb），设计性能优异的单元结构。
随后，采用几何相位和传输相位相结合的方式，设计超透镜单元结构的排布方式与空间取向，采用单层超透镜实现了太赫兹波的宽频带聚焦，有效消除了色差现象。
- 石英晶体：在太赫兹波段，石英晶体是一种常用的材料，能作为原材料生产太赫兹透镜、石英棱镜、太赫兹窗镜、石英波片等。
Ultra Thin Polymer Films for photolithography
Ultra Thin Polymer Films for photolithographic ApplicationsByJuan SchneiderV-P Applications TechnologyMarch 20051.0 Introduction:The aim of this document is to give the necessary information to understand the value of Nanometrix breakthrough on ultra thin film deposition and the solutions for photoresist deposition for the semiconductor industry.Nanometrix ultra thin film deposition technology has produced in the last months layers of photoresist and dielectric polymers with unprecedented thickness control and atomic scale surface roughness bringing attraction to major players of the semiconductor and ultra thin polymer film industries.This document starts with a brief description of photolithography followed by a description of the most common way of photoresist deposition, the spin coating and its characteristics.Nanometrix technology of monolayer, polymer deposition and our resent breakthroughs on ultra thin polymer photoresist coating are explained. Data is also presented from PMMA and PVPH photoresist coatings as well as data obtained with a common dielectric polymer, PVDF. All three polymers were deposited onto silicon wafers of 6 and 8 inches as well as on different flexible substrates of 12 inches wide at a production speed of one m2 per minute. AFM, TEM and spectroscopy techniques are commonly used for analysis and characterization.Finally, a resume of main characteristics of Nanometrix ultra thin photoresist deposition, the conclusion and references end this paper.2.0 Lithography:Photolithography means light-stone-writing in Greek. Is the process by which patterns on a semiconductor material can be defined using light. It is the meansby which the features of integrated circuits are built.Lithography is the single most important driver of Moore’s law. Reducing the size of features patterned, each generation of lithography equipment has enabled faster microprocessors and smaller, less costly integrated circuits on semiconductor wafers [3]. Without the continuous improvements in lithography process and equipment technology that have occurred over the past 30 years, personal computers, cell phones, and the Internet would not be available today. Many important key elements, components and processes make photolithography possible. Wavelength, accurate lenses, photomasks, aligners,photoresist are some of the key components of a complex manufacturing process.Nanometrix technology and applications are mainly on one of these key components: polymer deposition and specifically the photoresist deposition on wafers.A resist is applied to the surface most of the time using a spin coating machine. This device holds the wafer of semiconductor, using a vacuum, and spins it at high-speed (3000-6000 rpm) for a period of 15-60 seconds. The rotation causes the resist to be spread across the surface of the wafer with excess being thrown spun off. [2]After photoresist deposition, this photo sensitive polymer layer is exposed to light. A photomask is used as patterning template leaving an image on the photoresist.The image produced in the photoresist is transferred to other layers on the wafer. These layers are then etched to form a permanent pattern in the film.Once the permanent pattern is created, the photoresist can be removed leaving the desired patterned on the surface of the wafer.3.0 Next Generation LithographyThe quality of the integrated circuit and features are intrinsically dependent on the quality of the photoresist layer.Actual processes and future ones depend on the surface roughness and the thickness consistency over the entire wafer surface. DUV technology is strongly concerned by this issue. EUV is even more concerned by surface roughness of the photoresist. [4]Next generation lithography moves quickly with different approaches. Before the end of last century, DUV was the next generation technology. At the beginning of 2001, EUV and e-beam technology started pulling the interest and major breakthroughs were accomplished, such as reflective photomasks for EUV at 13.4 nm light as well as immersion technology [7]. Most of these experimentation and steps forward have been accomplished by joining forces to tackle a specific problem by the leaders of the industry. Millions of dollars have been and are spent in the race of following Moore’s trend. The most promising next generation processes are EUV, e-beam, Scalpel, Ion-beam and x-ray lithography [6]. Other new paths are at the horizon like nano imprint, Dip-pen, AFM, carbon Nano tubes and molecular memory and semiconductor devices.4.0 PhotoresistPhotolithography light works by passing through a glass mask containing opaque chrome patterns.The patterned light then exposes photoresist, which is subsequently developed to produce small features in the resist.Photoresist is a Light Sensitive Polymer which produces patterns on the Wafer.The Steps in photoresist processing include:- Spin coating the resist on wafers- Pre-baking the resist to remove excess solvent - Exposing the wafers- Post-exposure baking to complete the chemical reactions within the resist - Developing patterns in the resistFigure 1: photoresist current processing under UV lightFabrication of state-of-the-art semiconductor chips requires 20-40 lithography pattern levels all of them etched after the photoresist have been deposited and photosensitized. To align the many layers of an integrated circuit, alignment marks are printed on the wafer.UV exposureNegative resist Positive resist5.0 Method of applying photoresistThe most common method so far for applying photoresist in the lithography industry is the spincoating.Before the photoresist is applied to the substrate, the wafer is cleaned to remove any traces of contamination from the surface of the wafer such as dust, organic, ionic and metallic compounds.Photoresist is applied to the wafer using a spincoating machine. This machine holds the wafer using a vacuum chuck, and spins it at high-speed (approx 3000 rpm) for a period of 60 seconds.A small quantity of resist is initially dispensed onto the wafer. The rotation causes the resist to spread across the surface of the wafer. Preparation of the resist is concluded by a bake step, where the wafer is a gently heated in a hotplate to evaporate the solvents and to partially solidify the resist.6.0 Spinning ArtifactsSpin coating is essentially a pulling process where the photo-resist is spread out and thin down due to the high speed rotation of the silicon wafer. Complex hydrodynamic environment, disjoining pressure, surface tension and evaporation are common physical and chemical variables playing all at the same time during spin coating. 300 mm wafers are the next trend for industrial production causing even more stresses on the film formation due to a bigger diameter of surface to be coated. Striations, edge bead, streaks, material losses and environmental impact are some of the most important problems of spin coating [2]. Here you have a brief description and characteristics of these artifacts:6.1 Striations– ~ 30 nm variations in resist thickness due to non-uniform drying ofsolvent during spin coating– ~ 80-100 mm periodicity, radially out from center of wafer6.2 Edge Bead– residual ridge in resist at edge of wafer– can be up to 20-30 times the nominal thickness of the resist– radius on wafer edge greatly reduces the edge bead height– non-circular wafers greatly increase the edge bead height– edge bead removers are solvents that are spun on after resistcoating and which partially dissolve away the edge bead6.3 Streaks– radial patterns are caused by hard particles whose diameters are greater than the resist thickness6.4 Loses and Cost– Resist cost (~$500/gallon) and most of the resist is spun off the wafer6.5 Environmental Impact– Environmental Impact–Most of the Resist is “wasted”, in the sense that it is spun off the wafer. The high “waste” is required to achieve uniformity in thickness and proper coverage.–Because of precise resist formulations, solvent evaporation during spinning and contamination concerns, it is impossible to recycle resist that is spun from the wafer.Figure 2: Common spin coating resist thickness vs. RPMUnder 100 nm, spin coating starts facing a brick-wall. Thinning down polymers to less than 100 nm in thickness demand supersonic speed at the edge of a 300 mm wafer while the center doesn’t move. Surface tension and disjoining pressure start being extremely important because at few nanometers thick, the solution surface starts seeing each other collapsing with a force that is inversely proportional to the 6th exponential.At the end, spin coating is a complex process that is facing tremendous problems of surface roughness and thickness control due to its intrinsic concept. On the opposite, Nanometrix process is gentle, works by compression and it is a bottom up process.Film Thickness vs RPM0.511.522.533.544.55100020003000400050006000Spin Speed, RPMF i l m T h i c k n e s s , u m7.0 Nanometrix Solution7.1 Nanometrix technologyNanometrix Technology was created to bring to the industry what monolayers have to offer the best: build matter at the molecular scale while driving the process at the human scale. Making monolayers is not a new story. However, making monolayers and ultra thin films in a continuous, well packed, uniform way is essentially a breakthrough.The construction of ultra thin films by the Schneider-Picard (SP) Method is performed by sliding elements on the surface of a liquid until they meet the ultra thin film formation line. In this process, the natural flatness of liquid surface is used, and the flow of it is the driving force that packs the elements one against the other in an orderly and continuous manner. At the same rate the elements are being deposited onto the interface and packed at the formation line, the monolayer is transferred from the liquid surface onto a solid substrate. The whole process works in a dynamic and continuous equilibrium.A monolayer generator was created for this purpose and works in automatic mode. The rate of production exceeds one square meter per minute, which means typically, for a molecular scale element, one billion elements per second that are integrated in the monolayer matrix.So far, monolayers of many kinds and types of elements have been produced and displayed and ultra-thinpolymer films are made at thesame production rate.In this case, the polymer filmthickness can be tailored tohave different values starting at1 nm thick with a meanroughness at the atomic scale(0.1nm), see figures 3, 4 and 5.Thicker polymer films can beengineered with typically asurface roughness under 1 nm,see figures 4 and 5.Figure 3: an AFM image of PMMA ultra-thin film. Theappearance is featureless, except a small dust. Thethickness is 1.2 nm, which fits perfectly the calculated value.The roughness is 0.1 nm.7.2 Diagram of Nanometrix polymer depositionDuring the SP method of ultra thin film deposition, polymer solution is applied at the gas-liquid interface. The receiving phase or the sub-phase is a moving, flat liquid. After injection, the polymer solution thins down in different ways depending on the physico-chemical characteristics of the solvents and liquids. Evaporation and immersion are the main ways in which the solvent concentration, after spreading the polymer, fades down from the gas-liquid inter-phase. This dynamic continuous process transforms the polymer solution from liquid to solid. The solution starts in the liquid phase, moves toward a gel phase and ends at the formation line in the solid phase being deposited simultaneously to the wafer.Evaporationflowing sub-phase Solid plateFigure 4: an AFM image of PMMA ultra-thin film. An area wasanalyzed. The thickness is 1.2 nm, which fits perfectly the calculatedvalue. The roughness is 0.1 nm.Figure 5: A trench was dug in the monolayer. This proved difficult due to the thinness of the layer.The thickness is 1.2 nm.7.3 The substrateSubstrates onto which is deposited the monolayer or the ultra-thin film can be of different nature. Theoretically there are no limitations. Practically, we have deposited the ultra thin films on metal foil, plastic rolls or rigid material like glass slides and silicon wafers.Specifically, for the semi-conductor industry, our process did coat wafers and substrates of different nature and sizes. For the 300 mm generation, the Schneider-Picard (SP) method shows a perfect match.Therefore, this is a continuous process that can be integrated in the present standard line production equipment to create ultra high quality and ultra thin polymer surfaces of dielectrics, barriers and photoresist.Figure 6:an AFM picture of Nanometrix polymer (PVDF) ultra-thin film deposition. The film is 25 nm thick. The average roughness is one nanometer only. The vertical scale has been magnified to see better the monomer vertical orientation.Figure 7: two grooves were dig with the AFM tip into the Nanometrix ultra-thin film matrix. The width of the largest groove (seen in the box) is about 400 nm. A narrower groove was also carved into and parallel to this one. The groove is 50 nm wide only. In both cases, the groove edges are straight and square. The trenches were easily made at the first attempt with in AFM standard operating conditions.7.4 Main characteristics of Nanometrix SP Method for Ultra Thin Polymers Wafer size: Independent on the size of the surface, namely all wafer size currently used by the semiconductor industry.Compression instead of Tension: The polymer is assembled in a compression mode, up to 300 Atm for a film with few nanometers thick. No stress and forces pulling apart the polymer are implicated in the process.Surface roughness: Surface roughness of the polymer, down to the order of an Angstrom.Polymer thickness: It can be tailored with ease to have different values starting at 1 nm thick.Industrial production: Production rate is in the order of m2 per minute.Materials and process:Continuous and gentle.So far, no limits on material handling have been found with the SP method.Loses: Almost 100 % of the polymer solution is used avoiding important loses.8 ConclusionMoore’s Law predicts further development on miniaturisation and major players search continuously for better, faster and less expensive. Nanometrix’ breakthrough SP Method of Ultra Thin Film deposition is the ability to bring a continuous and scalable process for the industry at the molecular scale and up with unprecedented thinness and surface roughness, key factors for the development of DUV and EUV lithography and beyond. Nanometrix is in good position to satisfy actual needs and future ones with barriers, dielectrics and photoresist coatings for present and next generation nanoscale lithography.9 References[1] Manufacturing with DUV lithographyBy S. J. Holmes, P.H. Mitchell and M.C. Hakey, 1996[2] Resist Coating TutorialR.B. Darling[3] The Intel Lithography RoadmapPeter Silverman, Technology and Manufacturing Group, Intel CorporationIntel Technology Journal Vol. 6, Issue 2, 2002[4] Fundamental Challenges for Lithographic RoadmapR. Meagley, Manager, Lithography, Fab Materials Operation (FMO), Intel Corporation[5] Sematech Favoring Scalpel and EUV as Next Best BetsSematech Public Relations DepartmentLucent Technologies, 2002[6] ASML, Immersion TechnologyASML Public Relations Department, 2005[7] Semiconductor ManufacturingRon Leckie & Carl JohnsonInfrastructure, SEMInvest 2001 Tutorial。
大学物理实验报告 英文版
大学物理实验报告Ferroelectric Control of Spin PolarizationABS TR AC TA 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, for instance, 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 thi s 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 polarizationcan be electrically switched to point toward or away from the Fe electrode. We used a half-metallicLa0.67Sr0.33MnO3(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 isspin-coated on the BTO(1 nm)/LSMO(30 nm) bilayer. The nanoindentation is performed with a conductive-tip atomic force microscope, and the resulting nano-hole is filled by sputter-depositingAu/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 latticemis match 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 (CTA FM) (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(1.5 nm)/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. Thelow-resolution, high-angle annular dark field (HAADF) image of the entire heterostructure shows the sharpness of the LSMO/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 of body-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 boundary between 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 BaM4,5 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 the Fe/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 (4.2 K). Figure 2A (middle) shows a typicalresistance–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 hysteresis loop m(H) of a similar unpatterned sample measured with superconducting quantum interference 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 the high-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]. W ithin the simple Jullière model of TMR (23) and considering the large positive spin polarization of half-metallic LSMO (25), thisnegative 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 thi s effect in ourFe/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 betweenspin-polarized empty and filled states is possible. For the Fe/BTO interface, large changes have beenpredicted 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 theJullière model (23), these changes in TMR correspond to a large (or s mall) spin polarization at theFe/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–. Large values 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 havereported the influence of the electrically controlled ferroelectric barrier polarization on thetunnel-current spin polarization. This electrical influence over magnetic degrees of freedom representsa new and interfacial magnetoelectric effect that is large because spin-dependent tunneling is verysensitive 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 in datastorage. Nat. Mater. 6,813 (2007).2.I. Žutić, J. Fabian, S. Das Sarma, Spintronics: Fundamentals and applications. 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 of ferromagnetism. 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 magnetis m in thin-filmferromagnets. Science315, 349 (2007).8.T. Maruyama et al., Large voltage-induced magnetic anisotropy change in a fewatomic 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 ferromagnetism in(Ga,Mn)As. Nat. Mater. 7, 464(2008).11. C. Bihler et al., Ga1−x Mn x As/piezoelectric actuator hybrids: A model system formagnetoelastic 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, Giant sharp andpersistent converse magnetoelectric effects in multiferroic epitaxial heterostructures. Nat.Mater. 6, 348 (2007).15.T. Kanki, H. Tanaka, T. Kawai, Electric control of room temperature ferromagnetismin 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 ferromagnetis m using amagnetoelectric multiferroic. Nat. Mater. 7, 478 2008).17.S. Sahoo et al., Ferroelectric control of magnetis m in BaTiO3∕Fe heterostructures viainterface strain coupling. Phys. Rev. B 76, 092108 (2007).18. C.-G. Duan, S. S. Jaswal, E. Y. Tsymbal, Predicted magnetoelectric effect inFe/BaTiO3 multilayers: Ferroelectric control of magnetism. Phys. Rev. Lett. 97, 047201 (2006).19.M. Fechner et al., Magnetic phase transition in two-phase multiferroics predictedfrom 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 a half-metal/ferroelectricjunction. 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 from tunnelingexperiments. Appl. Phys. Lett. 82, 233 (2003).26.J. P. Velev et al., Magnetic tunnel junctions with ferroelectric barriers: Prediction offour resistance states from first principles. Nano Lett. 9, 427 (2009).27. F. Yang et al., Eight logic states of tunneling magnetoelectroresistance inmultiferroic tunnel junctions.J. Appl. Phys. 102, 044504 (2007).28.Materials and methods are available as supporting material on Science Online.29.V. Garcia et al., Giant tunnel electroresistance for non-destructive readout offerroelectric 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 electrically controlledindentation with an atomic force microscope for nanocontact elaboration. NanoLett. 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 kagome systemsYBaCo4O7 and YBaCo3M O7 (M=A l, Fe). Phys. Rev. B 80, 085111 (2009).34.M. Abbate et al., Soft-x-ray-absorption studies of the location of extra chargesinduced 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: Giantelectroresistance 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 tunnel junctions withhalf-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 microscopymeasurements. This study was partially supported by the France-U.K. Partenariat HubertCurien Alliance program, the French Réseau Thématique de Recherche Avancée Triangle de la Physique, the European Union (EU) Specific Targeted Research Project (STRep) Manipulating the Coupling in Multiferroic 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 (A NR) Oxitronics, French ANR A licante, the European Enabling Science and Technology through European Elelctron Microscopyprogram, and the French Microscopie Electronique et Sonde Atomique network. X.M.acknowledges support from Comissionat per a Universitats i Recerca (Generalitat de Catalunya).。
表柔比星超顺磁性氧化铁纳米粒(EPI-SPION)经皮递药抗肿瘤作用研究
表柔比星超顺磁性氧化铁纳米粒（EPI-SPION）经皮递药抗肿瘤作用研究皮肤癌是人类最常见的恶性肿瘤,每年的新患人数超过100万,尤其在浅色人种中,皮肤癌的发病率日趋增高。
传统的肿瘤化疗及放疗普遍存在靶部位药物浓度低、药物毒副作用大等缺陷,治疗的效果不理想。
目前对类似恶性黑色素瘤这一类浅表性皮肤癌仍欠缺有效的药物治疗手段。
在纳米药物抗肿瘤研究中,磁性靶向递药系统的给药方式一般采用静脉注射,要求药物粒子能自由通过最小的毛细血管,粒子易被肾脏排泄或聚积到骨髓中,导致抗肿瘤药物的毒性加大。
另一方面磁性药物靶向治疗的磁场装置设计和实施也是一个限制磁性靶向递药系统发展和应用的关键问题。
本研究尝试采用经皮递药系统克服上述不利限制,探索皮肤肿瘤治疗的新途径。
超顺磁氧化铁纳米粒(SPION)的研究从最初侧重于MRI造影发展到目前的靶向给药、治疗和造影等。
具有肿瘤靶向SPION由于其在MRI诊断、热疗及靶向给药中潜能,是目前靶向药物研究中很热门的课题。
本研究以表面具有功能基团的γ-Fe2O3纳米粒为内核,将具有伯氨基的肿瘤化疗药物表柔比星(EPI)嫁接于γ-Fe2O3纳米粒,得到直径十几个纳米、具有超顺磁性的表柔比星超顺磁性纳米粒(EPI-SPION),以磁导向经皮靶向递药技术和色素瘤细胞模型为主要研究对象,阐述在外加磁场作用下,EPI-SPION经皮给药穿透皮肤角质层及靶向皮肤肿瘤细胞的能力和作用机制。
结合肿瘤药物治疗学、分子生物药剂学和材料化学的思路和技术方法,从组织及细胞水平探讨超顺磁性纳米粒经皮吸收靶向肿瘤细胞的转运机制和动力学行为。
首先通过共沉淀-交联耦合法制备功能化γ-Fe2O3纳米粒,经仿生化表面修饰后将药物嫁接至纳米粒表面,得到性质稳定的EPI-SPION递药系统,并试图通过调控外磁场作用方式和强度,高效递送抗肿瘤药物至肿瘤局灶区。
我们前期得到的载药SPION水合粒径大小为16±2nm,主要理化性质研究表明,该纳米粒具有类球形单畴晶格形貌,保持良好磁性效能,饱和磁化强度大于70 emu/g,矫顽力几乎为零,具有良好的超顺磁性能。
Epitaxial Growth of Thin Films in Solid-Solid Inte
பைடு நூலகம்421
20 Epitaxial Growth of Thin Films
Harald Brune
This chapter gives an introduction to the epitaxial growth of thin films on solid substrates. The term epitaxy refers to the growth of a crystalline layer on (epi) the surface of a crystalline substrate, where the crystallographic orientation of the substrate surface imposes a crystalline order (taxis) onto the thin film. This implies that film elements can be grown, up to a certain thickness, in crystal structures differing from their bulk. If film and substrate have the same crystal lattices, but different lattice constants, the film will be under strain, that is, it will have a slightly different lattice constant than in its own bulk. Both effects, together with the electronic hybridization at the interface, lead to novel properties.
调制型AuSn薄膜微观形貌及合金化工艺
第58卷第3期 2021年3月欲纳电子技术Micronanoelectronic TechnologyVol. 58 No. 3March 2021〇〇I：1()_1325〇/j_ cnki. Wndz. 2021. 03. 012♦加工、测量与设备,调制型A ii/Sn薄膜微观形貌及合金化工艺沓世我U2，李淑华U2，李保昌U2，罗俊尧U2，杨瞾1’2’3(1.广东风华高新科技股份有限公司，广东肇庆526060;2.新型电子元器件关键材料与工艺国家重点实验室，广东肇庆526060;3.华南理工大学材料科学与工程学院，广州 510641)摘要：采用A u和S n单质金属靶，通过直流磁控溅射法制备调制型Au/S n薄膜（薄膜层数为 3〜21)，经快速退火后，实现单质多层薄膜的合金化。
主要研究了 A u/S n薄膜微观形貌和合金 化工艺控制。
结果表明，当固定薄膜总厚度为2 p m时，320 °C下退火10 m in后，膜层表面粗糙 度与薄膜层数呈反比。
薄膜层数较少（n二3)、调制周期厚度较大时，由于A u与S n间扩散不完 全，合金化不充分，造成薄膜表面起伏较大，其均值粗糙度最高达到188.5 nm。
随着薄膜层数 不断增加，调制周期厚度减小到纳米级，薄膜也更加致密、平整，n =21时，320 °C下退火10 m in后均值粗糙度仅为29.7 nm。
优化合金化工艺过程中，采用了不同的优化方法，包括增加退火温度、延长保温时间、降低薄膜总厚度和调制周期，最终在膜层厚度为700 nm、退火温 度为320 °C、退火时间为10 m in的工艺条件下，获得了表面致密平整、合金化充分及金锡质量 比约为8() : 20的合金薄膜，均值粗糙度仅为23. 5 nm。
关键词：A u/S n薄膜；磁控溅射；微观形貌；合金化；调制型中图分类号：0484. 1文献标识码： A 文章编号：1671-4776 (2021) 03-0264-07Microstructure Morphology and Alloying Process ofModulated Au/Sn Thin FilmsTa Shiw o1’2，Li S huhua〗’2，Li Baochang1’2，Luo Junyao〗’2，Yang Zhao1’2’3(1. G u a ng d o n g F eng h ua A d v a n c e d T echnology H o ld in g Co., L td., Z h aoq in g526060? C h in a；2. S ta te K e y L ab o ra to ry o f A d v a n c e d M a teria ls and E lectronic C om ponents^, Z h a o q in g526060, C hina；3. School o f M a teria ls Science a n d E n g in e e r in g，South C hina U n iv ersity o f T e c h n o lo g y，G uangzhou510641, C h in a)Abstract ：The Au and Sn single m etal targets were used to prepare m odulated A u/Sn thin films (the num ber of thin film layers is 3 _ 21 ) by DC m agnetron sputtering. After rapid annealing, the alloying of single substance m ultilayer thin films was realized. The m icrostructure m orphology and alloying process control of the A u/S n thin films were mainly studied. The results show that when the fixed total thickness of the thin film is 2 ^tm, after annealing at 320 °C for 10 m in, the thin film surface roughness is inversely proportional to the num ber of thin film layers. When the num ber of thin film layers is small (n = 3) and the m odulation period thickness收稿日期：2020-09-25通信作者：杨翌，E-mail: **************************264沓世我等：调制型A u/S n薄膜微观形貌及合金化工艺is larger, the diffusion between Au and Sn is not com plete, and then the alloying is not sufficient, thus thin film surface fluctuates greatly and the mean m axim um roughness is up to 188. 5 nm. W ith the num ber of the thin films increases, the modulation period thickness reduces to nanoscale, and the thin film becomes denser and sm oother. When n - 2\ y the mean roughness is only 29. 7 nm after annealing at 320 °C for 10 min. During optimizing the alloying process, different optimization m ethods were used, including increasing the annealing tem perature, extending the holding tim e, reducing the total thin film thickness and m odulation period. Finally, under the process conditions of a thin film thickness of 700 nm, an annealing temperature of 320 °C and annealing time of 10 min, an alloy thin film with a dense and smooth surface, sufficient alloying and a mass ratio of Au and Sn of about 80 ：20 was obtained, the mean roughness is only 23. 5 nm.Key words：A u/S n thin film；m agnetron sputtering; m icrostructure m orphology；alloying；modulation typeEEACC：0520〇引言随着微电子技术与半导体封装要求的不断提高，人们越来越重视用于高可靠性封装的焊接材料。
电子磁性材料 Fe 3 Si薄膜的研究进展
电子技术・ Electronic Technology132 •电子技术与软件工程 Electronic Technology & Software Engineering【关键词】Fe 3Si 薄膜 制备方法 研究现状1 引言纵观软磁材料的发展历程，关于金属软磁材料国内外的研究主要在Finement 合金、Nanoperm 合金和Hitperm 合金软磁材料，它们具有高的饱和磁化强度、低频段较高的初始磁导率和低损耗等优点，主要可用于音频、视频磁头和卡片阅读器磁头等方面。
而Fe-Si 材料蕴藏丰富、高效环保，已经成为通讯、电力工业和电子工业中不可或缺的软磁材料。
Fe 3Si 薄膜材料具有磁导率高、矫顽力Hc 比较小、饱和磁感应强度Br 高、高频下具有优异磁性等优点。
再着Fe 3Si 薄膜具有巨磁阻的性质，当外磁场作用时，电阻会发生剧烈变化，有望用于巨磁阻磁头、巨磁阻存储器、各种磁传感器上。
2 Fe 3Si薄膜的制备方法随着Fe 3Si 薄膜制备工艺的不断发展，目前国内外Fe 3Si 薄膜的制备方法主要有分子束外延法、脉冲激光沉积法、离子束合成法、射频溅射法等。
3 Fe 3Si薄膜的国内外研究现状3.1 国外研究现状1993年5月，H.Liou 等[1]开始通过多腔室分子束外延法在GaAs(001)衬底上外延生电子磁性材料Fe 3Si 薄膜的研究进展文/高赐国 谢晶 谢泉 刘栋长超薄Fe 3Si 薄膜的研究，制备的薄膜厚度从2-210纳米层（ML ），膜厚即使薄到2ML 也具有铁磁性，当膜厚为10ML-50ML 时，样品的矫顽力随膜厚的减小逐渐增大，而当膜厚为2ML-10ML 时，矫顽力变化趋势则表现出相反特性，且随膜厚增加，其磁滞回线矩形度越好。
2004年6月，D.Nakagauchi 等[2]通过脉冲激光沉积在硅和石英衬底上生长了铁磁性Fe 3Si 薄膜，当衬底温度为300℃时，硅和石英衬底都能很好的生长出单一的多晶Fe 3Si 薄膜，而当衬底温度超过400℃时，硅衬底中的硅会扩散到薄膜中去，形成Fe 3Si 和FeSi 混合相。
The Materials Science of Thin Films
The Materials Science of Thin Films薄膜材料科学薄膜材料科学是一个广泛的领域，可以应用于多种工业和科学领域。
与传统的物理学、化学和材料科学不同，薄膜材料科学的研究主要集中在二维和三维薄膜的制备、特征和性质研究。
在薄膜科学的研究中，材料的物理、化学和电学性质是需要考虑的重要参数。
在这篇文章中，我们将探讨薄膜材料科学的重要性及其应用。
1.薄膜的制备薄膜制备的方法有很多种，其中包括物理气相沉积、化学气相沉积、电化学沉积、溶胶-凝胶法和分子束外延等方法。
这些方法都是为了在一个非常小的区域内制造材料厚度很薄的膜，这些薄膜在可见光谱和电学特性方面与材料的块体状相似。
2.薄膜的结构薄膜可以是单晶、多晶或非晶态的。
单晶薄膜的结构与块体材料相似，有一个确定的晶面方向。
多晶薄膜由很多不同晶体方向的小结晶组成。
非晶态薄膜在结构上没有规律性。
这些不同结构形式的薄膜在各自领域中都有不同的应用。
3.薄膜的性质与块体材料相比，薄膜材料的一些基本电学特性会有显著的变化，这些变化可以归因于材料中谷能和导带能位于多晶颗粒之间。
激子的引入导致了光响应和电子传输的快速削减，从而产生了更快的响应速度。
薄膜还可以表现出良好的光吸收特性，可以在太阳电池和其他光电设备中应用。
薄膜对于某些元素和化合物的高选择性敏感性也使其成为很好的气体传感器材料。
总之，薄膜材料科学是一个广泛的领域，可以应用于多种工业和科学领域，比如在计算机芯片、有机LED、太阳能电池、传感器、纳米加工和生物医学设备等领域。
未来，薄膜科学将继续发展，我们也期待着更多的创新和应用，为社会和人类的未来做出贡献。
Ultrathin Films for Energy Conversion and Storage
Ultrathin Films for Energy Conversionand Storage近年来，随着能源紧张和环境污染等问题日益突出，人们对新能源技术的研究越来越深入，其中超薄膜技术的研究与应用成为一个热点领域。
超薄膜技术是一种能将材料减薄至亚纳米级别的制备技术，具有薄、轻、柔性和高表面积等优点，被广泛应用于能源转换和储存领域。
一、超薄膜在太阳能领域的应用太阳能光伏发电是目前最为普及的新能源技术之一，然而现有太阳电池的能量转换效率并不高，而且成本较高。
超薄膜技术的应用可以极大地提高这种太阳电池的能量转换效率。
利用超薄膜技术制备出的材料可以大幅提高光电转换效率，因为薄膜可以增加太阳能的吸收量，并且几乎不会减少光子的可用能量。
同时，超薄膜技术制造太阳电池的成本也可以降低，因为这种技术相比传统的硅基太阳电池成本更低。
二、超薄膜在储能领域的应用随着电动车和储能系统的兴起，超薄膜技术在储能领域的应用也越来越广泛。
超薄膜材料在电池领域中的应用可以提高电池的能量密度，这对于电动车和储能系统等应用来说是至关重要的。
相比传统的电池，超薄膜电池具有更大的表面积和更高的能量密度，能够在相同或更小尺寸下储存更多的能量。
此外，利用超薄膜技术制备出的电池具有更好的充电和放电速度，可以满足电动车、移动通讯和储能等行业的需求。
三、超薄膜在电解水制氢领域的应用水的分解产生的氢气可以成为非常清洁、无污染的能源，因此远程地区和岛屿上的氢能电力系统也已经成为了一个具有重要意义的研究领域。
超薄膜铁氧体材料可以用于催化水的分解，产生氢气。
与传统的电解水系统相比，利用超薄膜技术制备的电解水系统可以在更低的电压下产生氢气，同时还可以实现更高的氢气产量和更长的寿命。
因此，这种技术在未来的氢能源发展中应具有很高的应用前景。
总之，利用超薄膜技术制备出的材料在能源转换和储存领域具有广泛的应用前景。
虽然目前这种技术的研究仍处于早期阶段，但已经有许多有前途的材料在这个领域得到了广泛关注。
Epitaxial Growth of Semiconductor Thin Films
Epitaxial Growth of Semiconductor ThinFilms半导体薄膜的外延生长随着集成电路的迅猛发展，半导体材料的研究和制备技术也日益成熟。
半导体薄膜作为现代电子器件的基本材料，具有广泛的应用前景。
而半导体薄膜的外延生长技术也是半导体材料研究的一个重要方向。
本文将从半导体薄膜的概念、外延生长技术的分类和应用等方面进行探讨。
一、半导体薄膜的概念半导体薄膜是指在特定的材料基底上生长的薄层半导体材料。
由于薄膜厚度不同，其电学、光学性质也各异，因此，半导体薄膜具有很多优良的特性。
例如，可以控制其电学性质、改变其光学特性等。
半导体薄膜广泛应用于太阳能电池、光电器件、模拟集成电路、量子器件等领域。
二、外延生长技术的分类外延生长技术是指在单晶硅或其它半导体材料上，通过化学反应或热解反应的方式，沉积单晶半导体材料，从而完成半导体薄膜的生长过程。
外延生长技术可以分为熔体外延、气相外延、分子束外延等三种主要方式。
1.熔体外延熔体外延是指将原料直接熔化，当温度达到体系的熔点时，在基底上外延生长一个薄膜。
熔体外延适用于制备晶格匹配良好、相互溶解度大的半导体材料。
例如，可以利用熔体外延技术在氧化铜基底上制备GaAs、InP等薄膜。
2.气相外延气相外延是指利用半导体单质和半导体化合物在高温状态下分解产生的气态化学物质沉积在基底上完成外延生长。
气相外延适用于制备高质量的半导体材料。
例如，可以利用气相外延技术制备氧化硅、氧化锌、氢化硅等半导体薄膜。
3.分子束外延分子束外延是指利用高真空环境下的纯单质或其相应的化合物，通过分子束加速器将单个分子或分子团束投射到基底上，完成单晶薄膜的生长。
分子束外延技术具有高纯度、较小厚度等优点，适用于制备高质量、低缺陷密度的半导体薄膜。
例如，可以利用分子束外延技术在GaAs基底上制备AlGaAs、InGaAs等半导体薄膜。
三、外延生长技术的应用外延生长技术已经成为半导体制备的重要手段之一，在众多的应用领域中得到了广泛应用。
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