2014年3月21日金曜日

早稲田大学　逢坂哲彌 研究室の博士論文のコピペ疑惑

注：正しい方法で行えば「コピペ」もOK（弁護士ドットコムより）
他者著作物との類似性が見られた博士論文  （計２３報）：　
（コピペを効率的な博士論文執筆方法として取り入れた可能性のある賞されるべき事例）
　常田聡 研究室：　小保方晴子、松本慎也、古川和寛、寺原猛、岸田直裕、副島孝一、寺田昭彦（ラボ内コピペ）　（計７名）
　西出宏之 研究室：　義原直、加藤文昭、高橋克行、伊部武史、田中学、小鹿健一郎 （計６名）
　武岡真司 研究室：　藤枝俊宣、小幡洋輔、寺村裕治、岡村陽介（ラボ内コピペ）  （計４名）
　逢坂哲彌 研究室：　奈良洋希、蜂巣琢磨、本川慎二（計３名）
　平田彰 研究室：　吉江幸子（ラボ内コピペ）、日比谷和明（ラボ内コピペ） （計２名）
　黒田一幸 研究室：　藤本泰弘 （計１名）
 (早稲田大学リポジトリ）　（その他の早稲田理工の研究室も網羅的に調査中）

当記事の公益目的：　理化学研究所の調査委員会によりSTAP細胞論文における捏造・改ざんの研究不正や他者著作物からの文章のコピペが認定された小保方晴子氏は早稲田大学理工学術院の先進理工学研究科で学位を取得した後、理化学研究所研究員として採用されていました。小保方晴子氏の早稲田大学における博士論文についても、冒頭２０ページ近くの文章がNIHのサイトからのコピペであること、各章のリファレンスまでもがコピペであり本文と全く対応しておらず本文中にはリファレンス番号が記載されていないこと、複数の実験画像がバイオ系企業サイトに掲載されている実験画像と類似していることなどの多数の問題点が判明しています。これらの当然気付かれるべき問題点は早稲田大学における博士論文の審査では見過ごされていました。よって、小保方氏のSTAP細胞論文における様々な問題は、小保方氏個人が責められるべきものではなく、早稲田大学の教育環境や学位審査システムの特質性にもその要因が在ります。STAP細胞論文自体の研究や、その研究結果の再現性確認実験には多額の公的研究費や研究者の貴重な時間が費やされました。公益目的の観点から、二度と同様の問題が起こらないように対策をとるためには、早稲田大学の教育環境や学位審査システムを精査する必要があります。その手がかりを得るために、当記事では、自主的に網羅的調査をしようとしない早稲田大学に代わり、読者の調査協力の下に第三者の観点から「他者の著作物からのコピペが博士論文を効率的に書くための一方法として早稲田大学で普及していたのかどうか。」を網羅的に検討することにします。また、コピペが博士論文などの著作物を効率的に執筆するための一方法として認められるのかどうか、推奨されるべきかどうかの問題は社会一般公共の利害に関することから、専ら公益目的の観点から早稲田大学の事例をもとに考えていきたいと思います。

適切な引用（コピペ）とは？：　文化庁は、以下の７項目を、他人の主張や資料等を「引用」する場合の要件としています。
ア 既に公表されている著作物であること
イ 「公正な慣行」に合致すること
ウ 報道，批評，研究などの引用の目的上「正当な範囲内」であること
エ 引用部分とそれ以外の部分の「主従関係」が明確であること
オ カギ括弧などにより「引用部分」が明確になっていること
カ 引用を行う「必然性」があること
キ 「出所の明示」が必要（コピー以外はその慣行があるとき）
(文化庁長官官房著作権課 著作権テキスト 平成２２年度版 　PDFファイル　の　「§8. 著作物等の「例外的な無断利用」ができる場合 ⑧ ア、「引用」（第32条第1項」　より引用）

早稲田大学が先進理工学研究科の博士論文について調査開始（2014年4月7日）
2014年4月7日：　皆様のご協力のもと、当ブログにおいて早稲田大学の博士論文のコピペ問題を検証し続けたことにより、早稲田大学が先進理工学研究科の280本の全ての博士論文を調査することを決定しました。また、このことを数多くの大手新聞やNHKを含む大手放送局もニュースとして取り上げました。博士論文のコピペ問題を検証してきたことが、公共性を助成し公益目的を有するということが一般的にも認められました。ご協力ありがとうございました。博士論文の序章（イントロダクション、背景）における他者著作物からの丸ごとコピペが適切な引用にあたるのかどうかについて、早稲田大学がどのように判断するか注目したいですね。

以下、関連ニュースです。
2014年4月7日（日本経済新聞）：　早稲田大、博士論文280本対象に不正調査 小保方氏が学位取得の先進理工学研究科で
2014年4月7日（産経新聞）：　全博士論文を対象に調査 小保方氏所属の早大先進理工学研究科
2014年4月7日（The Huffington Post Japan）：　小保方さん問題で早稲田大学、博士論文280本を調査 不正あれば学位取り消しも （写し）
NHK：　早大 小保方氏出身の研究科 論文調査
日テレニュース：　早大 他の博士論文２８０本でも不正を調査
ＴＢＳ： 早大・小保方氏の所属学科、全博士論文の不正調査へ
2014年4月7日（その他）：　スポニチ、千葉日報、日本海新聞、SankeiBiz、日刊スポーツ、Sanspo、デイリースポーツ、北國新聞、ZAKZAK、財経新聞
2014年4月8日 Retraction Watch: Waseda University checking dissertations for plagiarism in wake of STAP stem cell misconduct finding
2014年4月15日 The Japan News by Yomiuri: Waseda graduate school probes 280 doctorate theses

以下、関連サイトです。
2014年3月14日：　早稲田大学の理工系におけるコピペ文化について
2014年3月26日（日刊ゲンダイ）　：　コピペどころか論文買う学生も…横行する「卒論ゴースト」
2014年3月27日：　早稲田大学の理工系の非コピペ文化について／電気・情報生命工学科の学生から （写し）
小保方晴子氏の博士論文のコピペ問題に関する報道
2014年3月18日（J-CASTニュース）：　早大で次々に「論文コピペ疑惑」が浮上 小保方氏は先輩の手法を見習った？
2014年3月18日（J-CASTニュース）：　「小保方博士論文」審査員のハーバード大教授「読んでないし頼まれてもいない」
2014年3月20日（日刊工業新聞）：　米ハーバード大教授、小保方氏の博士論文読まず
2014年3月20日（J-CASTニュース）：　ハーバード大教授「小保方氏の博士論文読んでない」 衝撃発言に東浩紀氏「本当なら早稲田は終わりだ」
2014年3月20日（朝日新聞）：　小保方さんの博士論文「読んでない」 学位審査の米教授
2014年3月21日（東京スポーツ）：　小保方氏「最後の味方」も不穏な発言
2014年3月26日（時事通信）：　早大が本格調査へ＝小保方氏の博士論文
2014年3月27日（弁護士ドットコム）：　小保方さんに教えてあげたい!? 弁護士が伝授する「論文引用」の正しいやり方 （写し）
2014年3月27日（弁護士ドットコム）：　小保方さん「コピペ論文」で揺れる早稲田大学――法学部に広がる「モカイ文化」とは？ （写し）
調査１：奈良洋希氏（早稲田大学の逢坂哲彌氏の研究室）の博士論文における文章のコピペについてのまとめ

著者：　奈良洋希　（現　早稲田大学 理工学術院 次席研究員（研究院講師））
論文題目：　「相分離電解質を用いたリチウム二次電池材料の開発」
http://dspace.wul.waseda.ac.jp/dspace/handle/2065/28688
出版日：　2008年
審査員：
 (主査) 早稲田大学教授工 学博士( 早稲田大学) 逢 坂哲彌
 (副査) 早稲田大学教授工 学博士( 早稲田大学) 西 出宏之
 (副査) 早稲田大学教授工 学博士( 早稲田大学) 黒 田一幸
 (副査) 早稲田大学教授工 学博士( 早稲田大学) 菅 原義之
 (副査) 早稲田大学教授博 士( 工学) 早稲田大学本 間敬之
 (副査) 早稲田大学准教授博 士( 工学) 早稲田大学門 間聰之
 (副査) 学外審査員 Rome 大学教授Ph.D (Uni v.Rome ) Bruno Scrosati

概要 （写し）
審査報告書 （写し）
論文本文 （写し）（写し２）（写し３）
逢坂哲彌研究室の奈良洋希氏の博士論文の第一章が、
約67000文字、10000単語のコピペであることが判明。
図表もほとんどがコピペ。
引用さえしていれば他者の有料論文も丸ごとコピペして無料公開してもよいのでしょうか？

奈良洋希氏の博士論文のchapter 1 の、1.1 の下記文章の一部や図は、J.-M. Tarascon氏の論文からのコピペです。

  同一文章１ ：黄色でハイライトされた部分がJ.-M. Tarascon氏の論文（Nature 414, 359-367 (15 November 2001) | doi:10.1038/35104644）と同一文章。 黄緑色がWoodbank Communications Ltdのサイト（http://www.mpoweruk.com/chemistries.htm）の文章と同一文章。
1.1. Back Ground(1, 2)
　A battery is composed of several electrochemical cells that are connected in series and/or in parallel to provide the required voltage and capacity, respectively. Each cell consists of a positive and a negative electrode (both sources of chemical reactions) separated by an electrolyte solution containing dissociated salts, which enable ion transfer between the two electrodes. Once these electrodes are connected externally, the chemical reactions proceed in tandem at both electrodes, thereby liberating electrons and enabling the current to be tapped by the user.The amount of electrical energy, expressed either per unit of weight (Wh/kg) or per unit of volume (Wh/l), that a battery is able to deliver is a function of the cell potential (V) and capacity (Ah/kg), both of which are linked directly to the chemistry of the system. The cell potential is determined by Gibbs free energy change, -ΔG.
zEF = -ΔG
where z is electron number for reaction, F is Faraday constant. When the anode materials which have strong reducing power and the cathode materials which have strong oxidizing power are combined, -ΔG increases, and the electromotive force increases. The reducing and oxidizing power are expressed by the electrode potential, therefore the electromotive force is determined by the potential difference between the anode and cathode.
　The capacity of a battery is determined by electrochemical equivalent, because the electrochemical reaction follows Faraday law. Therefore the high electromotive force and capacity battery that is high energy density battery can be achieved by applying the anode which has low redox potential and small electrochemical equivalent and the cathode which has high redox potential and small electrochemical equivalent.
　As described above, the voltage and current of a battery is directly related to the types of materials applied in the electrodes and electrolyte. The propensity of an individual metal or metal compound to gain or lose electrons in relation to another material is known as its electrode potential. Thus the strength of oxidizing and reducing agents are indicated by their standard potential. Compounds with a positive electrode potential are used for anodes and those with a negative electrode potential for cathodes. The larger the difference between the electrode potentials of the anode and cathode, the greater the electromotive force of the battery and the greater the amount of energy that can be produced by the battery.
　High specific energies are theoretically available with most metals in the first, second third principal groups of the Periodic Table, but only metallic lithium(3)　has found wide application. Among the various existing batteries (Fig. 1.1.1), lithium based batteries – because of their high energy density and design flexibility– currently outperform other systems, accounting for 63 % of worldwide sales values n portable batteries(4).


コピペ図 ：Figure 1.1.1 Comparison of the different battery technologies in terms of volumetric and gravimetric energy density.(1)　は、J.-M. Tarascon氏の論文のFig.1のコピペ



奈良洋希氏の博士論文のchapter 1 の、1.2 polymer electrolytes において、文献は引用してあるものの、 Felix B. Dias氏の論文（Journal of Power SourcesVolume 88, Issue 2, June 2000, Pages 169–191、DOI 10.1016/S0378-7753(99)00529-7）
http://www.sciencedirect.com/science/article/pii/S0378775399005297
からの長文、図、表の不適切な引用（数パラグラフ丸ごとコピペ）があります。

同一文章２ ：黄色でハイライトされた部分がFelix B. Dias氏の論文と同一文章。 
1.2. Polymer electrolytes(5)  
1.2.1. Basic properties of polymer electrolytes  
Although, tremendous research activities by Sony leading to commercialization of lithium batteries incorporating liquid electrolytes took place over the last decade(6), many advantages of solid polymer electrolytes over their liquid counterparts such as organic solutions and inorganic and molten salts can be seen. The possibility of internal shorting, leaks, and producing combustible reaction products at the electrode surfaces, existing in the liquid electrolytes, is eliminated by the presence of a solid polymer electrolyte. Nevertheless, the polymer electrolytes should exhibit ionic conductivities, at least, of the order of 10-3 to 10-2 S/cm at room temperature and play the role of a separator, played by the liquids. The polymer should also allow good cycle lives, low temperature performances, and good thermal and mechanical strengths in order to withstand internal temperature and pressure buildup during the battery operation. The polymers, in general, being light-weight and non-combustible materials can be fabricated to requirements of size and shape, thus offering a wide range of designs. Since stable thin films of the polymers can be easily made, high specific energy (low-mass) and high specific power (less volume) batteries can be expected for use in electric devices and EV. Internal voltage drops may be low, about 50 mV at 10 mA/cm2, when films as thin as 40 μm are made(7, 8).
 The field of polymer electrolytes has gone through three stages: “dry solid systems”, “polymer gels”, and “polymer composites”. The “dry systems” use the polymer host as the solid solvent and do not include any organic liquids. The “polymer gels” contain organic liquids as plasticizers which with a lithium salt remain encapsulated in a polymer matrix, whereas the “polymer composites” include high surface area inorganic solids in proportion with a “dry solid polymer” or “polymer gel” system. In this thesis, more detailed aspects of the polymer gels system with individual examples are discussed below. The advancement of the field can be found in literature publications and reviews(7, 9-11). One of the important properties of a polymer electrolyte leading to its development activity is the ionic conductivity. Temperature dependence on conductivity of amorphous polymer electrolytes generally follows the Vogel-Tammann-Fulcher (VTF) equation （式　略）where, T0 is the glass transition temperature of the polymer electrolyte measured by DSC, T is the temperature of measurement, A is the pre-exponential factor, and E is the activation energy which can be evaluated either from the configurational entropy theory or the free-volume theory, and hence relates to the segmental motion of polymer chains(12). The ionic conductivity is usually measured by AC impedance techniques(13).
 Another important property of the polymer electrolyte is the lithium ion transference number (tLi+) which ideally for lithium battery applications should be unity. A value of tLi+ lower than 1 would tend to develop concentration gradients at electrode surfaces leading to limiting currents. Thus, both the parameters, ionic conductivity and lithium ion transference number are important in order to choose a polymer electrolyte for a practical lithium battery. The maximum power obtainable in a lithium cell can be related to the conductivity of the electrolyte, whereas, the maxi-mum limiting current that can be drawn from the cell and the cycleability of the cell can be related to the tLi+. The lithium transference number (and its associated diffusion coefficient) measurements are usually made using techniques such as concentration cell method(14, 15), Tubandt method(16, 17), LiNMR(18), and electrochemical(19, 20), and electrophoretic NMR techniques(21, 22). The techniques have their advantages and disadvantages and differ by theoretical models used for interpretation of data. Analysis of some of the problems and limitations associated with some of the techniques have been described by Fritz and Kuhn(23).　An understanding to the interactions of the various species in the polymer electrolytes has made possible to choose appropriate polymer hosts, complexing salts and salt concentrations so that the ionic conductivities are optimized. Evidence of salt dissolution into a polymer host is mainly based on spectroscopic methods such as IR and Raman(24, 25). The anion-cation and ion-polymer associations in both crystalline and amorphous phases have been studied using both these which monitor changes in the vibrational modes of the polymer host and the anions(26-29). For example, in the case of lithium triflate (LiCF3SO3)(28) and lithium bis(trifluoromethane sulfonyl)imide [LiTFSI] complexes(27) with oxyethylene containing polymer chains, SO3 symmetric stretching modes, CF3 bending modes, and Li-O stretching modes were associated to ion-ion interactions and information on ion pairing and ion association was obtained(30).
 X-ray and neutron diffraction methods have given de-tailed information on the structural aspects of coordination around the ions in crystalline phases of polymer electrolytes(31-33). The EXAFS has revealed the local environment of the ion(34) in both crystalline and amorphous phases, including the existence of ion pairs such as [MX]0, [M2X]+, [MX2]-, etc. and the determination of bond lengths of the ion-polymer interactions(35). The EXAFS technique is more useful when theoretical models depicting the most probable situations are at hand so that fitting of the EXAFS spectrum of known structures similar to the chemical environment of the polymer electrolyte is possible. 
Stability of polymer electrolytes 
 The electrochemical stability of an electrolyte is one of the essential parameters when rechargeable lithium batteries are concerned. The instability in the electrolyte is known to bring out irreversible reactions and capacity fading in the battery(36). The stability commonly expressed as “electrochemical stability window” of the polymer electrolyte (units of volts) should be the same as that of the electrode potential or higher so that overpotentials during re-charge are compensated by the higher stability window. For example, for a 4-V lithium battery a window of at least 4.5 V vs. Li / Li+ in the polymer electrolyte is required for compatibility with a lithium metal anode and lithium-ion intercalating electrode materials. Thermal and mechanical stabilities of a polymer electrolyte during charge-discharge cycles are vital for a safe and endurable battery. During charge and discharge, heat is known to get generated in the battery which increases the surface area of an existing passive layer at the electrode surface(37). The heat can also melt and degrade the polymer electrolyte within the battery and cause internal short circuits. Increase in heat due to environmental factors during storage increases the self discharge reactions in the battery and shorten life. 
1.2.2. Polymer gel electrolytes  
In 1973, the first measurements on conductivities of poly(ethylene oxide) (PEO) complexes with alkali metal salts were made by Wright et al.(38, 39). It was after Armand that the potential of these new materials were realized for future battery applications(40). PEO with high molecular weight of about 5 x 106 and 80 % crystallinity was usually employed as the polymer host to form complexes with lithium salts. Apart from the ability of the sequential oxyethylene group, -CH2-CH2-O- in PEO to complex with lithium salts, polymer hosts containing sequential polar groups such as O-, -NH- and -C-N- in the polymer chain were also found to dissolve lithium salts(41).
 The lithium salt complexed-poly(ethylene oxide) (PEO)(42) and -poly(propylene oxide) (PPO)(43) are the most widely investigated “dry solid polymer” electrolyte systems in all solid state lithium batteries. The main reason to choose these two polymer hosts is because they form more stable complexes and possess higher ionic conductivities than any other group of solvating polymers without the addition of organic solvents. Complex formation in PEOn, -salt (n = number of ether oxygens per mole of salt) is governed by competition between solvation energy and lattice energy of the polymer and the inorganic salt(44). Low lattice energies of both the polymer and the complexing salt have been found to increase stabilities in the resultant polymer electrolyte(44).
 On dry solid polymer systems, to enhance the ionic conductivity, the improvement of host polymers was attempted by the reduction of its crystallinity.(7, 45-50) The performances of electrochemical cells incorporating the “dry solid” polymer electrolytes and lithium metal electrodes were not satisfactory, and cycle lives were as low as 200 to 300 cycles. The poor performance of the cells was mainly attributed to the poor conductivity of the electrolytes, along with the reactivity of the anion of the electrolyte towards the lithium metal electrodes. Therefore, the polymer gel systems, whose crystallinity is reduced by adding plasticizer, have attracted attention.
 In 1975, Feuillade and Perche demonstrated the idea of plasticizing a polymer with an aprotic solution containing an alkali metal salt in which the organic solution of the alkali metal salt remained trapped within the matrix of the polymer(51). Such mixings have resulted into formation of gels with ionic conductivities close to those of the liquid electrolytes, arising from similar conductivity mechanisms taking place in both the systems. However, one would expect slightly lower conductivities in the polymeric gels, if more viscous solvents are used, as compared to those used in the liquid electrolytes. Since then, various polymeric hosts such as, poly-(vinylidene fluoride) (PVdF)(52), poly(vinylidene carbonate) (PVdC), poly(acrylonitrile) (PAN)(53, 54), poly(vinyl chloride) (PVC)(55), poly(vinyl sulfone) (PVS)(56), poly(p-phenylene terepthalamide) (PPTA)(57), and poly(vinyl pyrrolidone) (PVP), have been found to form electrolytes with conductivities ranging between 10-4 and 10-3 S/cm at 20 oC (Table 1.2.1)(7). These systems are presently expressed in various terms such as “plasticized polymer electrolytes”, “polymer hybrids”, “gelionics” and “gel electrolytes”. The electrolytes are easily prepared by heating a mixture containing the appropriate amounts of the polymer, solvents and a lithium salt to about 120 - 150 oC, a temperature above the glass transition temperature of the polymer, to form viscous clear liquids. Films of the gels are usually made by solution casting in hot state allowing the solution to cool under pressure of electrodes. Less-evaporating solvents, such as ethylene carbonate (EC), propylene carbonate (PC), dimethyl formamide (DMF), diethyl phthalate (DEP), di-ethyl carbonate (DEC), methylethyl carbonate (MEC), dimethyl carbonate (DMC), γ-butyrolactone (BL), glycol sulfite (GS), and alkyl phthalates have been commonly investigated as “plasticizer” solvents for the gel electrolytes(7). The solvents have been used separately or as mixtures. The phenomenon of plasticization was found to increase the amorphous phase in the polymer system with a single glass transition temperature as low as -40 oC, the latter varying with the amounts of solvent and polymer containing in the gels. It was generally observed that up to 80 % of solvent could be trapped into the polymer matrix. The high permittivity solvents allowed a greater dissociation of the lithium salt and increased the mobility of the cation. Much research work in the area of gel electrolytes can be found in literature reviews(7, 44, 58). 
PEO-based gels  
 Plasticization of high molecular weight P(EO)n-LiX electrolytes with PC and/or EC was found to form soft solids with poor mechanical stabilities, although room temperature conductivities as high as the order of 10-3 S/cm were obtained(59, 60). The poor mechanical stability was accounted to be mainly due to the solubility of the PEO in the solvents(60). Cross-linking of polymers by methods such as, UV(61), thermal radiation(62), photo-polymerization(63), and electron beam radiation polymerization(64) was found to reduce the solubility of the polymers with the organic solvents and also helped to trap the liquid electrolyte within the polymer matrix. Low molecular weight PEO was cross-linked and plasticized with 50 % PC by Borghini et al.(65). This material showed good mechanical proper-ties and conductivities of the order 10-4 S/cm at 20 oC were two orders of magnitude higher than the unplasticized amorphous cross-linked PEO-LiClO4 complex. In general, the room temperature conductivities of gels based on polymers and copolymers prepared by crosslinking methods were found to be in the range of 10-5 - 10-4 S/cm(65). Cross-linking of polymers by methods such as, UV(61), thermal radiation(62), photo-polymerization(63), and electron beam radiation polymerization(64) was found to reduce the solubility of the polymers with the organic solvents and also helped to trap the liquid electrolyte within the polymer matrix. Low molecular weight PEO was cross-linked and plasticized with 50 % PC by Borghini et al.(65). This material showed good mechanical proper-ties and conductivities of the order 10-4 S/cm at 20 oC were two orders of magnitude higher than the unplasticized amorphous cross-linked PEO-LiClO4 complex. In general, the room temperature conductivities of gels based on polymers and copolymers prepared by crosslinking methods were found to be in the range of 10-5 - 10-4 S/cm(65).
PAN- and PVDF-based gels The PAN-based(67-69) and PVdF-based(70) gel electrolytes are the most widely investigated polymer gel electrolyte systems and remarkable conductivities of the order of 10-3 S/cm at 20 oC have been obtained. A fully amorphous gel of PAN-LiClO4 (1 : 0.2) in EC showed room temperature conductivity of 1 x 10-3 S/cm and an estimated activation energy of 86 kJ/mol(71). Because of the absence of any oxygen atoms in the PAN polymer matrix, the PAN-based gels were found to have lithium ion transference numbers greater than 0.5(72). With a greater dissociation of salts such as, LiTFSI and LiTFSM in PAN-based gels, transference numbers as high as 0.7 could be obtained. The delocalization of electron density around the large organic anions in the salts is believed to promote greater dissociation and increase the transference number. Strong polar groups in the polymer are undesirable because of the complexation of lithium ion by both the polymer matrix and the solvent(73). A probable interaction of the lithium ion with the polymer chain and the solvent is shown in Fig. 1.2.1 for a PAN-based gel system(74). Thus, the polymer in the gel electrolytes would mainly act as an encapsulating matrix with electrostatic interactions with the solvated lithium salt so that the lithium ion mobility is least hindered. Yang et al.(71) studied a PAN-LiClO4-DMF gel in order to understand the interaction of the Li+ ions with the PAN chain. Using FTIR they showed the Li+ ions to form bonds with C=N groups of the PAN as well as C=O of the DMF which would support the structure in Fig. 1.2.1. Using the same technique, in PAN-LiClO4-DMF, Wang et al. reported interactions between Li+ ions and the oxygen and/or nitrogen atoms of DMF along with interactions between the O atoms of DMF and N atoms of the nitrile of PAN(75). Unfortunately, no interactions between LiClO4 and PAN could be detected with the use of the FTIR technique. Jiang et al. recently studied PVdF-based gels consisting of EC, PC and a LiX salt(76). As expected they found the presence of plasticizers EC and PC to significantly disorder the crystalline structure and reduce crystallinity in the gels than in the parent polymer host, PVdF. They also reported the mechanical strength of the resulting gel to be depended on the PVdF content, whereas the conductivity to be mainly influenced by the viscosity of the medium and the concentration of the lithium salt. Room temperature conductivities as high as 2.2 x 10-3 S/cm for the LiN(CF3SO2)2 salt containing mixture were reported(76). Increase in gel properties were observed when mixtures of polymer hosts or copolymers were used in the gel preparation. Increase in gel properties were observed for PVdF co-polymerized with hexafluoropropylene (HFP). The PVdF-HFP co-polymer in the gel showed greater solubility for organic solvents, and lower crystallinity and glass transition temperature than the PVdF polymer alone in the gel(77).Others Wieczorek and Stevens studied blends polyether, PMMA, and LiCF3SO3(78). The electrolytes showed a maximum room temperature conductivity of 3 x 10-5 S/cm. Morita et al. also prepared gels consisting of PMMA grafted with PEO and containing Li salts, and room temperature conductivities of the order 10-3 were reported(79). Addition of crown ethers such as 12-crown-4 and 15-crown-5 to the PMMA-PEO-Li+ gel was found to increase the conductivity to some extent, whereas the Li ion mobility was found to increase significantly with the addition of the 15-crown-5 ether(79). Lithium transference numbers greater than 0.5 were found in the gel electrolytes based on poly(methyl methacrylate)(80) and poly(tetrahydrofuran)(81). The values were found to be dependent on the amounts of organic solvent present in the gel. Blends of polymers, PVC and PMMA were studied using PC as the plasticizer and LiCF3SO3 as the salt(82). Due to insolubility of PVC in the solvent PC, phase separation was observed. The inclusion of PVC into PMMA helped to increase the mechanical stability of the gel, but unfortunately decreased the lithium ion conductivity. The ions were found to preferentially move towards the plasticizer-rich phase or the PMMA-rich phase. Dimensionally stable gels consisting of PEG-PAN-PC-EC-LiClO4 were prepared by Munichandraiah et al.(83). Compared with gels PEO-PC-LiClO4 and PAN-PC-EC-LiClO4, the PEG containing gels showed lower room temperature conductivities, but higher mechanical stabilities. PVC-based electrolytes consisting of LiClO4, LiTFSM, or LPF6 salt and THF/PC mixture as plasticizer gave ionic conductivity of the order 10-4 S/cm at 20 oC with respective lithium ion transference numbers of 0.26, 0.40, and 0.45(84). Polymer gels consisting of PVC, LiTFSI, and solvents such as dibutyl phthalate (DBP) and dioctyl adipate (DOP) gave room temperature conductivities of the order 10-4 S/cm(85).
 Low molecular ethers such as poly(ethylene glycol) dimethyl ether (PEGDME) was used as a solvent for salts such as LiN(CF3SO2)2, LiCF3SO3, and LiPF6 to which poly(vinylidene fluoride)-hexafluoropropene copolymer was added as the encapsulating polymer matrix(86). PEGDME was also added to a copolymer formed from triethylene glycol dimethacrylate (TRGDMA) and acrylonitrile (AN) and LiCF3SO3 salt. Room temperature conductivities in the range from 10-5 to 10-4 S/cm were reported(87). The conductivity was found to increase in line with the molar ratios of both AN : TRGDMA and (EO) : LiCF3SO3(87). In the gel mixture consisting of poly(p-phenylene terephthalamide) (PPTA), polyethylene glycol (PEG), polycarbonate in PC-EC, and a lithium salt, conductivities as high as 2.2 x 10-3 S/cm at room temperature were observed at 0.8 M LiBF4 salt per mole of PPTA content(57). Above 1 M of the salt content, the conductivity was found to fall rapidly suggesting the ion conductivity to be due to LiBF4 interaction at the amide bond sites of the PPTA(57).
 In the polymer gel electrolytes, the incorporation of liquid electrolytes into homo- or co-polymer hosts has allowed room temperature conductivities as high as 10-3 S/cm. The mechanical stability of the gels is determined by the ratios of the polymer and solvent (plasticizer) in the gel. Highly mechanically stable gels with 10 - 12 wt.% of plasticizer and 70-80 wt.% of polymer, with room temperature conductivities between 10-4 and 10-3 S/cm have been generally observed. Plasticizer solvents such as EC and PC have been much studied due to their high polarity and low vapor pressure, which also allow greater plasticizing effect to the polymer host. Most of the studied plasticizers have shown electrochemical instabilities at lithium metal electrode surfaces, and thus a search for new plasticizers is seen necessary. Due to greater dissociation of the lithium salt, much higher lithium ion transference numbers than in the solvent-free electrolytes (about 0.6) have been observed.
コピペ表 ：Table 1.2.1 Ionic conductivities of some gel polymer electrolytes and polymer composite electrolytes.　は、Felix B. Dias氏の論文のTable.3のコピペ







コピペ図 ：Figure 1.2.1 Probable interaction of the lithium ion in gel electrolytes.　は、Felix B. Dias氏の論文のFig.6のコピペ








奈良洋希氏の博士論文のchapter 1 の 1.3 Anode materials の、1.3.1から1.3.2の途中までの文章は、 Martin Winter氏の論文（Issue Advanced Materials Advanced Materials Volume 10, Issue 10, pages 725–763, July, 1998）
http://wenku.baidu.com/view/7499ca59be23482fb4da4c41のコピペです（コピペ元の文献を引用してあるものの、長文の不適切な引用（数パラグラフ丸ごとコピペ）です。）

同一文章４ ：黄色でハイライトされた部分がMartin Winter氏の論文と同一文章。
1.3. Anode materials(88, 89) 
1.3.1. Carbon anodes Since the lithium ion secondary battery with carbon anode hadcommercialized by Sony Corporation in 1991(6), the carbon anode has beeninvestigated. At present mostly carbons are used as the negative electrode ofcommercial rechargeable lithium batteries: i) because they exhibit both higherspecific charges and more negative redox potentials than most metal oxides,chalcogenides, and polymers; and ii) due to their dimensional stability, they showbetter cycling performance than Li alloys. The insertion of lithium into carbon,habitually named “intercalation”, proceeds according to Equation (1).
　（式　略） Due to electrochemical reduction (charge) of the carbon host, lithium ionsfrom the electrolyte penetrate into the carbon and form a lithium/carbonintercalation compound, LixCn. The reaction is reversible. The quality of sites capable of lithium accommodation strongly depends onthe crystallinity, the microstructure, and the micromorphology of the carbonaceousmaterial(90-107). Thus, the kind of carbon determines the current/potentialcharacteristics of the electrochemical intercalation reaction, and also the risk ofsolvent co-intercalation. Carbonaceous materials suitable for lithium intercalation are commerciallyavailable in hundreds of types and qualities(90, 91, 108-111). Many exotic carbons havebeen synthesized in laboratories by pyrolysis of various precursors, some of themwith a remarkably high specific charge. Fullerenes have also been tested(112, 113).Carbons that are capable of reversible lithium intercalation can roughly beclassified as graphitic and non-graphitic (disordered). Graphitic carbons arecarbonaceous materials with a layered structure but typically with a number ofstructural defects. From a crystallographic point of view the term “graphite” isonly applicable to carbons having a layered lattice structure with a perfect stackingorder of graphene layers, either the prevalent AB (hexagonal graphite, Fig. 1.3.1) orthe less common ABC (rhombohedral graphite). Due to the small transformationenergy of AB into ABC stacking (and vice versa), perfectly stacked graphite crystalsare not readily available. Therefore, the term “graphite” is often used regardless ofstacking order. The actual structure of carbonaceous materials typically deviates more or less from the ideal graphite structure. Materials consisting of aggregates ofgraphite crystallites are named graphite as well. For instance, the terms naturalgraphite, artificial or synthetic graphite, and pyrolytic graphite are commonly used,although the materials are polycrystalline. The crystallites may vary considerablyin size. In some carbons, the aggregates are large and relatively free of defects, forexample, in highly oriented pyrolytic graphite (HOPG). In addition to graphiticcrystallites, other carbons also include crystallites containing carbon layers (orpackages of stacked carbon layers) having significant misfits and misorientationangles of the stacked segments to each other (turbostratic orientation orturbostratic disorder)(114). The latter disorder can be identified from an increasedaverage planar spacing compared to graphite(115). Non-graphitic carbonaceous materials consist of carbon atoms that aremainly arranged in a planar hexagonal network but without far-reachingcrystallographic order in the c-direction. The structure of those carbons ischaracterized by amorphous areas embedding and crosslinking more graphiticones(116) (Fig. 1.3.2). The number and the size of the areas vary, and depend onboth the precursor material and the manufacturing temperature. Non-graphiticcarbons are mostly prepared by pyrolysis of organic polymer or hydrocarbonprecursors at temperatures below ~1500 oC. Heat treatment of most non-graphitic(disordered) carbons at temperatures from ~1500 to ~3000 oC allows one to distinguish between two different carbon types. Graphitizing carbons develop thegraphite structure continuously during the heating process, as crosslinking betweenthe carbon layers is weak and, therefore, the layers are mobile enough to formgraphite-like crystallites. Non-graphitizing carbons show no true development ofthe graphite structure even at high temperatures (2500-3000 oC), since the carbonlayers are immobilized by crosslinking. Since non-graphitizing carbons aremechanically harder than graphitizing ones, it is common to divide thenon-graphitic carbons into “soft” and “hard” carbons(116).
1.3.2. Alloy anodes Various insertion materials have been proposed for negative electrodes ofrechargeable lithium batteries, for example, transition-metal oxides andchalcogenides, carbons, lithium alloys, and polymers. Table 1.3.1 shows that boththe specific charges and the charge densities of lithium insertion materials aretheoretically lower than that of metallic lithium. However, considering that thecycling efficiency of metallic lithium is ≤ 99 %, one has to employ a large excess oflithium(91, 117, 118) to reach sufficient cycle life. The practical charge density of asecondary lithium electrode is therefore much lower than the theoretical one, sothat it is comparable with the charge densities of alternative lithium-containingcompounds. However, the potential of the electrode materials also has to beconsidered because a higher potential versus Li/Li+ of the negative electrode meansa lower cell voltage. For instance, the potential of many Li alloys is ~0.3 to ~1.0 V vs.Li/Li+ whereas it is only ~0.1 V vs. Li/Li+ for graphite (Fig. 1.3.3). The replacement of metallic lithium by lithium alloys has been under investigationsince Dey(122) demonstrated the feasibility of electrochemical formation of lithiumalloys in liquid organic electrolytes in 1971. The reaction usually proceedsreversibly according to the general scheme shown in Equation (2).　（式　略）With only a few exceptions (such as hard metals, M = Ti, Ni, Mo, Nb), Li alloys areformed at ambient temperature by polarizing the metal M, for example, Al, Si, Sn,Pb, In, Bi, Sb, Ag, and some multinary alloys,(93, 120, 123-134) sufficiently negatively in aLi+- containing electrolyte. In most cases even the binary systems Li-M are verycomplex. Sequences of stoichiometric intermetallic compounds and phases LixMwith considerable phase range are usually formed during lithiation of the metal M,characterized by several steps and/or slopes in the charge diagram (Fig. 1.3.3). Theformation of Li-M phases is in many cases reversible, so that subsequent steps andslopes can also be observed during discharge. A Li+ ion transfer cell with the trademark Station, announced recently byFujifilm Celltech Co., Ltd.,(136) uses an “amorphous tin-based composite oxide(abbreviated TCO or ATOC)” for the negative electrode. The TCO combines both: i) apromising cycle life and ii) a high specific charge (>600 Ah/kg) and charge density(>2200 Ah/L). (137)
 The TCO is synthesized from SnO, B2O3, Sn2P2O7, Al2O3, and otherprecursors. However, only the SnII compounds in the composite oxide are said toform the electrochemically active centers for Li insertion. The oxides of B, P, or Al,which are electrochemically inactive, have glass forming properties and form anetwork stabilizing the dimensional integrity of the composite host duringsynthesis.(137) The improvement of cycle life by using composite amorphous lithiuminsertion materials, such as V2O5, together with a network former, such as P2O5, hasbeen reported in the literature.(138-144)
 In order to explain the high specific charge a mechanism can be suggestedin which the tin oxide reacts to form Li2O and metallic Sn.(145, 146) This reaction isassociated with large charge losses due to the irreversible formation of Li2O. In asecond step the Sn then alloys with lithium reversibly. On the other hand, accordingto Fujifilm Celltech(137) no Li2O was found after lithium insertion. However, the ideathat the high specific charge of the TCO is due to the alloying of metallic tin has ledto a renaissance of Li-alloy negative electrode research and development.(145)

コピペ表 ：Table 1.3.1 Characteristics of representative negative electrode materials for lithium batteries, calculated by using data from(117, 119-121). The values are for fully lithiated host materials except for the values in parentheses, which are for lithium-free host materials. Li4 denotes a four-fold lithium excess, which is necessary to reach a sufficient cycle life.　は、Martin Winter氏の論文のTable.1のコピペ 
コピペ図 ：Figure 1.3.2 Schematic drawing of a non-graphitic (disordered) carbonaceousmaterial.　は、Martin Winter氏の論文のFig.5のコピペ 
コピペ図 ：Figure 1.3.3 Charging curves of some matrix metals (M), compared to highlyoriented turbostratic mesophase pitch carbon fibers (P 100, FMI Composites-UnionCarbide, characterized in LiClO4 / propylene carbonate). Modified and redrawnfrom(135).　は、Martin Winter氏の論文のFig.3のコピペ

奈良洋希氏の博士論文のchapter 1 の 1.3 Anode materials の1.3.2の途中からの下記文章は、 Martin Winter氏の論文（Electrochimica Acta Volume 45, Issues 1–2, 30 September 1999, Pages 31–50）
http://wenku.baidu.com/view/a5c2261ea8114431b90dd8d6.htmlのコピペです（コピペ元の文献を引用してあるものの、長文の不適切な引用（数パラグラフ丸ごとコピペ）です。）

このMartin Winter氏の論文の購読には、本来$35.95かかりますが、奈良氏の論文にコピペされているので、無料で読める状態になっています。
同一文章５ ：黄色でハイライトされた部分がMartin Winter氏の論文（Electrochimica Acta）と同一文章。
 The electrochemical reduction of a tin electrode in a Li+-containingelectrolyte leads to the subsequent formation of a number of intermetallic phasesLixSny at high temperature (415 oC)(147) as well as at room temperature(145, 148-150).Diffusion data (Table 1.3.2) disclose that the Li+ cation mobility in lithium-richphases is quite acceptable both at 415 oC(148) and at room temperature(123, 132)allowing reasonable charge/discharge current densities. However, whereas veryslow measurements at near equilibrium conditions allow to characterize severalLixSny phases (particular at 415 oC, cf. data in Table 1.3.2), at room temperatureand under practical charging conditions only the lithium-poor phases Li2Sn5 andLiSn can be clearly distinguished in the constant current charge curve (Figure1.3.4) as well as in X-ray diffraction patterns.(145, 151) It has been suggested that thelithium-rich phases LixSny do not form long-range ordered structures, because theatom mobility is too low at room temperature in this case.(145, 151)
 The lithium-rich phases exhibit high theoretical specific charges and chargedensities (Table 1.3.2). On the other hand, the tin pulverizes rapidly when itexperiences full lithiation due to the large density decrease (is volume increase).Limiting the degree of lithium uptake to the stoichiometry of LiSn relieves the pulverization problem to some extent, because the differences in density betweenmetallic Sn (7.29 g/cm3) and the phases Li2Sn5 and LiSn are relatively small (Table1.3.2). Moreover, these phases are structurally related(152), so that the consequencesof reconstitution reactions are small as well. However, further alloying with lithiumbeyond the stoichiometry LiSn leads to a rapid density decrease (Table 1.3.2) andcauses major structural rearrangements, which considerably enlarge themechanical stresses on the host material. In order to improve the dimensionalstability and thus the rechargeability for higher lithium uptakes, the morphologyand chemistry of the host material has to be specifically designed.
 Host metal properties, such as particle size, shape, texture, porosity, etc.strongly affect the macroscopic dimensional stability during lithium alloying anddealloying and thus the cycling behavior.(135, 153, 154) Though the volume expansions ofthe metal hosts upon alloying with lithium are in the order of several 100 %, largeabsolute volume changes can be avoided, when the size of the metallic host particlesis kept small (Fig. 1.3.5).
 The practical feasibility of this concept was checked by employing tinanodes of different particle sizes. These have been prepared by electroplating oncopper substrates from aqueous solutions containing Sn2+ cations. The morphologyof the metals can be drastically changed by a control of the plating conditions, forexample, by variation (i) of the chemical composition of the solutions (e.g. byaddition of leveling and complexing agents), (ii) of the solution temperature, (iii) ofthe stirring conditions and (iv) of the plating current densities. The detailedpreparation procedures can be found elsewhere.(135, 154) The thickness of thedeposits is in the range of a few micrometers, i.e. too thin for practical application incurrent cylindrical and prismatic lithium ion cells. Nevertheless, since they arebinder-free they are good model substrates for understanding the relation betweenmetal anode properties and electrochemical performance.
 The cycling performance of lithium alloys can be significantly improved ifintermetallic and/or composite hosts are employed instead of pure metals. Thebasic idea is that at a certain stage of charge/discharge, i.e. at a certain electrodepotential, one (or more) components/phases of the composite or intermetallic areable to store lithium (`reactant'), i.e. expand/contract, whereas the othercomponents/phases are less active or even inactive(88), i.e. perform as a `matrix'buffering the expansions of the reactant (Fig. 1.3.6). Additional preconditions for asuccessful application of this concept are (i) that the reactant is finely dispersed inthe matrix, (ii) that the matrix allows the electrons and lithium ions to movebetween the reactive domains or is a mixed (electron and lithium cation) conductorand (iii) that the reactive domains are sufficiently small, so that the above discussedbenefits of `small particle size' are effective.

コピペ表 ：Table 1.3.2 Melting points, densities from X-ray data dx), specific charges, charge densities as well as plateau potentials and maximum chemical diffusion coefficients ((Dchem-max) of several LixSny phases at 415 oC (LiCl-KCl eutectic melt as electrolyte) and 25 oC (LiAsF6/PC or LiClO4/PC as electrolytes).　は、Martin Winter氏の論文（Electrochimica Acta Volume 45, Issues 1–2, 30 September 1999, Pages 31–50）のTable. 1のコピペ

コピペ図 ：Figure 1.3.4 Characteristic charge curve of electroplated Sn in 1 M LiClO4/PC,i =0.025 mA/cm2.　は、Martin Winter氏の論文（Electrochimica Acta Volume 45, Issues 1–2, 30 September 1999, Pages 31–50）のFig.2のコピペ

コピペ図 ：Figure 1.3.5 Model: 1st lithiation (is 1st alloying with lithium) of a loosely packed small particle size metallic material. Even 100% volume expansion of the individual particles does not crack them as their absolute changes in dimensions are still small.　は、Martin Winter氏の論文（Electrochimica Acta Volume 45, Issues 1–2, 30 September 1999, Pages 31–50）のFig.3のコピペ

コピペ図 ：Figure 1.3.6 Model: strong expansions of the `reactant' domains due to lithiation can be buffered by the inactive or less active `matrix' domains, thus keeping the extent of crack formation in the overall multiphase material small.は、Martin Winter氏の論文（Electrochimica Acta Volume 45, Issues 1–2, 30 September 1999, Pages 31–50）のFig.7のコピペ

奈良洋希氏の博士論文のchapter 1 の 1.3 Anode materials の1.3.2の途中からの下記文章の一部は、コピペです。　Ahn氏の論文（Mater. Trans. 43, 63 2002）。
同一文章６ ：黄色でハイライトされた部分がと同一文章。 
 Thus, the selection of an adequate matrix is the key to a successful Sncompound anode. Elements that are inactive against Li are assumed to suppressthe volume change effectively without much irreversible capacity. The use of Sn compounds with elements such as Ni(155-157), Fe(158-162), Cu(163-166), Mn(160, 167), andCo(160) has been investigated based on this assumption. Recently, an amorphousternary Sn-Co-C anode has been introduced to a practical use.(168)

奈良洋希氏の博士論文のchapter 1 の 1.4 の下記文章の一部や図は、R. C. Willemse氏の論文「Polymer, 40, 827 (1999) 827-834」からのコピペです。
コピペ図 ：Figure 1.4.1.Phase inversion as a function of the viscosity ratio, p = ηd/ηm, according to several empirical relations summarized. * and ●: points found for PE/PS systems in Ref. (170).は、R. C. Willemse氏の論文「Polymer, 40, 827 (1999) 827-834」のFig.1のコピペ 
同一文章７ ：黄色でハイライトされた部分がR. C. Willemse氏の論文と同一文章。
1.4. Phase separated structure
1.4.1. Polymer blend
 A kind and quantity of the polymer materials were rapidly developed in these dozens of years. We can not live without the polymer materials over a wide area. However, the polymer materials are said to be reached the age of puberty, the characteristic of polymer are demanded to meet a broad applications. However, the performance as the single material has already come to the limit. The composite, copolymerization, and polymer blend of polymers are needed as the metallic and inorganic materials had done. The multicomponent polymer will satisfy enough demanded performance by a suitable combination from many existing polymers.
Polymer blend composed of two immiscible polymers(169, 170)A co-continuous morphology is a non-equilibrium morphology that is generated during mixing of two polymers. As such, it is an unstable morphology, and it starts changing through filament break-up and retraction as soon as the fluid blend comes out of the mixer. However, the blend may remain co-continuous, if it is frozen fast. Considerable attention has been given to the conditions that make co-continuous morphologies possible in blends during the mixing. It has been generally suggested that co-continuity occurs at the phase inversion point. Existing empirical relations(171-173) and theories(174) give a volume fraction for phase inversion as a function of the viscosity ratio, as shown in Fig. 1.4.1. However, many experimental results(170) cannot be described with these relations (Fig. 1.4.1). By basing the phase inversion point on the viscosity ratio of the components only, these relations neglect the influence of the blending conditions and material properties, e.g. the interfacial tension. Moreover, they do not take into account any requirements as to the shape of the dispersed component necessary to obtain co-continuity. Especially at low volume fractions, a co-continuous network can only exist if the minor blend component consists of structures with an extended shape(170). These structures can be formed and remain so only under appropriate blending conditions. For this reason, it is to be expected that the existence of a co-continuous morphology in the blend will be strongly dependent both on the processing conditions, e.g. the stress levels, and the processing properties of the blend components, e.g. the matrix viscosity, viscosity ratio and interfacial tension. In the literature(170), an equation has been derived that describes the critical volume fraction of the minor phase for complete co-continuity as a function of the matrix viscosity, interfacial tension, shear rate and phase dimensions, and it was shown that a high viscosity of the matrix component was favorable for co-continuity over a broad composition range in blends of commercial grades of polyethylene and polystyrene.

奈良洋希氏の博士論文のchapter 1 の 1.4 の下記文章の一部は、Park, C氏らの論文「Polymer, Volume 44, Number 22, October 2003 , pp. 6725-6760(36)」からのコピペです。
同一文章８ ： 黄色でハイライトされた部分がPark, C氏らの論文と同一文章。
Polymer blend composed of a di-block copolymer (175) 
Block copolymers have recently received much attention not only thanks to the scale of the microdomains (tens of nanometers), their various chemical and physical properties but also due to the convenient size and shape tunability of microdomains afforded by simply changing their molecular weights and compositions. Many potential uses of block copolymers for different nanotechnologies have been proposed based on principally their ability to form interesting patterns. However, the main challenge of using block copolymers lies with control of microstructure. Achievement of precise microdomain location, orientation, and elimination of various defects requires introduction of external fields during the processing step. A variety of mechanical, electrical, magnetic biases and surface interactions have been proposed to manipulate and guide the microstructures of block copolymers.
Block copolymers consist of chemically distinct polymer chains covalently linked to form a single molecule. Owing to their mutual repulsion, dissimilar blocks tend to segregate into different domains, the spatial extent of the domains being limited by the constraint imposed by the chemical connectivity of the blocks. Area minimization at the interface (the IMDS) of two blocks takes place to lower the interfacial energy. From an entropic standpoint, the molecules prefer random coil shapes but the blocks are stretched away from the IMDS to avoid unfavorable contacts. As a result, of these competing effects, self-organized periodic microstructures emerge on the nanoscopic length scale. Various microdomain structures are achieved, depending on relative volume ratio between blocks and chain architecture as well as the persistence lengths of the respective blocks.
 In the simplest case of non-crystalline flexible coil AB di-block copolymers, the composition of the AB di-block (i.e. the volume fraction f of block A) controls the geometry of the microdomain structure. As shown in Fig. 1.4.2, for nearly symmetric di-blocks (f ~ 1/2) a lamellar phase occurs. For moderate compositional asymmetries, a complex bi-continuous state, known as the double gyroid phase, has been observed in which the minority blocks form domains consisting of two interweaving threefoldcoordinated networks. At yet higher compositional asymmetry, the minority component forms hexagonally packed cylinders and then spheres arranged on a body-centered cubic lattice. Eventually, as f → 0 or 1, a homogeneous phase results.(176)

奈良洋希氏の博士論文のchapter 1 の 1.4 の下記文章の一部は、D. John Mitchell氏らの論文「J. Chem. Soc., Faraday Trans. 1, 1983,79, 975-1000 DOI: 10.1039/F19837900975」からのコピペです。

同一文章９ ： 黄色でハイライトされた部分がD. John Mitchell氏らの論文と同一文章。
1.4.2. Liquid crystals(178) Polyoxyethylene surfactants are widely used as emulsifying agents and detergents. Like ionic surfactants they form micelles above a critical concentration in water (the c.m.c.) with liquid crystals frequently occurring at higher concentrations. They are unusual in having a lower consolute temperature, termed the cloud point. For many years this was attributed to the presence of giant micelles. However, it is now thought to correspond to the phase separation of a concentrated surfactant solution containing small micelles from a more dilute solution, perhaps due to the presence of significant intermicellar attractions.(179-181) Some surfactants even show a ‘double cloud point’, where a 1 % aqueous solution goes cloudy then clears and clouds a second time when the temperature is increased.(182)
 Over years a number of theories to describe micelle shapes have been published.(183, 184) These involve a balance of alkyl-chain/water repulsions and repulsion between adjacent head groups within the micelle, together with surface curvature and limitations due to alkyl-chain packing. Most lyotropic liquid crystals consist of ordered micelles.(185) It should be possible to extend the theories of micelle shapes to account for mesophase formation by addition of appropriate terms for intermicellar forces. Parsegian(186) has already successfully described the hexagonal/lamellar transition of ionic surfactants using Poisson-Boltzmann theory to account for electrostatic repulsions. Wennerstram and co-workers(187, 188) have developed this approach further. For non-ionic surfactants there is no quantitative theory to describe intermicellar repulsions. However, with poly(oxyethylene) alkyl ethers [n-CnH2n+1(OCH2CH2)mOH,CnEOm] it is possible to change systematically the surfactant chemical structure by alteration of m and n, hence varying head-group interactions and micelle size. By determining water/surfactant phase diagrams of a range of compounds, the validity of this theoretical description and the importance of each contributing factor (alkyl-chain conformations, curvature, head-group area etc.) for the structure and stability of mesophases can be assessed.
 The shape of micelles just above the c.m.c. is highly dependent upon surfactant type and solution conditions (concentration, electrolyte level, temperature). It appears that spherical micelles, rod micelles and oblate spheroid (bilayer) micelles all occur, but under different circumstances. The structures of normal hexagonal (H1), lamellar (Lα) and reversed hexagonal (H2) lyotropic mesophases are well known.(185, 189) The lamellar phase consists of equidistant parallel surfactant bilayers separated by water layers. In the hexagonal phases very long normal or reversed rod micelles are packed in a hexagonal array. Two different classes of cubic phases have been discovered,(185, 190) one occurs at compositions between micellar solution (L1) and H1 phases, consisting of ordered spherical micelles packed in body- or face-centred arrays (labelled I1), and the second occurs between Lα and H1 (labelled V1) or H2 (labelled V2). While the structures of these phases are not fully known, the best candidates(185, 190) are regular ‘bicontinuous’ networks where the chain/water interface has both positive and negative curvatures.
 The forces responsible for micellar shape and mesophase structure can be divided into intramicellar and intermicellar contributions. The former determine the shape just above the c.m.c. while the latter take account of intermicelle interactions at higher concentrations. The description given below is a summary and development of several treatments.(183-185, 190, 191)
 At the c.m.c. micelle shape is determined by the surface area per molecule (a) at the alkyl-chain/water interface and the interface curvature. The possible micelle shapes are spheres, rods (prolate spheroids) or discs (oblate spheroids, bilayer micelles). For a spherical micelle, the volume of the surfactant alkyl chain (v), micelle radius (r) and a are related by
 v = ar 3 ........................................................................................................(1)
 Since r cannot extend beyond the ‘all-trans’ length of alkyl chain (lt), while v is invariant, then a cannot fall below a certain limit (asc). From the known densities of alkanes, assuming a micelle radius of 1.25 Ǻ per C-C bond this is estimated to be asc = 70 Ǻ 2. Similar arguments apply to rod micelles, with the limit (arc ) being arc = 47 Ǻ 2. Thus if a ≥ 70 Ǻ, the possible micelle shapes are spheres, rods or discs. With a in the range 70 > a/ Ǻ 2 ≥ 47, only rod or disc micelles can form. With a < 47 Ǻ 2, rod micelles are excluded and only bilayer micelles (or reversed phases) can form. This limitation of micelle type is generally referred to as the ‘alkyl-chain packing constraint’.
 The detailed shapes adopted by rod and disc micelles are not known. Prolate and oblate ellipsoids, or rods with hemi-spherical ends and bilayers with hemi-cylindrical edges, are both popular pictorial representations. It has been argued that only ellipsoids with low eccentricity can occur because alkyl chains are unable to fill up completely the highly curved regions at the edge of ellipsoids with high eccentricity.(183) This conclusion involves an assumption that the alkyl-chain axis is, on average, normal to the interface. There seems to be no physical reason why a few chains at the ellipsoid–micelle edge should not be tilted to allow packing into the highly curved region, and thus the restraint is too severe.
 Given that the value of a is sufficiently large for any particular shape to occur, the actual form adopted is determined by surface curvature. The forces present at the chain/water interface (intramicellar forces) can be represented as occurring in different planes, as shown schematically in Fig. 1.4.3. The chain/water interface is at x. Interactions between hydrated head groups (electrostatic, solvation, steric) usually result in a repulsive force in some plane y, further into the water than x. The minimization of hydrocarbon-chain/water interactions gives an attractive force in plane x. These are dominant contributions.
 However, depending on the value of a, there will also be an apparent repulsive force in some plane z within the alkyl-chain region. This arises from the unfavorable entropy change caused by restrictions on the conformations of the chains with small a values. While no calculations of the magnitude of this effect have been published, it is obvious that it will increase with temperature and with alkyl-chain length, since the number of available conformations is proportional to 3n-2 (n = chain carbon number). The effect will also depend on the shape of the micelle. Without detailed model calculations of chain conformations it is not possible to estimate this contribution to the stability of spheres, rods or bilayers at large a value (a > asc). However, any restriction on the chain taking the all-trans conformation will increase the limiting values asc, and arc. Moreover, with a > asc, the average chain length within a binary is < lt/3. This will involve numerous gauche conformations which have an unfavorable enthalpy, so disfavoring the lamellar phase. For low values of a, spheres and rods cannot pack. On comparing reversed phases and the lamellar phase as a approaches the limit of the all-trans cross-sectional area within the lamellar phase (abc), the reversed phases will be favored at sufficiently high temperatures because they have more conformational states available to the chains. Thus in practice the alkyl-chain curvature effect is likely to transform bilayers into reversed structures when a is too small for rods or spheres to occur.
 The contributions listed above can be represented by an equation of the form
（式　略）
　where μo is the energy per surfactant molecule, and C(r) is a curvature free-energy expression, being a function of micelle radius; the alkyl-chain/water repulsion is given by γa, where γ represents surface tension, and all other contributions are included in g. One could employ expressions involving higher powers of a but these do not qualitatively alter the description,(183) nor does inclusion of an allowance for the area of the chain/water interface occupied by the head groups.(184)
　To summarize, the balance of forces indicated in Fig. 1.4.3 determines the aggregate shape within packing constraints. A large repulsion in plane y gives water continuous phases while small repulsions give reversed micelles. With large repulsions spherical micelles occur, where a (sphere) < a (bilayer) [provided that a (sphere) ≥ asc]. When a is just too small for spheres, one expects rod micelles (prolate ellipsoids), and as a is reduced further, oblate spheroids (bilayers) with a < arc. The particular values of a where shape changes occur will depend on the detailed nature of the forces in planes x and y (and z). Reversed phases are expected at still smaller values of a.
　There are two major effects of intermicellar forces. First, if we consider the micelles formed at the c.m.c. as hard-core particles, then as the volume fraction of micelles is increased, we will observe order/disorder transitions. Spherical micelles will close-pack into a regular cubic array. Long rod micelles will form a hexagonal array, while a lamellar phase will form with bilayers. The limiting volume fractions are 0.74 for face-centred cubic and 0.91 for a hexagonal phase. Theoretically, a lamellar phase can occur without water being present, i.e. limiting
　volume fraction = 1. (In practice it is well known that anhydrous soaps form a lamellar phase at high temperatures.(185, 189) The actual volume fractions for order/disorder transitions appear to be in the range 0.7-0.8 of the close-packed volume values(192) (i.e. ca. 0.5 and ca. 0.7 for spheres and rods, respectively). For lamellae, two large bilayers in any volume will be aligned to some extent because they cannot pass through each other. Thus the lowest theoretical volume fraction for a lamellar phase is just above the c.m.c. However, van der Waals attractions between bilayers can be sufficiently large to overcome entropy and interbilayer repulsions (193) causing phase separation of a lamellar phase where the water layers do not swell indefinitely. (Similar considerations apply to disordered oblate micellar solutions, hence the occurrence of the cloud point.(180, 185)) Thus instead of a dilute solution of large bilayers, a lamellar-phase + water dispersion occurs just above the c.m.c. Starting with various different micellar shapes, one would expect the following phase sequences (the numbers refer to the volume fraction of micelles required for the phase transition):
　　oblate spheroid  (disc micelles) → lamellar phase
　　prolate spheroid (rod micelles)  → hexagonal phase → lamellar phases
　　pherical phase → cubic phase → hexagonal phase → lamellar phase
A hexagonal phase occurs at volume fractions above the close-packing limit of spherical micelles because the rods can relieve some of the surface strain due to curvature better than the lamellar phase. In the second and third cases a lamellar phase occurs only when the close-packing volume of rods is exceeded. In practice, the limit of mesophase formation is often determined by the stability of a crystalline surfactant phase, particularly with ionic or zwitterionic compounds. This description has nothing to say about the crystalline state.
 Of course, the repulsions between micelles are not of the hard-sphere type, but occur at a range of distances from the micelle surface. This soft-core repulsion can extend to many micelle diameters if ionic surfactants with low concentrations of added electrolyte are present. For zwitterionic surfactants (lecithins) the existence of an apparent ‘hydration’ repulsion force has been demonstrated.(193) The force can be thought of as arising from interactions between hydrogen-bonded water network structures on adjacent micelle surfaces. Theoretical considerations(194) suggest that it could have an exponential fall-off with increasing distance from the surface. For polyoxyethylene surfactants we expect that this force will be accompanied by another repulsive force due to limitations on the conformation of EO groups arising from steric hindrance when micelles are close together.
 One possible consequence of the soft-core repulsions is the formation of cubic structures other than the face-centred variety. Recent results suggest that primitive or body-centred structures can occur in addition to face-centred cubic. Thus the minimum volume fraction for cubic structures is further reduced. On increasing volume fraction, one might expect to observe the sequence: primitive cubic → body-centred cubic → face-centred cubic → other phases, with the volume fraction for these transitions being in the ratio 0.52 : 0.68 : 0.74, respectively. However, the actual structures observed could vary from this sequence according to the form of the intermicellar repulsions.
 The soft-core repulsion between micelles has a second effect: it causes the value of a to decrease with increasing volume fraction (unless a is limited by alkyl-chain packing). This leads to an increase in micelle size. More importantly, a micelle shape transition of the type sphere → rod → bilayer can occur with a > ac, where the unfavorable curvature energy is balanced by a contribution from micelle interactions. This type of transition will be accelerated when a approaches the packing limits of spheres or rods. If the repulsions between the new shapes are sufficiently large, then an ordered phase can form immediately. Thus one could observe the sequences:
 spherical micelles  (disordered) → hexagonal phase → lamellar phase
 rod micelles (desordered)→ lamellar phase
What is observed in practice will depend on the relative magnitudes of the forces and their detailed dependence on area (a) or distance (micelle separation, i.e. concentration). An illustration of the possible phase sequences at different curvatures and as a function of volume fraction is given in Fig. 1.4.4. Note that the micelle shape transitions in disordered solutions occur over a range of volume fractions (dotted lines) while transitions involving mesophases occur at constant volume fraction. Also, it was emphasized that mesophases occur from interactions between micelles. If interactions are absent then ordered phases cannot form.
 So far the effects of entropy were ignored. This is because there was no obvious model to allow for the change caused by forming small micelles from large ones which includes the effects of alkyl-chain conformations and changes in water/head-group orientations. However, any effect is likely to scale as kT/N, where N is the aggregation number of the small micelle. For a given area per molecule, N will be a function of r2, where r is the micelle radius, which in turn is proportional to the alkyl-chain length. Thus the entropy contribution will decrease rapidly with increasing alkyl-chain length.


他にも、コピペの可能性あり。


調査２：蜂巣 琢磨氏（早稲田大学の逢坂哲彌氏の研究室）の博士論文における文章のコピペについてのまとめ

 著者：　蜂巣 琢磨
論文題目：　「Chemical Synthesis of Magnetic Nanoparticles and Their Application to Magnetic Recording」
http://dspace.wul.waseda.ac.jp/dspace/handle/2065/36373 （写し）
出版日：　2011年
審査員：　
  （主査） 早稲田大学教授 工学博士（早稲田大学） 逢坂 哲彌
  （副査） 早稲田大学教授 工学博士（早稲田大学） 黒田 一幸
  （副査）  早稲田大学教授 工学博士（早稲田大学） 菅原 義之
  （副査） 早稲田大学教授 博士 (工学 )（早稲田大学） 本間 敬之
  （副査） 早稲田大学准教授 博士 （工学）（ 早稲田大学 ） 門 間 聰 之
  （副査） 学外審査員 Rome 大学教授 Ph.D (Univ. Rome) Bruno Scrosati
  （副査）  Tel Aviv 大学教授 P h . D ( Te c h n i o n ) Yo s i S h a c h a m- D i a m a n d

概要 （写し）
審査報告書 （写し）
論文本文 （写し）

蜂巣琢磨氏の博士論文のchapter 1 の、1.1 の下記文章の一部は、Junichi SAYAMA氏らの論文からのコピペです。

 同一文章１ ：黄色でハイライトされた部分がJunichi SAYAMA氏らの論文「Journal of Magnetism and Magnetic Materials Volume 287, February 2005, Pages 239–244」と同一文章。
1. Magnetic Recording Devices for Hard Disk Drives 1.1 Hard Disk Drives (HDDs) 
In the advanced information society of today, information storage technology, which helps to store a mass of electronic data and offers high-speed random access to the data, is indispensable. Compared to the amount of available storage, the amount of information created continues to increase explosively, as shown in Fig. 1.1.1 [1]. The five-year period in 2006 to 2010, the amount of information created has increased six-fold about to 1200 Exabyte (Exa; E, 1 E = 1018). In 2011, the information created is expected to swell to 1800 Ebyte. Against this background, hard disk drives (HDDs), which are magnetic recording devices, have gained in importance because of their advantages in capacity, speed, reliability, and production cost. Magnetic recording devices are one of the most important keys to the advancement of information technology since its invention more than 100 years ago by Poulsen [2]. These days, the uses of HDD extend not only to personal computers and network servers but also to consumer electronics products such as personal video recorders, portable music players, car navigation systems, video games, video cameras, and personal digital assistances. 
1.2 Estimation of Increase in HDD Areal Density and Technology Developments 
　IBM introduced the first HDD of the IBM 350 Disk File in 1956, which was the storage unit of the IBM RAMAC 305 computer [3]. It was composed of fifty 61 cm aluminum disks coated with iron oxide paint, and provided an areal density of 0.002 Mega-bit in-2 ; in the last half century, the density has been increasing to beat the b
　Fig. 1.1.2 shows the estimated increase of HDD areal density and related technology development from 1985 to 2020. Such a rapid progress of HDD areal density was attained through epochal developments of magnetic thin-film materials [4].


蜂巣琢磨氏の博士論文のchapter 1 の、1.2 の下記文章の一部は、Electrochemical Nanotechnologies Nanostructure Science and Technology 2010, pp 67-85 Magnetic Heads Tokihiko Yokoshimaからのコピペです。

同一文章２ ：黄色でハイライトされた部分が同一文章です。 
　One of the most significant innovations from the viewpoint of material improvement is the electrodeposition of permalloy (Ni80Fe20), which was introduced by IBM in 1979 as the core material of a thin-film inductive head [5]. After the introduction by IBM in 1991 of the magnetoresistive element as a read head, the areal density of HDDs jumped from 30% per year to 60% per year [6]. Following these developments, the replacement of the magnetoresistive element by a giant magnetoresistive element in the read component has led to a jump in the HDD areal density growth rate of over 100% per year since 1999, which is very impressive progress. At the beginning of 2004, the areal density of commercial HDDs approached 100 Giga-bit in-2 [7], and demonstration of 170 Gbit in-2 [8] was reported. However, there has been an apparent slowdown in this increase of density. The perpendicular magnetic recording system that was advocated by Iwasaki in 1977 [9, 10] was recently introduced to the market by Toshiba [11]. This system is a critical innovation for developing high-performance HDD systems with a high recording density. The design of the magnetic recording head also was changed because of the change in the recording system. For example, a longitudinal recording media system with a ring-type head and recording layer exhibiting in-plane anisotropy has been applied to the conventional HDD. For a perpendicular media, the sputter deposition of CoCr alloy film, which was first reported by Iwasaki and Ouchi [12], and the control of grain orientation, grain size, and magnetic properties of film also have contributed to the progress in areal density [11–14]. To meet the strong demand for high-performance write heads to be used for high-density magnetic recording, soft magnetic materials with high saturation magnetic flux density, Bs, are being developed. Although the sputtering technique has become paramount, the electrochemical technique also has played an important role in the development of magnetic recording systems [15].

蜂巣琢磨氏の博士論文のchapter 1 の、1.2 の下記文章の一部は、Felipe García Sánchez氏の博士論文の「1.2 Magnetic recording crisis and challenges」における文章や、Wikipediaの”Perpendicular recording”の文章と類似しています。

同一文章２ ：黄色でハイライトされた部分が同一文章です。

In January 2006, Seagate Technology began shipping its first laptop sized, 2.5 inch HDD using perpendicular recording technology, the Seagate Momentous 5400.3. At this time the majority of the HDD utilize the new technology. In October 2007 Seagate Technology announced a new record of magnetic recording density of 421 Gbit in-2 . The company announced the results of a magnetic recording demonstration that used perpendicular recording heads and media created with currently available production equipment. The difference between longitudinal and perpendicular recording is the orientation of the anisotropy of media grains. In 2009, 927 Gbit in-2 for a prototype discrete track media was demon- strated by TDK, and 612 Gbit in-2 for a prototype only perpendicular type was demon- strated by Hitachi GST. On the contrary, in March 2010, 541 Gbit in-2 for 2.5 inch HDD was produced by Toshiba (MK7559GSXP, 750 GB of a double 2.5 inch disk HDD type) [16]. At the end of 2010, over 700 Gbit in-2 HDD areal density of the product was achieved. 
1.3 HDD Technology Candidates  
HDD areal density has been increased because of changing memory system from longitudinal type to granular perpendicular type, i.e. a thin magnetic film. Recording demonstrations at densities of 412 Gbit in-2 [17], and 519 Gbit in-2 [18] have been reported and system designs have been discussed for Tera-bit in-2 densities [19–22].
 In order to achieve ultra-high areal density over several Tbit in-2 , we must find a reasonable compromise between signal-to-noise ratio (SNR), thermal stability, and recording writability. Since each puts different requirements on a medium, the concern about them leads to a so-called ―recording trilemma‖, as shown in Fig. 1.1.3. To get a high SNR, a small grain size is better; to achieve good thermal stability for a magnetic recording device, a large grain size or high anisotropy is desirable; to achieve good writability, it is better to use materials with a small anisotropy.

蜂巣琢磨氏の博士論文のchapter 1 の、1.3 の下記文章の一部は、http://www.fgarciasanchez.es/thesisfelipe/node5.html　（1.2 Magnetic recording crisis and challenges）や
K O'Grady, H Laidler氏らの論文「Journal of Magnetism and Magnetic Materials Volume 200, Issues 1–3, October 1999, Pages 616–63」からのコピペです。
同一文章３ ：黄色でハイライトされた部分が同一文章です。  
 Thermal stability of a bit of information date is of critical importance particularly as bits are made smaller and media are made thinner. In conventional magnetic recording, the media is a granular film and a bit consists of several (N=500/1000) almost non-interacting magnetic grains. SNR considerations are extremely complex and derive from factors such as the shape of bits and cross-talk between neighboring bits or even neighboring tracks but from simple statistical estimation the SNR is proportional to N 1/2 . Therefore, the number of grains included in a bit could not be reduced in order to preserve. SNR and, consequently, the increasing bit density imply a reduction of the grain size. However, a reduction in the grain size leads to a reduction in the energy barrier KuV, separating two magnetization states, where Ku is the magnetic anisotropy constant and V the grain volume, which determines the thermal stability of the written information. When the energy barrier is comparable to the thermal energy (shown in Fig. 1.1.4) [23], the magnetization becomes unstable and the inversion of the magnetization by thermal fluctuations is likely to occur. This effect is known as superparamagnetism and the corresponding limitation of the density as superparamagnetic limit [23]. Thermal stability equal shows in Fig.1.1.3. kB is the Boltzmann's constant and T is a room temperature. Essentially, values of thermal stability are required to ensure the long-term stability of written information.

蜂巣琢磨氏の博士論文のchapter 1 の、1.3 の下記文章の一部は、Alex Taratorin's book "Magnetic Recording Systems and Measurements" Chapter 7 "Perpendicular Recording"からのコピペです。
同一文章４ ：黄色でハイライトされた部分が同一文章です。  
 The advantages of the perpendicular recording system over the longitudinal recording are multiple: (I) Higher thermal stability can be achieved by small in-plane grain diameter with cylindrical grain structure. (II) A vertical pole head in a recording media with a soft underlayer can generate twice the field of longitudinal recording head. This allows writing higher coercivity (Hc) device, further decreasing grain size and maintaining media thermal stability. (III) The read-back signals amplitude from perpendicular media with soft underlayer (SUL) is larger compared with equivalent longitudinal media, improving signal-to-noise ratio. (iv) Perpendicular media grains are strongly oriented. This results in smaller media noise and a sharper recorded transition. (v) The demagnetization field in the perpendicular media is small at the transition region. Bit patterned media (BPM), discrete track media (DTM), energy assisted magnetic recording (EAMR) and shingled writing recording (SWR) are considered to fulfill those requirements and expected as the future magnetic recording technologies (shown in Fig. 1.1.5). The use of lithography patterned media such as DTM and BPM may enable a recording density in excess of that achievable by conventional perpendicular recording, including the potential use of exchange spring media [24]. 2055 year falls on the 100th HDD birth anniversary. They are drawing the scenario of exceeding 1 Peta-bit in-2 over by Shiroishi from Hitachi GST (Peta; P, 1 P = 1,000 T). [25] The areal memory density corresponds to the capacity of 750 TB by one media of 2.5 inch HDD when converting it into an actual product.
 Next section described about next-generation perpendicular recording technolo- gies for ultra-high recording density of several Tbit in-2 or more, based on many publications reported.

Modern Electroplating - Mordechay Schlesinger, Milan Paunovic Electrochemical Nanotechnologies - Tetsuya Osaka, Madhav Datta, Yosi Shacham-Diamand Developments in Data Storage: Materials Perspective - S. N. Piramanayagam, Tow C. Chong

調査３：本川慎二氏（早稲田大学の逢坂哲彌氏の研究室）の博士論文における文章のコピペについてのまとめ

 著者：　本川慎二
論文題目：　「ＭＥＭＳ技術を利用した新規平面型構造を有する超小型直接形メタノール燃料電池の設計と作製」
http://dspace.wul.waseda.ac.jp/dspace/handle/2065/2923
出版日：　2005年
審査員：　
  （主査） 早稲田大学教授 工学博士（早稲田大学） 逢坂 哲彌
  （副査） 早稲田大学教授 工学博士（早稲田大学） 黒田 一幸
  （副査） 早稲田大学教授 工学博士（早稲田大学） 菅原 義之
  （副査） 早稲田大学教授 博士 (工学 )（早稲田大学） 本間 敬之
  （副査） 早稲田大学教授 工学博士(東北大学) 庄子習一

概要 （写し）
審査報告書 （写し）
論文本文 （写し） （写し2）

 本川 慎二　（指導教授 逢坂 哲彌）　2005年
同一文章１ ：黄色でハイライトされた部分が同一文章です。  （Web検索で多数ＨＩＴ、出所調査中）
Chapter 1 General Introduction
Fuel Cells
 A fuel cell uses the chemical energy of hydrogen to produce electricity and water, cleanly and efficiently. Fuel cells are unique in terms of the variety of their potential applications; they can provide energy for systems as large as a utility power station and as small as a smoke detector. Fuel cells have several benefits over conventional combustion-based technologies currently used in many power plants and passenger vehicles. They produce much smaller quantities of greenhouse gases and none of the air pollutants that create smog and cause health problems. If pure hydrogen is used as a fuel, fuel cells emit only heat and water as a byproduct. A schematic of a fuel cell is shown in Fig.1.1.

Chapter 1の下記文章は、 L.J.M.J. Blomen、M.N. Mugerwa 氏の著作物「Fuel Cell Systems (1993)」のS. Srinivasan氏らによる"Overview of Fuel Cell Technology"のページ（page37～）の文書からのコピペです。
 同一文章２ ：黄色でハイライトされた部分が同一文章です。  
The thermodynamic aspects of a fuel cell are as follows [2]. The fuel cell is an electrochemical device which converts the free-energy change of an electro chemical reaction into electrical energy. One may thus write the expression.
（式　略）
where, °E is the reversible potential of the r cell. The simplest and most commonly encountered fuel cell reaction is
（式　略）
The free-energy change of this reaction under standard conditions of temperature and pressure (T=25oC, PH2 = PO2 = 1 atm, H2O in liquid state) is 56.32 kcal mol-1. The number of electrons transferred in this reaction is 2. Thus, the reversible potential is 1.229 V. The variations of the reversible potential (Er) with temperature and pressure are expressed by the equations
（式　略）
　where,nΔ is the change in number of gas molecules during the reaction. The entropy change,SΔ for reaction (1.1) is -39 entropy units, while nΔ is -3/2. Thus, Eqs. (1.3) and (1.4) show that the reversible potential decreases with an increase of temperature, 
 while the behavior is opposite with and increase of pressure. At temperature above 100oC, water is produced as a vapor in the cell. The value of SΔ is considerably less when water is produced in this state than as a liquid. Thus ∂E/∂T is -0.25 mV/oC in the former case and is -0.54 mV/oC in the latter. The effect of p ressure on Er is also less when water is produced as a vapor than as a liquid (nΔ = -1/2 in the former and -3/2 in the latter case).
 Hydrogen is not a primary fuel. Thus, attempts were made in the 1960s to use primary fuels such as hydrocarbons (CH4 to C10H22) and coal in fuel cells. However, due to the high degree of irreversibility of the anodic oxidation reactions of the hydrocarbons, these attempts proved futile. Thus, these fuels were processed to produce hydrogen for low-temperature fuel cells and H2 and CO for the higher temperature fuel cells by steam-reforming reactions of the hydrocarbons and by gasification of coal. Several other types of fuels (methanol, ethanol, ammonia, and hydrazine) were also　researched in the 1960s. Practically all of these fuel cell reactions have a themodynamic reversible potential from 1.0 to 1.2V. 
Even if there were no efficiency losses in H2-O2 fuel cells (due to activation, mass transport, and ohmic overpotentials), heat would still have to be rejected from a fuel cell because SΔ is negative for reaction (1.2). Thus, the theoretical efficiency of H2-O2, fuel cells at 25oC C, based on the enthalpy change of the reaction (commonly referred to as the higher heating value by mechanical engineers), is 83%. There is only one fuel cell reaction where ΔS is positive, namely, 
（式　略） 
For this reaction, the theoretical efficiency is 137% at 150oC. The reaction 
（式　略） 
has entropy change of 0 e.u. Thus,r this reaction, the theoretical efficiency is close to unity. As a rough rule of thumb, if O nΔ is positive, the entropy change is positive (due to i ncreasing disorder), while if nΔ is negative (increasing order), SΔ is negative, and if nΔ is zero SΔ is also zero. 
The electrode kinetics aspects of a fuel cell are as follows. The vitally important role of electrode kinetics on the performance of fuel cells (particularly those operating at low and intermediate temperature, 25-200oC) is best illustrated by a typical cell potential versus current density plot as shown in Fig.1.3. Three distinct regions are illustrated in this plot. The predominant cause of the difficulties in attaining high energy efficiencies and high power densities in low- to medium-temperature fuel cells is the low electrocatalytic activity of most electrode materials for the oxygen electrode reaction. The hydrogen electrode shows a linear relationship of its half-cell potential versus current density plot from zero to the highest value of current density in the fuel cells using phosphoric acid (T= 200oC) , potassium hydroxide (T= 80oC), or a proton- conducting electrolyte (Nafion or Dow membrane, T= 85oC). This is not the case with the oxygen electrode, where a semi-exponential relation between its half-cell potential and current density is observed. Thus, at low current densities, the entire loss in the fuel cell potential from the reversible value is due to activation overpotential at the oxygen electrode.
Another problem encountered is that the reversible potential is not attained even at zero current density. This problem is again due to the oxygen electrode. The exchange current density for this reaction is so low that competing anodic reactions (for example, oxidation of the platinum electrocatalyst, corrosion of carbon, oxidation of organic impurities) play a significant role. The net results is that the open circuit potential (OCP) is a mixed potential, which is lower than the reversible potential for the H2-O2 fuel cell reaction by about 0.1 to 0.2V. Thus, even at close to zero current densities, the efficiency of a fuel cell is lower than its theoretical value by 8-16%. The cell potential (E) versus current density (i), from a current density 0 to the value at the end of the linear region, may be expressed by the relation
（式　略）
where,
（式　略）
The parameters b and i0 are the Tafel slope and exchange current density for the oxygen reduction reaction, and R accounts for the linear variation of overpotential (predominantly ohmic) with current density, which is observed in the intermediate range.
 The situation is more complex when organic fuels (hydrocarbons and alcohols) are used directly in fuel cells. The exchange current densities for these reactions are as low as or even lower than those for the oxygen reduction reaction. Thus, the low open circuit potentials and the exponential decreas e half-cell overpotentials with current densities at both electrodes account for their relatively poor performance. It is e of th worth while rationalizing the shape of the E versus I plot (Fig.1.3). by differentiating Eqs. (1.7):
（式　略）
At a low current density, the differential resistance of the cell is high because of the first term on the right side of Eqs. (1.9); consequently, there is a steep fall of cell potential with increasing current density. At higher current densities, b << R. Thus, in the intermediate range, the cell potential varies linearly with current density. At high current densities, mass transport limitations, due to the low rates of supplies of reactants to the electrocatalytic sites or of products away from these sites, predominate and cause the sudden drop of cell potential to near-zero values. Advances have been made in the fabrication of fuel cell electrodes with optimized structures; thus, mass transport limitations are rarely encountered at current densities up to a few A/cm2. In high-temperature fuel cells using molten carbonate (T= 650℃) and solid oxide (T=1000℃ electrolytes, the exchange current densities of the fuel cell reaction are quite high (>1mA/cm2), and thus the cell potential versus current density plot is linear throughout the entire current density range (0 to a few hundred mA/cm2). The typical several cells potential versus current density plot as shown in Fig.1.4
　Theoretical and experimental electrode kinetic studies of fuel cell reactions (1.1)-(1.6) have led to the engineering design, development, and demonstrations of fuel cell power plants exhibiting high levels of performances (high energy efficiency, high power density, and reduction in noble metal loading. The major accomplishments are (i) The design and fabrication of porous gas diffusion electrodes with optimized structures to enhance (a) diffusion of dissolved gases (H2, O2) to reaction sites, (b) electrochemically active sites, and (c) ionic transport through porous electrodes; (ii) use of supported electrocatalysis (Pt crystallites on high-surface-area carbon) to significantly reduce noble metal loadings, as compared with loadings when unsupported platinum electrocatalysts were used; (iii) inhibition of CO poisoning in phosphoric acid fuel cells by elevation of operating temperature; (iv) utilization of alloys and heat-treated metal-organic macrocyclics as electrocatalysts which exhibit higher exchange non-noble metal electrocatalysts for high-temperature fuel cells with molten carbonate or solid oxide electrolytes; and (vi) use of thin electrolyte layers to minimize ohmic over potentials.
 The efficiency (ε) of a fuel cell varies with current density in the same manner as that of the cell potential with current density (cf. Fig. 1.3) because
（式　略）
The power density of a fuel cell is expressed by the relation
（式　略）
The shape of P versus i or P versus E plot is a parabola if the E versus i relation is linear. The parabola is distorted for low- and intermediate-temperature fuel cells because of the semi logarithmic relation between cell potential and current density at low values of i and sudden drops of cell potential at high current densities.
 The irreversible losses in fuel cells load to waste heat generation (Q), which may be expressed by
（式　略）
 In this equation the first term represents the entropic loss, which cannot be overcome due to thermodynamic considerations, the second term is due to activation and mass transport overpotentials, and the third is due to ohmic heating effects.
Due to the heat generation in fuel cells, it is necessary to incorporate cooling subsystems in the fuel cell system. A considerable portion of this heat is efficiently used in one ore more of the following ways: fuel processing, which usually involves an endothermic reaction; space heating; hot water; and occasionally in high temperature fuel cells, for chemical processes and for enhancing electricity generation using gas turbines.
 The performance of bio fuel cell ended up in 268 μW/cm2 of maximum power density [3]. A. Heller et al [4] indicated the reason of low performance is poor dissolved oxygen in solution.

1.3.1 の出だしは
http://fgamedia.org/faculty/rdcormia/NANO52/pdf/MEMS%20Handout.pdf　（Microelectronics/MEMS/NEMS）からのコピペです。
同一文章３ ：黄色でハイライトされた部分が同一文章です。 
1.3 Miniaturized System by Using Micro Fabrication Technology1.3.1 MEMSMicro-Electro-Mechanical Systems (MEMS) is the integration of mechanical elements, sensors, actuators, and electronics on a common silicon substrate through micro-fabrication technology. While the electronics are fabricated using integrated circuit (IC) process sequences (e.g., CMOS, Bipolar, or BICMOS processes), the micromechanical components are fabricated using compatible "micromachining" processes that selectively etch away parts of the silicon wafer or add new structural layers to form the mechanical and electromechanical devices [24]. 　MEMS promises to revolutionize nearly every product category by bringing together silicon-based microelectronics with micromachining technology, making possible the realization of complete systems-on-a-chip. MEMS is an enabling technology allowing the development of smart products, augmenting the computational ability of microelectronics with the perception and control capabilities of micro-sensors and micro-actuators and expanding the space of possible designs and applications. Microelectronic integrated circuits can be thought of as the "brains" of a system and MEMS augments this decision-making capability with "eyes" and "arms", to allow micro-systems to sense and control the environment. Sensors gather information from the environment through measuring mechanical, thermal, biological, chemical, optical, and magnetic phenomena. The electronics then process the information derived from the sensors and through some decision making capability direct the actuators to respond by moving, positioning, regulating, pumping, and filtering, thereby controlling the environment for some desired outcome or purpose. Because MEMS devices are manufactured using batch fabrication techniques similar to those used for integrated circuits, unprecedented levels of functionality, reliability, and sophistication can be placed on a small silicon chip at a relatively low cost.

1.4.1
Yohtaro Yamazaki氏の論文（Electrochimica Acta Volume 50, Issues 2–3, 30 November 2004, Pages 663–666）

同一文章４ ：黄色でハイライトされた部分が同一文章です。
1.4.1 Micro polymer electrolyte fuel cells
The micro electro mechanical systems (MEMS) technology has been developed at the various requests of environmental and internal sensors, machining of silicon and metal derivatives, optical and biomedical systems, and micro-fluidics. The prospective potential for miniaturization and economical mass production of small fuel cells is the main reason for the use of MEMS processes to fabricate the micro fuel cells; however, we can expect more basic effect of the micro fabrication process on the performance of the fuel cell of various sizes. Most conventional cell stacks of PEMFC consist of separators with gas channels, carbon paper sheets, carbon particles, and catalysis supported by carbon particles. In those stacks, the molecules of the fuel must move from the gas channels to the reaction points on the catalysis particles in the anode. In the cathode, the molecules of the air must be transported from the channels to the surface of the catalysts while the water molecules move in the reverse direction. These molecules move through the pores in the electrodes by the diffusion process. Therefore, the efficiency of the material transportation is not high. By using the MEMS technology, we can reduce these diffusion paths by preparing fine channels in which the materials pass not by diffusion but by bulk flow. This may considerably increase the performance of the cell stack [25].
　The micro electro mechanical systems (MEMS) technology has been developed at the various requests of environmental and internal sensors, machining of silicon and metal derivatives, optical and biomedical systems, and micro-fluidics. The prospective potential for miniaturization and economical mass production of small fuel cells is the main reason for the use of MEMS processes to fabricate the micro fuel cells; however, we can expect more basic effect of the micro fabrication process on the performance of the fuel cell of various sizes. Most conventional cell stacks of PEMFC consist of separators with gas channels, carbon paper sheets, carbon particles, and catalysis supported by carbon particles. In those stacks, the molecules of the fuel must move from the gas channels to the reaction points on the catalysis particles in the anode. In the cathode, the molecules of the air must be transported from the channels to the surface of the catalysts while the water molecules move in the reverse direction. These molecules move through the pores in the electrodes by the diffusion process. Therefore, the efficiency of the material transportation is not high. By using the MEMS technology, we can reduce these diffusion paths by preparing fine channels in which the materials pass not by diffusion but by bulk flow. This may considerably increase the performance of the cell stack [25].
　The components of a novel miniature fuel cell/fuel reformer system fueled by liquid gases such as butane and propane were prototyped by MEMS technology and tested. In this system, fuel, air and water are supplied to the fuel reformer by utilizing the vapor pressure of the liquid gas for the reduction of power consumption by peripherals and the simplification of the system [29]. A Micro-Polymer Electrolyte Fuel Cell (μ-PEFC) with “alternating structure” was demonstrated. The alternating structure has a series of single cells formed in one plane, and the polarization of each single cell is alternately inverted. This structure has self-assembled cell interconnection on micro-machined silicon substrates.



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投稿者 時刻:

7 件のコメント:

  1. 匿名2014年3月25日 15:33

    蜂巣　琢磨（主査　逢坂哲彌教授）　2011年
    博士論文　「Chemical Synthesis of Magnetic Nanoparticles and Their Application to Magnetic Recording」
    この論文もコピペがある。

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  2. 匿名2014年3月25日 19:36

    >>蜂巣　琢磨
    Modern Electroplating - Mordechay Schlesinger, Milan Paunovic
    Electrochemical Nanotechnologies - Tetsuya Osaka, Madhav Datta, Yosi Shacham-Diamand
    Developments in Data Storage: Materials Perspective - S. N. Piramanayagam, Tow C. Chong

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    1. 11jigen2014年3月25日 21:27

      ご指摘ありがとうございます。追記します。
      K O'Grady, H Laidler氏らの論文「Journal of Magnetism and Magnetic Materials Volume 200, Issues 1–3, October 1999, Pages 616–63」

      Alex Taratorin's book "Magnetic Recording Systems and Measurements" Chapter 7 "Perpendicular Recording"

      などからもコピペがあるようです。引き続き調べます。

      削除
  3. 匿名2014年3月25日 21:34

    本川 慎二　（指導教授 逢坂 哲彌）　2005年
    ごく軽くチェックしました。まだあるとは思います。

    1st pragraph
    A fuel cell uses the chemical energy of hydrogen to produce electricity and water, cleanly and efficiently. Fuel cells are unique in terms of the variety of their potential applications;

    途中からはgoogle booksでFuel Cell Systems L.J.M.J. Blomen、M.N. Mugerwa のパクリでしょう
    The situation is more complex when organic fuels (hydrocarbons and alcohols) are used directly in fuel cells. The exchange current densities for these reactions are as low as or even lower than those for the oxygen reduction reaction.

    1.3.1 の出だしはmems
    http://fgamedia.org/faculty/rdcormia/NANO52/pdf/MEMS%20Handout.pdf

    1.4.1は The micro electro mechanical systems (MEMS) technology has been developed at the various requests of environmental and internal sensors, machining of silicon and metal derivatives, optical and biomedical systems, and micro-fluidics.
    で検索すると出てくるワードの文書

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  4. 匿名2014年3月26日 18:36

    上記エントリ本文中の
    「
    調査２：蜂巣 琢磨氏（早稲田大学の逢坂哲彌氏の研究室）の博士論文における文章のコピペについてのまとめ

    著者：　逢坂哲彌
    論文題目：　「Chemical Synthesis of Magnetic Nanoparticles and Their Application to Magnetic Recording」
    」
    は、
    著者：　逢坂哲彌
    ではなく
    著者：　蜂巣琢磨
    ではないでしょうか。
    単純な転記ミスと思われます。

    返信削除
  5. 11jigen2014年4月18日 14:25

    奈良洋希氏、蜂巣琢磨氏の博士論文の本分がデータベースから削除されました。

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  6. 匿名2015年7月10日 20:01

    Great post. I have in fact enjoyed reading your website posts. I have been googling blogs and sites in related manner recently and i have to state you have a nice
    Electrolyte Analyzer

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