Nano Plating - Microstructure Formation Theory of Plated Films and a Database of Plated Films

Nano-Plating Microstructure Control Theory of Plated Film and Data Base of Plated Film Microstructure Elsevier Intern...

Author: T Watanabe

94 downloads 1123 Views 37MB Size Report

This content was uploaded by our users and we assume good faith they have the permission to share this book. If you own the copyright to this book and it is wrongfully on our website, we offer a simple DMCA procedure to remove your content from our site. Start by pressing the button below!
Report copyright / DMCA form

DOWNLOAD PDF

Nano-Plating Microstructure Control Theory of Plated Film and Data Base of Plated Film Microstructure

Elsevier Internet Homepage - http://www.elsevier.com Consult the Elsevier homepage for full catalogue information on all books, journals and electronic products and services. Elsevier Titles of Related Interest Books CHRISTIAN The Theory of Transformations in Metals and Alloys ISBN: 0-08-0440193

NALWA Handbook of Surfaces and Interfaces of Materials, 5-Volume set ISBN: 0-12-513910-1

MURAKAMI Metal Fatigue: Effects of Small Defects and Nonmetallic Inclusions ISBN: 0-08-044064-9

OHRING The Materials Science of Thin Films (2nd Edition) ISBN: 0-12-524975-6

NALWA Handbook of Thin Films, 5-Volume Set ISBN: 0-12-512908-4 Journals Free specimen copy gladly sent on request. Elsevier Ltd, The Boulevard, Langford Lane, Kidlington, Oxford, OX5 1GB,UK Acta Materialia Intermetallics Journals of Alloys and Compounds Journal of Materials Processing Technology Materials Chemistry and Physics

Materials Letters Scripta Materialia Surface and Coatings Technology Thin Solid Films

Series Semiconductors and Semi Metals Solid State Physics Thin Films and Nanostructures (formerly known as Physics of Thin Films, volumes 1-20; Thin Films, volumes 21-29) To Contact the Publisher Elsevier welcomes enquiries concerning publishing proposals: books, journal special issues, conference proceedings, etc. All formats and media can be considered. Should you have a publishing proposal you wish to discuss, please contact, without obligation, the publisher responsible for this title. Iain Craig Publishing Editor Elsevier Ltd The Boulevard, Langford Lane Kidlington, Oxford OX5 1GB, UK

Phone: +44 1865 843631 Fax: +44 1865 843987 E-mail: [email protected]

General enquiries, including placing orders, should be directed to Elsevier’s Regional Sales Offices – please access the Elsevier homepage for full contact details (homepage details at the top of this page).

Nano-Plating Microstructure Control Theory of Plated Film and Data Base of Plated Film Microstructure

Tohru Watanabe Department of Applied Chemistry Tokyo Metropolitan University Minami-ohsawa, Hachioji-shi Tokyo, Japan

2004

Amsterdam – Boston – Heidelberg – London – New York – Oxford Paris – San Diego – San Francisco – Singapore – Sydney – Tokyo

ELSEVIER B.V. Sara Burgerhartstraat 25 P.O.Box 211, 1000 AE Amsterdam, The Netherlands

ELSEVIER Inc. 525 B Street, Suite 1900 San Diego, CA 92101-4495 USA

ELSEVIER Ltd The Boulevard, Langford Lane Kidlington, Oxford OX5 1GB UK

ELSEVIER Ltd 84 Theobalds Road London WC1X 8RR UK

q 2004 Elsevier Ltd. All rights reserved. This work is protected under copyright by Elsevier Ltd, and the following terms and conditions apply to its use: Photocopying Single photocopies of single chapters may be made for personal use as allowed by national copyright laws. Permission of the Publisher and payment of a fee is required for all other photocopying, including multiple or systematic copying, copying for advertising or promotional purposes, resale, and all forms of document delivery. Special rates are available for educational institutions that wish to make photocopies for non-profit educational classroom use. Permissions may be sought directly from Elsevier’s Rights Department in Oxford, UK: phone (+44) 1865 843830, fax (+44) 1865 853333, e-mail: [email protected]. Requests may also be completed on-line via the Elsevier homepage (http://www.elsevier.com/locate/permissions). In the USA, users may clear permissions and make payments through the Copyright Clearance Center, Inc., 222 Rosewood Drive, Danvers, MA 01923, USA; phone: (+1) (978) 7508400, fax: (+1) (978) 7504744, and in the UK through the Copyright Licensing Agency Rapid Clearance Service (CLARCS), 90 Tottenham Court Road, London W1P 0LP, UK; phone: (+44) 20 7631 5555; fax: (+44) 20 7631 5500. Other countries may have a local reprographic rights agency for payments. Derivative Works Tables of contents may be reproduced for internal circulation, but permission of the Publisher is required for external resale or distribution of such material. Permission of the Publisher is required for all other derivative works, including compilations and translations. Electronic Storage or Usage Permission of the Publisher is required to store or use electronically any material contained in this work, including any chapter or part of a chapter. Except as outlined above, no part of this work may be reproduced, stored in a retrieval system or transmitted in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, without prior written permission of the Publisher. Address permissions requests to: Elsevier’s Rights Department, at the fax and e-mail addresses noted above. Notice No responsibility is assumed by the Publisher for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions or ideas contained in the material herein. Because of rapid advances in the medical sciences, in particular, independent verification of diagnoses and drug dosages should be made.

First edition 2004 Library of Congress Cataloging in Publication Data A catalog record is available from the Library of Congress. British Library Cataloguing in Publication Data A catalogue record is available from the British Library. ISBN: 0-08-044375-3 1 The paper used in this publication meets the requirements of ANSI/NISO Z39.48-1992 (Permanence of Paper). W Printed in The Netherlands.

Preface Purpose of Publication and Background Plating technology has existed for more than 2,000 years and has been practiced steadily over the last two millennia. Modern plating technology is highly advanced, and has developed to cover a wide range of applications, e.g. in addition to the traditional use for surface finishing, plating technology can now offer novel processes to fabricate highperformance films or fine microstructural bodies in the microelectronics industry. This rapid progress reflects the potential for the electroplating method to become one of today’s leading-edge technologies. The development of plating technology, however, has been slow due to the lack of a sound theoretical foundation and has often relied on a method of trial and error to obtain films with desirable properties. Despite numerous attempts, a unified theory has yet to be established that can provide this control in a consistent manner. This book is a compilation of vast amounts of experimental results that we have obtained in collaboration with our students over the last 34 years. We will introduce the concept of our new “Microstructure Control Theory” for plated films, which is entirely different from previous theories. We will then describe various experimental results that prove the validity of our theory. Finally, a large collection of experimental data on plated metal/alloy systems will be presented with a special emphasis on their microstructure. In the early days of research, we started modestly with fundamental studies on the microstructure of electroless films and on epitaxial phenomenon occurring between a plated film and a substrate. We presented these results as an occasional scientific paper in technical society meetings. During the course of these studies, we became increasingly aware of the possibility of producing amorphous films by plating methods. In 1988, motivated by this curiosity, we proposed a theory that explains why plated films can form an amorphous phase (cf. J. Surface Finish Soc. Japan 40, 375 (1989)). This theory was found to be also applicable to crystalline materials, giving us an opportunity to develop a theory explaining how the microstructure of plated pure metal/alloy films evolves. From this theoretical exercise, it has become apparent that the microstructure of plated films is closely connected to their equilibrium phase diagram. This concept finally led to the establishing of a theory of microstructure control for plated films. In 1999, we found that the microstructure of plated films can be categorized into 7 types. Each microstructural type is unique, and thus could be controlled independently (The 2nd Thin Film Basic Seminar, “Plating Methods”, Japan Surface Science Society (1999) p. 115; Materia 40, 871 (2001)). Based on this microstructure classification, we are now able to focus on each individual microstructure type and develop its control theory independently. We believe that our efforts are finally paying dividends, as if tangled threads are being separated. v

vi

Preface

It is clear that in recent years much progress on theoretical work dealing with the microstructural control of plated films has been made. After leading plating technology discussion groups for the last 10 years, we sensed an urgent need for a specialized book on the microstructure control theory among researchers/engineers in the plating industry. For this reason, we take the liberty to introduce this book although topics described within are still far from completion. At the same time, we hope that the accomplishments made in plating technology during the 20th century can be conveyed through this book into the 21st century. To describe an effective manner of conducting plating research, we explain research and experimental methods of plating. We list the microstructures of 53 types of plated pure metals and alloys, which should serve as a useful database for plated films. The unique feature of this database is that most of the plating baths are chosen to be simple and contain no additives. In addition, amorphous materials are used as substrates to avoid the effect of the substrate structure, and single-crystal substrates are chosen to study the epitaxial growth phenomenon. As mentioned above, the coverage in this book is still incomplete. Therefore, our mission is to continue collecting experimental data to fill the many gaps left in this book. At an opportune time, therefore, we will update the book accordingly. We believe that the materials contained within this book represent a crystallization of our students’ efforts in conducting plating experiments over the last 34 years. We wish to express our sincere thanks to all the students and collaborators involved and list their names below. Thanks are also due to many companies, which funded, totally or partially, our research activities over the years. Finally my sincere personal thanks go to Dr. Shohei Nakahara for his invaluable assistance in the translation of this book from Japanese into English. Faculty Professor Emeritus Yoshimi Tanabe, Tokyo Metropolitan University. Professor Emeritus Nobuyoshi Baba, Tokyo Metropolitan University. Professor Shohei Nakahara, University of Limerick Professor Emeritus Ryouichi Urao, Ibaragi University. Professor Takeshi Nakata, Shibaura Institute of Technology. Secretary Yoshinari Misaki. Laboratory Faculty and Senior Associates Professor Kazuhisa Ishibashi, Tokai University. Shigeo Urai, Toyama National College of Technology. Shouzo Asano, Senju Metal Co., Ltd. Visiting Professors and Scientists Wei-Ping Yu (China), Dr. Go¨ran Holmbom (Sweden), Dr. Torben Tang (Denmark),

Preface

vii

Dr. Xin-Tan Hu (China), Su-Wei Yao (China), Dr. Imre Bakonyi (Hungary) and M.Sc. Anett Alsted Rasmussen (Denmark). Research Students (Domestic) Shinji Katou, Kazuo Shimizu, Suguru Abe, Yasutaka Mogi, Dr. Susumu Arai, Hiroko Furusawa, Katsuhiko Tashiro, Hiroshi Tsukamoto, and Hirokazu Takagi. Trainee (Foreign) Yin-Qi Zhang. Company Doctor Course Graduates Motonobu Onoda (Nippon Piston Ring Co., Ltd) and Kiyoshi Itoh. Doctor Course Graduates Tokyo Metropolitan University Seiji Kamasaki, Yasuo Shimizu, Masayuki Kakegawa, Hisakazu Ito, Hisashi Furuya, and Naoki Fukumuro, and Feng Wang. Shibaura Institute of Technology Koichiro Inoue. Master Course Graduates Tokyo Metropolitan University Ken Tone, Hiroshi Matsubayashi, Toshihisa Sudo, Hiroki Shimizu, Hiroshi Ikebuchi, Hiroyasu Kojima, Hideo Suda, Hiroshi Imai, Yukio Numakura, Yoshiaki Ikeda, Naoya Hasegawa, Akira Narita, Hiroyuki Yamaguchi, Norimoto Usuzaka, Satoru Katsumata, Kazuyoshi Arai, Kenji Takahashi, Akira Suzaki, Hai-Ying Liang, Io Mizushima, Toshie Murai, Naoyuki Igarashi, Isao Satou, Kei Imafuji, and Takayuki Ikeda. Yuuichiro Miura and Shouji Matsuda. Shibaura Institute of Technology Takanobu Kanayama, Kenji Ikejima, Ichiro Wada, and Hideki Kotsuji. Tomoya Teshigawara, Takako Terakado, Yuki Makino, and Nobuto Sasaki. Undergraduates Tokyo Metropolitan University Toshiaki Ogura, Toshiie Kurihara, Minoru Kanda, Hiroshi Nakajima, Isao Kawaida, Mitsuru Kasahara, Kouhei Kitukawa, Akiyuki Kuniyoshi, Masao Soranishi, Kenzou Matsui, Yukio Matsumoto, Isao Ikegaya, Minoru Uehara, Hiroshi Takashio, Fumiaki Komatsu, Takakazu Fukuchi, Fujio Kawashima, Zenzaburo Naito, Kazuo Hosokawa, Hideo Tomita, Mamoru Arai, Nobuyuki Kataigi, Takashi Inoue, Kazuo Fukuda, Satomi Inoue, Takashi Eguro, Yoshio Kubota, Kazuyoshi Nishizawa,

viii

Preface

Minoru Odaka, Masaya Naoi, Satoru Shinohara, Nobuo Tomizawa, Susumu Morohashi, Shinji Tobita, Yukihisa Hiroyama, Takayuki Anma, Jun Tanaka, Takuya Naoe, Keiichi Nozawa, Tetsuo Sakamoto, Wakana Wasa, Takeshi Hirose, Hiroyuki Hoshina, Toshiko Kitagawa, Naoyuki Kanami, Yoshiaki Fukuda, Tomokazu Takasaka, Shusuke Minami, Shinichi Miyazaki, Saeko Yoshioka, Hroo Sawanobori, Mari Tomita, Megumi Osada and Yuji Hanaie. Shibaura Institute of Technology Atushi Mituo, Jun Kubo, Sadanori Tanemura, Junko Hirose, Kenji Asada, and Kazuyuki Koide. Akihiro Inami. Current Students (2004) Doctor Course Tokyo Metropolitan University Takashi Sugizaki (Meltex Inc.), Naoki Okamoto. Master Course Tokyo Metropolitan University Jin-Song Qiu, Sayaka Doi, Kaori Hosoiri, Atushi Kondo and Mituhisa Funatu. Shibaura Institute of Technology Takuya Yoshihara Takahiro Makino and Miki Yachidate. Undergraduate (Senior) Tokyo Metropolitan University Keigo Hoshina and Satoko Shoda. Shibaura Institute of Technology Kyohei Komori, Shinya Okahara, and Kazuya Kitazawa. Companies that Provided Research Funds Alps Electric Co., Ltd. Nippon Piston Ring Co., Ltd. Nihon Steel Corp. Sumitomo Metal Industries, Ltd. Hitachi Ltd. Sumitomo Heavy Industries, Ltd. Ebara-Udylite Co., Ltd. Asahi Glass Co., Ltd. Mitsui Mining & Smelting Co., Ltd. N. E. Chemcat Corp. Mitsubishi Aluminum, Ltd. Mitsubishi Materials Co., Ltd.

Preface

ix

Kyocera Corp. Electroplating Engineering of Japan Ltd. Hitachi Cable Ltd. Limited Liability Company Taw Konika Technology Center Corporation Toppan Prating Co., Ltd. Fuji Electric Co., Ltd This book was originally published in Japanese in February 2003. The title was “FINE PLATING: Microstructure Control and Analysis Methods for Plated Films” published by Technical Information Association, Co., Ltd. This book is an English version. Dr. Tohru Watanabe Department of Applied Chemistry Graduate School of Engineering Tokyo Metropolitan University Minami-ohsawa, Hachioji-shi, Tokyo, 129-0397 Japan

This Page Intentionally Left Blank

Contents Preface

v

CHAPTER 1 MICROSTRUCTURE CONTROL THEORY OF PLATED FILM 1.1. Introduction 1.2. Research Process of Plating Technology 1.3. Review of Previous TMC and Practical Plated Films 1.4. Theory of Microstructure Control for Plated Films 1.4.1 Metallurgical Structure (Crystalline, Solid-Solution, Intermetallic Compound, Meta-Stable Phase, Amorphous Phase, Mixed Phases) 1.4.2 Surface Morphology (Leveling, Brightness, Surface Irregularities/ Form, Dendrite, etc.) 1.4.2.1 Surface Morphology Change with Increasing Film Thickness 1.4.2.2 Surface Morphology Change with Current Density (Overpotential) 1.4.2.3 Surface Morphology Change with the Type of Anions 1.4.2.4 Surface Morphology Change with Solution Temperature 1.4.2.5 Surface Morphology Change with Solution Agitation 1.4.2.6 Surface Morphology of Alloy Films 1.4.2.7 Summary of the Formation Principles of Surface Irregularities 1.4.2.8 Formation of Dendrites 1.4.2.9 Effect of Brighteners 1.4.3 Grain Size (Grain Size, Granular, Spherical, Columnar, Needle-Like, etc.) 1.4.3.1 Grain Size of Pure Metal Deposits 1.4.3.2 Grain Size of Alloy Deposits 1.4.4 Preferred Orientation (Texture) (Film Normal/Film Plane Direction and Assembled Structure) 1.4.4.1 Experimental Methods for Determining the Texture of Plated Films 1.4.4.2 Relationship Between the Texture of Various Plated Films and Plating Conditions 1.4.4.3 Effect of Plating Conditions on the Film Texture 1.4.4.4 Texture of Electroless Films 1.4.4.5 Summary xi

3 4 7 10 10 14 14 24 29 32 32 32 36 37 38 39 39 41 46 50 50 57 66 68

xii

Contents

1.4.5 Bonding with Substrate and Crystallographic Matching (Epitaxy) 1.4.6 Residual Stress (Compressive, Tensile Stresses, Cracks) 1.4.7 Anomalous Morphology 1.4.7.1 Nodules 1.4.7.2 Pits 1.4.7.3 Cracks 1.4.7.4 Formation of Layer Structure 1.4.7.5 Initial Layer 1.4.7.6 Whisker References

68 74 79 79 82 82 83 88 91 91

CHAPTER 2 FILM FORMATION MECHANISM IN ELECTRODEPOSITION 2.1. Formation of Electrolytic Films 2.1.1 The Initial Nucleation and Growth Stages of Pure Metal Films 2.1.2 Structure of Electrolytic Binary Alloy Films 2.1.2.1 Formation of Phase Separation Type 2.1.2.2 Formation of Solid –Solution Alloy 2.1.2.3 Formation of Intermetallic Compounds 2.1.2.4 Meta-Stable Phase 2.1.2.5 Amorphous Phase 2.1.2.6 Mixed Phase 2.2. Electroless Films 2.2.1 Electroless Plating Method 2.2.1.1 Electroless Plating Method and Microstructure 2.2.1.2 Formation Mechanism of Electroless Films 2.2.1.3 Bonding Between Electroless Films and Substrate 2.2.2 Immersion Coating 2.2.2.1 Film Formation Mechanism in Immersion (Displacement) Coating 2.2.2.2 Bonding between Immersion Coating and Substrate 2.2.3 Contact Plating References

132 136 136 138

CHAPTER 3 PLATING IN ORGANIC ELECTROLYTE 3.1. Introduction 3.2. Experimental Methods 3.3. Selection of Plating Electrolyte 3.3.1 Surface Morphologies of Plated Films

143 144 146 146

97 97 106 107 108 109 109 113 119 121 121 121 122 127 132

Contents

3.3.2 Crystal Structure of Plated Films 3.4. Cobalt Deposition from Various Organic Solvent Solutions 3.4.1 Grain-Refining Effect 3.4.2 Inclusions in Plated Films 3.4.2.1 H2O Bath 3.4.2.2 DMSO Bath 3.4.2.3 FA Bath 3.4.3 Relationship Between Solvent Decomposition Product and Grain Size 3.4.4 Surface Morphology of Plated Cobalt Films 3.4.5 Crystal Orientation and Form References CHAPTER 4 MICROSTRUCTURAL CHANGES IN PLATED FILMS DURING HEAT TREATMENT 4.1. Introduction 4.2. A Change of Meta-Stable and Non-Equilibrium Phases to Stable Crystals 4.3. Transformation of Amorphous Phase to Equilibrium Phase References CHAPTER 5 CONTROL OF MACROSTRUCTURE IN PLATED FILMS AND FABRICATION OF THREE-DIMENSIONAL MICROSTRUCTURE 5.1. Introduction 5.2. Macrostructure Control in Plated Films 5.2.1 Columnar Structure 5.2.2 Fine-Crystal Structure (Nanocrystals) 5.2.3 Amorphous Structure 5.2.4 Single Crystals 5.2.5 Multi-Layer (ML) Film 5.2.6 Multi-Layer (ML) Containing Alternating Layers of Crystalline and Amorphous Phases 5.2.7 Epitaxial Multi-Layer 5.2.8 Amorphous/Crystalline Graded Structure 5.2.9 Composite Coating 5.2.10 Graded Composite Coating 5.2.11 Other Structures 5.3. Fabrication of Three-Dimensional Microstructural Body References

xiii 150 152 161 162 162 163 165 166 168 169 170

175 176 182 192

197 197 197 197 199 199 199 200 200 202 202 202 202 204 204

xiv

Contents

CHAPTER 6 CHARACTERIZATION METHODS FOR PLATED FILMS 6.1. Structure of Metal Substrates 6.1.1 Deformation and Annealing Texture 6.1.2 Surface Deformation Layer 6.1.3 Single-Crystal Substrate 6.1.4 Pre-Treatment Methods for Polycrystalline Substrates 6.1.4.1 Annealing Treatment 6.1.4.2 Electropolishing 6.2. Structural Determination of Plated Films 6.2.1 XRD Method 6.2.2 Structural Analysis by TEM 6.2.2.1 Observation of the Initial Stages of Electrodeposition 6.2.2.2 Plan-View Observation of Thick Plated Films 6.2.2.3 Crystallographic Matching Relationship Between Plated Film and Substrate 6.2.2.4 Preparation of TEM Cross-Section Samples 6.3. Grain-Size Measurements 6.3.1 Direct Measurements by TEM Observations 6.3.2 X-Ray Diffraction Method 6.4. Observations of Surface Morphology 6.4.1 Scanning Electron Microscope (SEM) 6.4.2 Replica Method for TEM Observations 6.4.3 Observations by AFM 6.4.4 Measurements by a Surface Profilometer 6.5. Measurement of Preferred Orientation 6.5.1 Definition of Preferred Orientation and its Description 6.5.2 Measurement of Texture in Plated Films References

223 223 235 235 238 240 240 246 248 248 248 250 250 252

CHAPTER 7 DATABASE FOR THE MICROSTRUCTURE OF PLATED FILMS 7.1. Introduction 7.2. Plating Methods 7.2.1 Plating Conditions 7.2.2 Substrate Materials 7.2.2.1 Polycrystalline Copper Sheets 7.2.2.2 Single-Crystal Copper ({100}, {110}, and {111} planes) 7.2.2.3 Stainless Steel Sheets (SUS-304) 7.2.2.4 Electrolytic Ni-25 at.% P Alloy Films

257 257 257 258 258 258 258 259

209 209 209 212 212 212 213 215 215 216 218 221

Contents

Splat-Quenched Amorphous Alloy Foils (Fe –Si –B or Fe –Co – Si – B Alloys Manufactured by Nippon Amorphous Metals Co., Ltd.) 7.3. Microstructure Observations and Measurements Methods for Plated Films 7.4. Database for the Microstructure of Plated Films 7.4.1 Pure Metals 7.4.1.1 Electrolytic Ag 7.4.1.2 Electrolytic Au 7.4.1.3 Electrolytic Cd 7.4.1.4 Electrolytic Co 7.4.1.5 Electrolytic Cr 7.4.1.6 Electrolytic Cu 7.4.1.7 Electrolytic Fe 7.4.1.8 Electrolytic Ni 7.4.1.9 Electrolytic Sn 7.4.1.10 Electrolytic Zn 7.4.2 Pure Alloys 7.4.2.1 Electrolytic Ag – Cd 7.4.2.2 Electrolytic Ag – Co 7.4.2.3 Electrolytic Ag – Cu 7.4.2.4 Electrolytic Ag – Sn 7.4.2.5 Electrolytic Ag – Zn 7.4.2.6 Electrolytic Al– Mn 7.4.2.7 Electrolytic Au – Cu 7.4.2.8 Electrolytic Au – Ni 7.4.2.9 Electrolytic Au – Pd 7.4.2.10 Electrolytic Au – Sn 7.4.2.11 Electrolytic Cd –Sn 7.4.2.12 Electrolytic Cd –Zn 7.4.2.13 Electrolytic Co –Cu 7.4.2.14 Electrolytic Co –Fe 7.4.2.15 Electrolytic Co –Mo 7.4.2.16 Electrolytic Co –Ni 7.4.2.17 Electrolytic Co –Sn 7.4.2.18 Electrolytic Co –W 7.4.2.19 Electrolytic Cr– H 7.4.2.20 Electrolytic Cu –Ni 7.4.2.21 Electrolytic Cu –Pb 7.4.2.22 Electrolytic Cu –Sb 7.4.2.23 Electrolytic Cu –Sn

xv

7.2.2.5

259 259 260 260 260 274 296 308 318 323 333 344 358 369 384 384 387 398 407 416 417 423 427 437 446 455 457 459 466 476 482 495 501 510 518 525 532 533

xvi

Contents

7.4.2.24 Electrolytic Cu – Zn 7.4.2.25 Electrolytic Fe –Mo 7.4.2.26 Electrolytic Fe –Ni 7.4.2.27 Electrolytic Fe –W 7.4.2.28 Electrolytic Fe –Zn 7.4.2.29 Electrolytic In – Sn 7.4.2.30 Electrolytic Ni –B 7.4.2.31 Electrolytic Ni –Mo 7.4.2.32 Electrolytic Ni –P 7.4.2.33 Electrolytic Ni –S 7.4.2.34 Electrolytic Ni –Sn 7.4.2.35 Electrolytic Ni –W 7.4.2.36 Electrolytic Ni –Zn 7.4.2.37 Electrolytic Sn –Zn 7.4.3 Electroless Plating 7.4.3.1 Electroless Ni– B 7.4.3.2 Electroless Ni– P 7.4.4 Displacement Plating 7.4.4.1 Displacement Ag 7.4.4.2 Displacement Au 7.4.4.3 Displacement Cd 7.4.4.4 Displacement Cu 7.4.4.5 Displacement Zn INDEX

544 546 555 566 577 589 599 605 611 624 630 642 652 655 657 657 666 674 674 680 686 690 692 697

Chapter 1

Microstructure Control Theory of Plated Film 1.1. 1.2. 1.3. 1.4.

Introduction Research process of plating technology Review of previous TMC and practical plated films Theory of microstructure control for plated films 1.4.1 Metallurgical structure (crystalline, solid-solution, intermetallic compound, meta-stable phase, amorphous phase, mixed phases) 1.4.2 Surface morphology (leveling, brightness, surface irregularities/ form, dendrite, etc.) 1.4.2.1 Surface morphology change with increasing film thickness 1.4.2.2 Surface morphology change with current density (overpotential) 1.4.2.3 Surface morphology change with the type of anions 1.4.2.4 Surface morphology change with solution temperature 1.4.2.5 Surface morphology change with solution agitation 1.4.2.6 Surface morphology of alloy films 1.4.2.7 Summary of the formation principles of surface irregularities 1.4.2.8 Formation of dendrites 1.4.2.9 Effect of brighteners 1.4.3 Grain size (grain size, granular, spherical, columnar, needle-like, etc.) 1.4.3.1 Grain size of pure metal deposits 1.4.3.2 Grain size of alloy deposits 1.4.4 Preferred orientation (texture) (film normal/film plane direction and assembled structure) 1.4.4.1 Experimental methods for determining the texture of plated films 1.4.4.2 Relationship between the texture of various plated films and plating conditions 1.4.4.3 Effect of plating conditions on the film texture 1.4.4.4 Texture of electroless films 1.4.4.5 Summary 1.4.5 Bonding with substrate and crystallographic matching (epitaxy)

3 4 7 10 10 14 14 24 29 32 32 32 36 37 38 39 39 41 46 50 50 57 66 68 68

2

Nano-Plating

1.4.6 1.4.7

References

Residual stress (compressive, tensile stresses, cracks) Anomalous morphology 1.4.7.1 Nodules 1.4.7.2 Pits 1.4.7.3 Cracks 1.4.7.4 Formation of layer structure 1.4.7.5 Initial layer 1.4.7.6 Whisker

74 79 79 82 82 83 88 91 91

Chapter 1

Microstructure Control Theory of Plated Film 1.1

INTRODUCTION

The ultimate goal of conducting research in the electroplating technology field is to investigate how to produce plated films with desirable mechanical, physical, and chemical properties that meet particular application requirements. In the past, a vast number of papers have been published in the field of electroplating technologies, and their primary emphasis has been placed on electrochemical studies. At the same time, various attempts have been made to establish a unified theory that allows one to control both the microstructure and properties. Attempts to establish such a theory have been unsuccessful due to the complexities in achieving controlled experimental conditions. It is often the case that while fixing one parameter, others change uncontrollably, i.e. all the experimental parameters cannot be fixed at the same time. Furthermore, not all the chemical reactions occurring during electrodeposition processes are well understood. Because of these experimental difficulties and uncertainties, a reliable theoretical work that attempts to find a link to electroplating technologies has been largely hindered. From our extensive electrodeposition research in the past, we have come to the conclusion that all the physical properties of electroplated films must originate from their microstructure. For example, in deriving the microstructure – alloy composition relationship in electroplated alloy films, we found that the microstructural details provided more relevant information than the knowledge of the chemical reactions occurring during electrodeposition. If we realize that “electroplated films are metallurgical materials and thus their properties are closely related to their equilibrium phase diagram”, “we can explain the microstructure logically using a metallurgical concept”. Understanding of the microstructure led to the development of a theory that suggests a way of controlling the microstructure of plated films. This theory will be called a theory of microstructure control (TMC). Other physical properties, such as the surface morphology, grain size and texture, can also be controlled independently in a similar manner. Based on our previous studies as well as reports from the literature, we constructed seven categories of microstructures in plated films, as depicted in Figure 1.1. Once the TMC is fully developed for each microstructural category, we will be able to control the physical properties of all categories of films. In this book, we will attempt to explain the TMC concept and discuss the validity of the theory using microstructural data for various plated films. 3

4

Nano-Plating

Figure 1.1. Seven microstructure types observed in plated films.

1.2

RESEARCH PROCESS OF PLATING TECHNOLOGY

A process diagram describing electrodeposition research is summarized in Figure 1.2. In the past, most of research has been conducted by varying plating conditions by a trial and error method based on an empirical rule stemming from a large number of experimental

Microstructure Control Theory of Plated Film

5

Figure 1.2. A process diagram for conducting electrodeposition research.

results. The properties of the resulting film were then matched with the properties of interest, applicable to a specific research objective. The problem with this approach is that the numbers of plating conditions are extremely large, as described below. The possible variables to be considered prior to plating include: (a) (b) (c) (d) (e)

Substrate material (metal, non-metal (ceramics, plastic)). Pre-treatment of substrate materials (deoxidation, degreasing, washing, drying). Type of metals (pure metal, alloy). Solution concentration (concentration, alloy composition). Type of metallic salts (effect of anion type).

6

Nano-Plating

(f) (g) (h) (i) (j) (k) (l) (m) (n) (o) (p)

Type and amount of complexing agents. Solution pH, pH adjuster, and pH buffer. Additives (brightener, leveling agent, stress reducer, supporting electrolyte). Current density (DC, pulse, cathodic/anodic current density). Overpotential. Shape of plating tank. Type of anode materials (soluble, non-soluble). Inter-electrode distance. Agitation (static, stirrer, bubbling, ultrasonic, vibration). Plating temperature. Other conditions.

Conditions changing during electroplating include: (a) (b) (c) (d) (e) (f) (g) (h) (i) (j) (k) (l)

Metal concentration. Consumption of additives. pH. Structural change of complexing agent. Solution temperature. From deposition on the substrate material to deposition on the plated film. Surface morphology change with increasing film thickness. Potential change due to constant current condition. Current density change due to constant potential. Generation of hydrogen gas on cathode. Gas evolution at anode. Other conditions.

In addition to the above variables, we need to accurately know the type and concentration of all impurities present in an electrolyte. For example, the commercial high-purity chemicals used for making a plating solution generally list the type and concentration of metallic impurities, but do not mention organic impurities. The concentration of organic impurities is an important parameter in plating because even a small quantity, in the order of a few ppm, is known to significantly affect the film structure. An analysis of organic impurities, however, is not straightforward. Furthermore, the experimental factors mentioned above are for pure metals only, meaning that additional variables have to be taken into account for the case of alloy deposits. The combinations of these variables amount to astronomical numbers. For this reason, it is almost impossible to develop a theoretical framework that relates the microstructure to the film properties. In practice, therefore, electrodeposition research ends up applying the traditional trial and error method.


7

Since the trial and error method is extremely inefficient, many researchers have primarily relied on electrochemical techniques, which measure the deposition overpotential and analyze the state of metal/complex ions in an electrolyte. Based on electrochemical/ chemical information, a theoretical framework has been formulated to explain the electrodeposition phenomenon. Theories thus far proposed, however, are often controversial, and no unified theory has yet been established. As seen in Figure 1.2, all the physical properties of plated films are directly connected to their microstructure. Previous microstructural studies are incomplete because their characterization was often performed by an X-ray diffraction method. The X-ray method generally assumes that the film microstructure is uniform. This assumption directly contradicts the fact that plated films are not uniform, particularly along the film thickness direction. The X-ray method therefore cannot provide the correct description of the film microstructure. Another difficulty encountered in the X-ray technique is a lack of structural information for non-crystalline or thermodynamically meta-stable phases. If X-ray diffraction data were not listed in the Joint Committee of Powder Diffraction Standard (JCPDS) card or were not previously reported, a detailed structural analysis by the X-ray method was often impossible. Phase composition and impurity content in plated films provide important information for interpreting the microstructural results. An elemental analysis of plated films has been conducted for metals and some selected non-metallic elements, such as phosphorus or boron, but not for light elements, e.g. carbon or nitrogen. A simple chemical analysis was sufficient when plated films were used primarily for surface finishing, but recent applications requiring a high-performance function demand a more detailed chemical analysis of various elements in the films. A small amount of organic impurities in an electrolyte is known to affect the film microstructure and their physical properties significantly. Compositional information for such impurities is essential in obtaining a correlation between the microstructure and impurity content. For example, one of the most prominent light elements in plated films is hydrogen. Hydrogen not only changes the microstructure of some metals, but also alters their mechanical properties. In addition, hydrogen is closely connected to the generation of residual stresses and cracks in plated films.

1.3

REVIEW OF PREVIOUS TMC AND PRACTICAL PLATED FILMS

Before we discuss our TMC concept, let us review previous theories of electrodeposition phenomenon. According to Gerischer (1960), hydrated or complexed metal ions present in a bulk electrolyte diffuse toward a cathode (substrate) through the potential gradient. On the cathode, the ions are stripped from the hydrate or complex ions at the so-called Helmholtz double layer. The stripped metal ions then become neutral atoms (adatoms) after

8

Nano-Plating

undergoing a discharge process with electrons on the cathode. These adatoms are subsequently adsorbed on the substrate, finally forming a crystal. Kossel (1927) and Stranski (1928) proposed a mechanism of crystal growth, whereby the adatoms migrate over the substrate via surface diffusion and finally arrive at active sites, such as kinks or steps, on the cathode surface. Gerischer (1960) described how crystal growth proceeds via spiral growth around the emerging sites of screw dislocations on the surface. Fischer (1960) and Seiter et al. (1960) proposed a two-dimensional nucleation model. Contrary to the model proposed by Kossel (1927) and Stranski (1928), Fischer (1960) and Seiter et al. (1960) assumed that the active sites of the cathode surface are occupied by adsorbed impurity atoms/molecules and therefore the normal crystal growth involving the successive adsorption and depositing of atoms at the active sites is inhibited. Consequently, crystal growth proceeds by the formation and expansion of two-dimensional nuclei. Gerischer (1960) assumed that an exchange reaction occurs between slightly ionized atoms and the cathode surface. Bockris (1964) and Bockris and Razmney (1967) called these atoms “adions”. Since such an exchange reaction involves ionized atoms rather than neutral atoms (adatoms), Vetter (1961), Haruyama (1963), Ohno (1988), Ohno and Haruyama (1991), and Winand (1994) thought that the whole crystal growth processes should be affected by the deposition overpotential. The deposition overpotential h consists of the discharge overpotential hds, diffusion overpotential hd, reaction overpotential hr, and crystallization overpotential hcr. These overpotential values then determine the microstructure of plated films. The concept of overpotential is derived from electrochemical studies and is the only accepted theory at present. Here we call it the overpotential theory (OT). The OT concept can be used to control grain sizes/surface irregularities in both pure metal and alloy deposits. We will briefly outline the OT concept here. At high overpotentials, the deposition rate becomes very high, accompanied by an increase in the adion concentration. The increased adions promote more nucleations, which refine the grains. The fine grains allow the film surface to become smoother. Conversely, at low overpotentials, the deposition rate is slow, thus allowing the grains to grow larger. The formation of large-grained films results in the development of rough surfaces. In other words, the OT assumes that surface irregularities in plated films are controlled by the grain size. This prediction appears to be reasonable, but in practice, an increase in the overpotential does not necessarily refine the grains. Since a pulse plating technique can achieve very high overpotentials, it was thought that extremely fine-grained or even amorphous films (Ohno and Haruyama, 1991) could be obtained by this technique. No successful cases have yet been reported by the pulse plating technique. In fact, amorphous films have been obtained even at low overpotentials, while crystalline films were obtained at high overpotentials. The OT appears to be inconsistent in explaining these results. In addition, the OT does not deal with a method of controlling metallurgical structures for solid-solution alloys or intermetallic compounds. There are a few papers claiming that the preferred orientation


9

of plated films can also be varied by the overpotential (Pangarov, 1962, 1965; Pangarov et al., 1963). It turned out, however, that the preferred orientation cannot be explained by the OT alone in a consistent manner (Sard et al., 1966). Winand (1994) reported that the grain size and surface morphology of plated films are affected by the anomalous adsorption of anions where the magnitude of the adsorption force determines the degree of inhibition in crystal growth. He constructed a map of the grain size and surface morphology as a function of the degree of growth inhibition by the anions, the current density, and the metal ion concentration. As discussed above, the OT cannot explain all the microstructural aspects of plated films consistently. A reason for the large gap between the theory and experiments has been attributed to the effect of incorporated impurities in the film (Eichkorn and Fischer, 1967) or to difficulties involved in measuring the state of an electrolyte (the thickness and condition of a diffusion layer and an electric double layer) on the electrode surface. Since it is almost impossible to separate the four types of overpotential described above and to measure them individually, it has been thought impossible to control the film microstructure (Winand, 1994). As a result, the development and improvement of plating technologies continue to rely on the traditional trial and error method. Plated films are often said to contain a large concentration of impurities (Eichkorn and Fischer, 1967). We consider that such a statement reflects an electrochemists’ view on plated films and is sometimes misleading. Electrochemists traditionally regarded any foreign species in plated films as an impurity, although metallurgists might consider them as light elements in alloy films. As discussed previously, a theoretical analysis of electrodeposition phenomena has been primarily performed using an electrochemical approach, such as the OT. The electrochemical approach can analyze a behavior of ionized species easily but cannot provide a clear understanding of the behavior of non-metallic elements, such as B, C, N, P, S, especially when the ionization state of these elements is not known. In addition, a quantitative analysis of these non-metallic elements is generally difficult. Therefore, it is much simpler to regard non-metallic elements as impurities rather than as light elements in alloy films. It is clear that there is a lack of metallurgical concepts in the analysis of plated films. In the chemistry/electrochemistry field, only materials having metal – metal bonds are considered metals, but those metals containing non-metallic elements, such Ni3P and Ni3B, are called inorganic crystals. In a metallurgical sense, these inorganic compounds containing non-metallic elements are also regarded as metals, more specifically alloys. It should be emphasized that the microstructure of plated metals, especially alloys, can be understood only using a metallurgical approach. We believe that a lack of metallurgical conception has been a hindrance in the understanding of the microstructure of plated films. The reduction step of individual metal ions to neutral atoms at a cathode may be consistent with a model proposed by Gerischer (1960). However, one important step that

10

Nano-Plating

has been largely ignored in the reduction step is the possible evolution of a large amount of heat during the ion-discharge process. For example, an electrical charge of 1 C corresponds to deposition of about 1016 atoms per second, accompanied by diffusion of a comparable number of anions toward the anode. Since heat generated from each discharged atom amounts to several eV, the cathode surface is expected to be highly heated. Consequently, the deposited film will undergo a rapid thermal quenching process. It is easy to see that this quenching process determines the resulting microstructure of plated films. This concept will be further elaborated in Section 1.4.1.

1.4 THEORY OF MICROSTRUCTURE CONTROL FOR PLATED FILMS (WATANABE, 1989a,b, 1990)

The microstructure of plated films can be classified into seven categories (Watanabe, 1998, 2001) as illustrated in Figure 1.1. Although there will be some overlap among these seven microstructures, it is possible to control each microstructure independently. Hereafter, we will explain a theory of individual microstructure control. 1.4.1 Metallurgical structure (crystalline, solid-solution, intermetallic compound, meta-stable phase, amorphous phase, mixed phases) Figure 1.3 summarizes all possible structures appearing in metallurgically processed pure metals as well as binary alloys. Here, alloys include not only metal – metal systems, but also metal –non-metal (B, P, As, N, etc.) systems. Even metal – hydrogen systems are included as alloys. Hydrogen is known to significantly affect the structure, as well as their physical properties (Raub, 1930; Loebichir et al., 1980; Furuya et al., 1981; Okinaka and Straschil, 1986; Jerkiewicz et al., 2000). Figure 1.3(a) depicts the structure of a pure metal. In general, the crystal structure of pure metals cannot be changed; only their grain size can be modified. For those pure metals having a phase transformation, however, it may be possible to change the structure by a plating method. The structure of binary alloys is shown in Figure 1.3(b) – (f). The structure in Figure 1.3(b) is a eutectic alloy, which separates into two phases. Solid-solution alloys of substitution type are shown in Figure 1.3(c) (note that there are solid-solution alloys of interstitial type). The structures of metastable and amorphous phases are shown in Figure 1.3(c) and (d), respectively. Finally, the structure of intermetallic compounds is illustrated in Figure 1.3(f). These structural models are called the microstructure in a metallurgical sense, and also appear in plated films. As discussed before, plated films are metallurgical materials and thus the microstructure of plated alloy films have a close relationship to the equilibrium (binary) phase diagrams. We have performed a detailed study on the microstructure of various binary alloy films obtained by electrodeposition to see how these alloy films are related to those listed in the equilibrium phase diagram. Similar studies that attempt to correlate the microstructure of


Figure 1.3. Various crystal structures (microstructures) generated in metallurgically processed pure metals and binary alloys.

11

12

Nano-Plating

plated films to the phase diagram were performed in our group (Watanabe, 1989a,b, 1990, 1992 –1993, 1994, 1995, 1998, 2001; Watanabe et al., 1989, 1999; Onoda et al., 1990; Watanabe and Masumoto, 1990; Narita and Watanabe, 1991; Watanabe and Arai, 1991; Liang et al., 1995, 1999; Arai and Watanabe, 1996; Mizushima et al., 1996) as well as by Aotani (1950, 1951, 1953), Raub (1953), Raub and Sautter (1957), Shimizu and Tanabe (1976a,b), Tanabe et al. (1976), Shimizu and Tanabe (1978), Shimizu et al. (1978), Enomoto et al. (1982), and Isaki et al. (1987). From the above studies, the following general trend was observed. If both metals have low melting point temperatures and have mutual solubility over the entire composition range, this binary alloy deposit ends up with the same structure as those listed in the equilibrium phase diagram, regardless of plating conditions. However, if the alloy system contains metals with a high melting point, and form intermetallic compounds, the structure of the resulting alloy deposits are often very complex and may not agree with the equilibrium phase diagram. This disagreement is exemplified by the absence of expected intermetallic compounds, the formation of meta-stable or amorphous phases, and the formation of supersaturated solid-solution alloys. The presence of meta-stable phases in alloy deposits often complicates an analysis of X-ray diffraction peaks. However, metastable phases are thermodynamically unstable and thus can be transformed to stable phases by heat treatment. This subject will be dealt with in Chapter 4. The reason why such meta-stable or amorphous phases are formed by electrodeposition methods can be explained as follows. In electrodeposition processes, an energy process involving the discharge of one metal ion amounts to few eV, which corresponds to the temperature of several 10,000 8C per atom. Furthermore, since the electric field (Haruyama, 1963) in the order of 107 V/cm is present at the electrolyte – substrate interface, joule heating (Hishino and Ro, 1973, 1974; Hishino, 1976) also occurs. Every time a metal ion is discharged, a high-temperature adatom is generated. In this way, a number of high-temperature adatoms migrate over the substrate surface via a surface diffusion mechanism and finally form a solid film (see Figure 1.4 (Watanabe, 2001)). A film thus formed is a solid, which is stable at high temperatures. At the same time, the thermal energy generated at the interface dissipates towards the plating solution as well as towards the substrate. The temperature of the plating solution is generally kept at room temperature, or at most at 100 8C. The heat capacity of this solution is very large and thus the cooling action of a plating solution is generally very effective. Pure metals having high melting point temperature have a fast cooling rate. This rapid cooling rate causes multiple nucleations, which result in the formation of fine-grained deposits. Conversely, low melting point metals having a slow cooling action, take a longer time to solidify and the corresponding surface diffusion distance becomes long. This results in the formation of large grains. The film purity also changes the diffusion distance of adatoms and thus affects the grain size, hence, the purer the film, the larger the grain size, and vice versa.


13

Figure 1.4. A mechanism for the formation of polycrystalline pure metal films during electrodeposition.

In alloy systems, the rapid cooling process at the cathode causes additional changes to the structure of plated alloy films. For alloys near the composition of their intermetallic compound, these will form as long as the crystallization energy of the compound is supplied. If sufficient crystallization energy is not supplied, due to rapid cooling, that particular compound cannot be generated. Instead, their meta-stable phase (Enomoto et al., 1982; Isaki et al., 1987; Povetkin and Devyatkova, 1996; Liang et al., 1999; Watanabe et al., 1999) or amorphous phase (Watanabe, 1989a, 1992; Watanabe et al., 1989; Onoda et al., 1990; Narita and Watanabe, 1991) may be formed. Unknown X-ray diffraction peaks are expected to appear from such a meta-stable phase, peaks which may not have been reported previously. Sometimes supersaturated solid-solution alloys also appear. The plating solution temperature and the deposition overpotential supply additional thermal

14

Nano-Plating

energies and thus are considered to be factors for determining the final crystal structure of plated films. The formation of meta-stable or amorphous phases will be described in more detail in Chapter 7. From the discussion above, it is clear that plated films are not only metallurgical materials, but also supercooled solids. For this reason, the structure of plated films does not necessarily coincide with that found in the equilibrium phase diagram. This can be summed up as The microstructure of plated films can be determined primarily by the type of elements and their composition is independent of plating conditions. The use of the term, primarily, suggests the possibility that the high deposition overpotential or the solution temperature might also affect the structure by providing additional energy to the thermal energy generated upon discharge of atoms. Solution agitation may also alter the deposition rate, or the cooling effect for discharged adatoms, and hence affect the final microstructure. 1.4.2 Surface morphology (leveling, brightness, surface irregularities/form, dendrite, etc.) The surface morphology of plated films varies with plating conditions. We will present micrographs showing the morphological changes and explain its probable mechanism. 1.4.2.1 Surface morphology change with increasing film thickness. Figures 1.5– 1.12 show a change in the surface morphology with film thickness for electroplated Sn, Cd, Zn, Ag, Cu, Ni, Co, and Fe films, respectively. Amorphous alloy substrates are used for all the deposits. For all cases, the surface irregularities increase with increasing film thickness. Figure 1.13 illustrates a mechanism as to how these surface irregularities develop with increasing film thickness. Figure 1.13(a) shows the cross-sectional view of the amorphous substrate. A film deposited immediately after the plating started is flat and uniform in thickness (see Figure 1.13(b)). Note that prior to plating, the solution composition on the cathode surface is the same as that of the bulk electrolyte. After the potential is applied to this substrate, metal ions in close contact with the cathode are immediately reduced by electrons present at the cathode. This reduction will form fine-grained metal films, which cover the substrate surface uniformly. At the same time, the discharge induces the creation of a metal-ion denuted layer (MIDL) immediately above the substrate surface (see Figure 1.13(b)– (d)). The thickness of the MIDL will vary over the substrate surface, depending on the extent of the discharge event. The thickness variation of the MIDL directly affects subsequent metal deposition. It is clear that the metal ion discharge process will no longer


15

Figure 1.5. A surface morphology change with varying film thickness and current density in electrolytic tin films grown on amorphous substrates.

16

Nano-Plating

Figure 1.6. A surface morphology change with increasing film thickness in electrolytic cadmium films grown on amorphous substrates from (A) sulfate and (B) chloride baths.


17

Figure 1.7. A surface morphology change with increasing film thickness in electrolytic zinc films grown on amorphous substrates at the current density of 100 A/m2 from a sulfate bath.

18 Nano-Plating Figure 1.8. A surface morphology change with increasing film thickness in electrolytic silver films grown on amorphous substrates at 150 A/m2. The solution temperature is 60 8C and pH is 2.


19

Figure 1.9. A surface morphology change with increasing film thickness in electrolytic copper films grown on amorphous substrates at (a) 500 and (b) 1000 A/m2 from a sulfate bath.

20 Nano-Plating Figure 1.10. A surface morphology change with increasing film thickness in electrolytic nickel films grown on amorphous substrates at (a) 100, (b) 500, and (c) 1000 A/m2 from (A) chloride bath and (B) sulfate bath.


21

Figure 1.11. A surface morphology change with increasing film thickness in electrolytic cobalt films grown on amorphous substrates from (A) sulfate and (B) chloride baths.

22 Nano-Plating Figure 1.12. A surface morphology change with increasing film thickness in electrolytic iron films grown on amorphous substrates at (a) 500, (b) 1000, and (c) 2000 A/m2.


23

Figure 1.13. A mechanism for the development of surface irregularities in plated films with increasing film thickness during electrodeposition.

24

Nano-Plating

occur uniformly, but takes place preferentially at the pointed sites (protrusions) of the substrate. Enhanced discharge events at these protrusions induce heating, which activates the electrode reaction by increasing the supply of anions and the emission of cations. At such activated sites, more atoms will deposit, thus locally increasing the film thickness. In the valley region, where metal ions are deficient, no significant discharge occurs and thus no film growth takes place. Such a localized metal ion discharge will become more pronounced with time and will promote surface roughening further. We consider that the above picture describes a formation mechanism of surface roughness with increasing film thickness in plated films. A dotted line in Figure 1.13(b)– (d) represents MIDL. It has been thought that metal ions are supplied through the so-called diffusion layer. We have a different theory on the supply of metal ions. We consider that the cathode surface consists of two distinct regions; a region that metal ions are supplied (i.e. the protrusions) and a region that metal ions are not supplied (i.e. the valley region). Film growth at the valley region occurs only by the overflow of atoms from the protrusion. The overflow distance is equivalent to the average distance traveled by adatoms via surface diffusion. This distance determines the shape of the pointed region; the shorter the diffusion distance, the more pointed the region is. Conversely, the longer the diffusion distance, the more rounded the pointed region is. The long diffusion distance allows the crystal to minimize the surface energy by converting a pointed region into a thermally equilibrated rounded shape. Surface roughness is finally determined by the size, shape, and distribution of these protrusions. For low melting point metals (Figures 1.5 –1.8), each protrusion is made up of one grain (see Figure 1.14), whereas for high melting point metals (Figures 1.9– 1.12), each protrusion is not necessarily one grain but may be an assembly of small grains (see Figure 1.4). For plated alloy films, polycrystalline grains determine the surface irregularities in most cases. The magnitude of surface irregularities can be expressed in terms of wavelength or height. For example, if the magnitude is larger than the wavelength of visible light, the films will cause random scattering to the light and thus appear semi-bright or dull. Conversely, films having surface irregularities, whose magnitudes are less than the wavelength of visible light, will appear bright. 1.4.2.2 Surface morphology change with current density (overpotential). The surface roughness of plated films is said to change with the overpotential (Haruyama, 1963; Ohno, 1988; Ohno and Haruyama, 1991; Winand, 1994). Indeed, the surface roughness is high at low overpotentials and becomes small at high overpotentials. Low-overpotential (low current density) depositions produce films with large surface irregularities, whereas high-overpotential depositions yield films with smooth surfaces. This trend was observed in electrodeposited tin films, whose thickness was below 2.6 mm, in nickel films (Figure 1.10), in cobalt films (Figure 1.15), in cadmium films


25

Figure 1.14. A schematic diagram showing how the surface morphology and crystal size of low-melting point metals evolve with a change in current density during electrodeposition.

(Figure 1.16), and in silver films (Figure 1.17). It is believed that the change in surface roughness is affected by a change in the overpotential. In this book we propose an entirely different mechanism for the development of surface roughness in plated films. We show schematically how the surface roughness changes with

26

Nano-Plating

Figure 1.15. A surface morphology change with increasing current density in 5 mm-thick electrolytic cobalt films grown on amorphous substrates from (A) sulfate and (B) chloride baths.


27

Figure 1.16. A surface morphology change with increasing current density in electrolytic cadmium films grown on amorphous substrates from (A) sulfate and (B) chloride baths. The film thickness is 6.73 mm, pH is 1.8, and the solution temperature is 25 8C.

28 Nano-Plating Figure 1.17. A surface morphology change with increasing current density in electrolytic silver films grown on amorphous substrates. The film thickness is 10 mm, pH is 2, and the solution temperature is 60 8C.


29

current density in Figure 1.18. In Figure 1.18(a), the number of metal ion discharge event per unit area or per unit time is shown to be small at low current densities. Even at low current densities, the metal ion discharge event will still occur preferentially at protrusions. These protrusions will grow, producing films with high surface irregularities. At high current densities (Figure 1.18(b) and (c)), high densities of metal ions (cations) are discharged on the cathode surface. Here it is important to remember that similarly charged particles repel each other if they are brought closer than a critical distance from each other. This is true for the case of metal depositions at high current densities. A high density of metal ions and electrons generated at high current densities will be redistributed over the surface according to the magnitude of their repulsive force. This cation redistribution over the substrate surface allows the discharge sites to be more uniformly distributed, making the surface of plated films smoother (Figure 1.18(b)) or wavier (Figure 1.18(c)). It can be concluded that one cause of surface roughness in plated films is the distribution of current densities over the cathode surface, but it cannot be attributed to a change in the overpotential. 1.4.2.3 Surface morphology change with the type of anions. Various metallic salts are used to prepare plating solutions and the selection of the salts is often determined by empirical rules. A metallic salt dissociates into metal ions and anions in a plating solution. Therefore, the effect of metallic salts on the film morphology is the same as the effect of the anions. For electrolytic cadmium films no morphological difference was observed between sulfate and chloride baths (see Figure 1.6). For electrolytic nickel (Figure 1.10) and electrolytic cobalt (Figures 1.11 and 1.15), a significant morphological difference was observed between sulfate and chloride baths. Here a smooth surface was obtained in a sulfate bath, whereas a rough surface with sharply pointed protrusions was obtained in a chloride bath. This difference was prominent at low current densities. For electrolytic zinc, the use of four kinds of baths and their mixed baths produced various characteristic surface morphologies (Watanabe and Minami, 2000) (see Figure 1.19). A mechanism as to how anions change the surface morphology appears to be complex and cannot be explained at the present time. There is a general theory (Winand, 1994) that supports an anomalous adsorption of anions on the cathode. According to our studies, the surface morphology is affected by the molecular weight and size of anions, which are related to the solution viscosity and the diffusivity of anions toward the anode. A plating solution containing anions of large molecular size, such as sulfate ions, is highly viscous and produces wavy surfaces. A solution containing small molecules, such as chloride ions, has low viscosity and produces a deposit with rough surface. If the molecular size of anions is large, the movement of cations toward the cathode and anions

30

Nano-Plating

Figure 1.18. A schematic illustration showing how a change in current density affects the surface irregularity of plated films.

toward the anode is easily interfered with. This interference will disperse the discharge event over the substrate surface, thus producing films with a smooth surface. For anions with small molecular size, the interference between the cations and anions is small, suggesting that the cation discharge process will occur without the dispersion of cations. In this way, the discharge will be concentrated on protrusions, thus producing rough surfaces.


31

Figure 1.19. A surface morphology change with increasing film thickness in electrolytic zinc films grown on amorphous substrates from (A) chloride, (B) chloride þ sulfate, (C) sulfate, (D) acetic, and (E) sulfate baths.

32

Nano-Plating

Watts’ baths containing both sulfate and chloride salts have been used for plating nickel. Nickel films with a smooth surface can be obtained from a sulfate bath, but the current efficiency is relatively low and the bath tends to produce pitted nickel films. In contrast, the chloride bath has a high current efficiency but produces nickel films with rough surfaces. To improve the current efficiency, while reducing the density of pits and smoothing the surface, both sulfate and chloride salts were mixed into one bath—this is the Watts bath. 1.4.2.4 Surface morphology change with solution temperature. The effect of solution temperature on surface morphology is shown in Figure 1.20 for electrolytic silver films, and in Figure 1.21 for electrolytic nickel films. The solution temperature is clearly one of the most important factors controlling the surface morphology. High solution temperature generally produces higher surface irregularities. Electrolytic silver films, which have a low melting point temperature, become large-grained with increasing solution temperature (see Figure 1.20). An increase in the grain size with increasing solution temperature is closely related to the improved supply rate of metal ions, the increased diffusion distance of cations, and the increased surface diffusion distance of adatoms. It can be easily seen that the increased diffusivity allows adatoms to migrate a long distance over the substrate surface, thus producing large grains. 1.4.2.5 Surface morphology change with solution agitation. Figure 1.22 illustrates the effect of solution agitation on the surface morphology of electrolytic nickel films obtained from a chloride bath. Agitation was performed by varying the speed of the magnetic stirrer. Nickel films from a non-agitated bath show severe surface irregularities (see Figure 1.22(a)). An increase in stirring speed caused the film surface to become even rougher and secondary surface irregularities were developed on the side face of the primary irregularities (see Figure 1.22(b)– (d)). The development of the secondary irregularities occurred because the solution agitation supplied metal ions to the side face and thus made metal ion discharge possible at such sites. 1.4.2.6 Surface morphology of alloy films. For alloy films, the microstructure (phase structure) is one of the most significant factors affecting the surface morphology. In general, the finer the grains, the smoother the surface. There are, however, some exceptions to this trend. These exceptions will be described in the following section. For most of the amorphous films, the surface is smooth and bright. Figure 1.23 shows the surface morphology and fracture surface of an electrolytic Ni –P alloy film. Contrary to crystalline films, the amorphous films do not have crystallographic facets and form a smooth surface in order to minimize surface energy. Amorphous films maintain their surface smoothness even with increasing film thickness.


Figure 1.20. A crystal size change with increasing solution temperature in electrolytic silver films grown on amorphous substrates. The film thickness is 10 mm, the current density is 150 A/m2, and pH is 2.

33

34 Nano-Plating Figure 1.21. A surface morphology change with increasing solution temperature in electrolytic nickel films grown from (A) sulfate, (B) Watts, and (C) chloride baths.


Figure 1.22. Effect of solution agitation on the surface morphology of electrolytic nickel films grown on amorphous substrates from a chloride bath. The solution agitation was achieved by changing the rotation speed of a magnetic stirrer.

35

36

Nano-Plating

Figure 1.23. The surface morphology and fracture surface of an electrolytic Ni–P amorphous alloy film.

1.4.2.7 Summary of the formation principles of surface irregularities. Various formation mechanisms of the surface morphology in plated films have been described above, and it has been shown that the surface morphology of pure metal deposits have to be discussed separately from those of alloy deposits. As will be described later, the grain size of low melting point metals is large, whereas that of high melting point metals is small. For low melting point metals, the surface roughness is proportional to the grain size. For high melting point metals and alloy films, the primary factor in the development of surface roughness is the current density distribution over the surface. Protrusions form at sites where the metal ion discharge is concentrated. The size and distribution of these protrusions thus formed determine the final surface roughness. The type of anions changes the shape of the protrusions. Anions generally change the wavelength of the surface roughness. The size of the protrusions is also affected by the solution temperature. An increase in the solution temperature makes the protrusions large and vice versa. Solution agitation also affects the surface morphology. For pure metals, the grain size can be a direct indication of the surface roughness, but there are some exceptions to this trend. For alloy deposits, it is often the case that the surface roughness is not related to the grain size.


37

1.4.2.8 Formation of dendrites (Idemoto, 1999). Figure 1.24 shows dendrites in an electrolytic gold film. Dendrites in plated films are often generated under conditions of low metal concentration/high current densities or of low current densities. An individual dendrite is generally a single crystal, but in some cases, it may be an assembly of fine grains. Even in plating systems that form dendrites, the films always start as a uniform layer and then a dendrite forms on top of this layer. This is consistent with the fact that metal ions present in close contact with the cathode are initially distributed

Figure 1.24. Dendrites formed on the surface of electrolytic gold films grown on (a) {100} and (b) {110} copper substrates.

38

Nano-Plating

uniformly over the substrate surface. After the layer formation, a metal-ion denuted zone (MIDL) is created over the surface. The thickness variation of the MIDL will then develop over the surface. Consequently, metal ions will be supplied preferentially through a thin MIDL region, where the discharge process occurs locally. Dendrites will then nucleate at such regions and their growth will be further accelerated through the tips, where the current density becomes very high. Furthermore, the temperature of the tip region is expected to rise due to the enhanced discharge activity (see Figure 2.6).

1.4.2.9 Effect of brighteners. Various organic or inorganic compounds have been added as additives to plating solutions to improve the brightness of electrodeposited films. Such additives are known as leveling agents or brighteners. There are a number of studies involving leveling agents/brighteners in plated films. A detailed review on this subject can be found in the article by Onical and Muresan (1991). These studies include a thermodynamic analysis of the adsorption mechanism of brighteners (Minami and Mayanma, 1993), AFM observations of brighteners (Schmidt et al., 1996), and a discussion on an adsorption mechanism of organic additives (Nichols et al., 1993). In the plating industries, two types of brighteners, known as Type I and Type II, are used. There is no clear definition for this classification. Organic reagents with large molecular weight are generally used as Type I additives, which are primarily growth inhibitors. Type I additives are adsorbed preferentially at the protrusions. The adsorbed additives then inhibit the growth of the protrusions and promote the growth at the valley regions. Such a growth process promotes smoothening of the film surface. It should be noted that these brighteners are negatively charged and thus their adsorption on a negatively charged substrate is considered anomalous. We consider that these brighteners are not adsorbed on the cathode. Similar to the role of anions on the surface morphology of plated films, their function is to affect the surface morphology by changing the transport property of various chemical species in a solution. It is convenient to consider Type I brighteners as anions having a much larger molecular dimension. A surface smoothening mechanism of Type I brighteners is a leveling effect but not a brightening effect. This leveling effect is to hide fine surface irregularities or scratches on the film surface. For this reason, Type I brighteners should be called levelers. Type II brighteners are generally inorganic reagents with small molecular weights. These additives decompose into light atomic elements, such as B, C, N, P, S, As, etc. at the cathode during electrodeposition. The decomposed products are then incorporated primarily into the grain boundaries of the deposits. This incorporation makes fine-grained deposits and generates bright film surfaces. Generally, atomic elements having a strong chemical bond to the plated metal tend to be incorporated into the deposit. Although these atoms are included into the film, they will be incorporated at the grain boundaries without


39

forming an ordered structure in the matrix. As the amount of incorporated atoms increases, the grain size decreases. Therefore, the function of Type II additives is to refine the grain by alloying. If the amount of these light elements becomes extremely large, plated films sometimes become amorphous. Since Type II additives increase the brightness, they are called brighteners. In practice, both Type I and Type II brighteners are used together in a plating solution to obtain deposits with smooth surfaces. 1.4.3 Grain size (grain size, granular, spherical, columnar, needle-like, etc.) 1.4.3.1 Grain size of pure metal deposits. It is generally accepted that the grain size of plated films can be controlled by the overpotential (Vetter, 1961; Haruyama, 1963; Ohno, 1988; Ohno and Haruyama, 1991; Winand, 1994). The use of high overpotentials is suggested to obtain films with smooth surfaces. At high overpotentials, the deposition rate becomes faster than the film growth rate. An increase in the deposition rate will promote higher nucleation rates, which in turn will produce fine-grained films. If this logic is true, the grain size should be small for films grown at high overpotentials. Figure 1.25 is planview TEM micrograph of electrolytic nickel films stripped off the substrate. The surface morphology of these films was presented in Figure 1.10. The microstructure of these films is considered to be columnar. According to this experiment, films with rough surfaces were obtained at low overpotentials, whereas those with smooth surfaces were produced at high overpotentials. In Figure 1.25, the grain size appears to be the same at the same film depth (thickness) regardless of the type of plating solutions used and the magnitude of current densities. In other words, the grain size does not depend on plating conditions (type of anions and current density). From these observations, it is correct to state, “High overpotentials should be used to obtain electrodeposits with smooth surfaces.” However, it is not correct to say, “Such surface smoothening resulted from the formation of fine grains.” The magnitude of surface irregularities varies with various plating conditions, such as current densities. If we assume that the surface irregularities are represented by the grain size, we can estimate grain sizes for pure metal deposits, as shown in Table 1.1. These estimated grain size values for 5 – mm thick films are obtained from Figures 1.5– 1.12. In Table 1.1, the melting point of each metal is also listed, in descending order. From the table, it is clear that the higher the melting point, the smaller the grain size and vice versa. As described above, heated metal atoms undergo a cooling process immediately after the discharge process, while diffusing over the film surface, and finally contribute to crystal growth. The diffusivity of high melting point metals is small or, more precisely, their average diffusion distance is short. In other words, high melting point metals tend to freeze faster than low melting point metals. The cooling rate determines the extent of the grain growth and the resulting grain size of plated films. According to experiments by Girin (1995), the grain size of plated films is related to the exchange current density,

40 Nano-Plating Figure 1.25. Plan-view TEM micrographs showing the grain-size distribution of electrolytic nickel films observed at three different depths (0.12, 1.3, and 8.0 mm) along the film thickness direction. The nickel films were grown at 100 and 1000 A/m2 from (A) sulfate and (B) chloride baths.

41

Microstructure Control Theory of Plated Film Table 1.1. Relationship between the melting point temperature and the crystal size of plated films. Metal

Crystal size at 5 mm thickness position (mm)

Melting point (8C)

Fe Co Ni Cu Ag Zn Cd Sn

? About 0.5 About 0.5 About 0.5 About 3 About 2 About 8 About 10

1537 1495 1453 1063 960 420 321 232

The crystal size was measured for 5-mm thick films based on an assumption that the surface irregularities seen in Figures 1.5–1.12 represent the crystal size.

which in turn is related to the melting point. Large-grained deposits are obtained for low melting point metals and vice versa. 1.4.3.2 Grain size of alloy deposits. The grain size of alloy films changes with the alloy composition. In particular, the incorporation of light elements, such as B, C, P, and S, has a dramatic grain-refining effect. Other binary alloy systems such as eutectic or peritectic alloys (Figure 1.3(b)) are of phase separation type and crystallize into two types of crystals where their grain size is determined by the composition. As an example, we show the surface morphology and grain size of electrolytic Fe –Ni alloy films as a function of current densities (overpotential) (Fukumuro et al., 1996) in Figure 1.26. The film composition did not change when the current density was varied while keeping the solution composition constant. However, it was possible to vary the film composition by changing the solution composition (Fukumuro et al., 1998). In Figure 1.26, we show SEM micrographs, cross-section TEM images and electron diffraction patterns for electrolytic Fe – Ni alloy films. Although the current density is increased 10-fold from 150 to 1450 A/m2 (the overpotential is 5 V), the alloy composition remained constant (Fe-48 at.% Ni) (compare Figure 1.26(a) with (b)). For these films, high current densities produced a smooth surface, whereas low current densities produced a rough surface. This trend is indeed consistent with the OT. However, contrary to the prediction made by the OT, the grain size for the films plated at 150 and 1450 A/m2 is almost the same. The films shown in Figure 1.26(c) and (d) are also deposited at the current densities of 150 and 1450 A/m2, respectively, and have a composition of about Fe –80 at.% Ni. Similar to the case of the Fe– 48 at.% Ni films, high current densities produced a smooth surface, whereas low current densities induced a rough surface. The grain size for the films plated at 150 and 1450 A/m2 is almost the same as that plated at 1450 A/m2.

42

Nano-Plating

Figure 1.26. Effect of film composition, overpotential, and current density on the surface morphology and crystal size of electrolytic Fe–Ni alloy films.


43

From this experiment, it can be concluded that surface irregularities change with the current density (overpotential) but that the grain size is determined by the film composition, independently of the overpotential. In Figure 1.27(a), we show the equilibrium phase diagram of Fe– Ni alloys and the grain size of plated Fe –Ni alloy films as a function of the alloy composition (Fukumuro et al., 1998) (see Figure 1.27(b)). Although both the Fe-rich and Ni-rich films have large grains, Fe –Ni films having mixed g- and a-phases possess the smallest grains. In general, for eutectic or peritectic alloy films, the higher melting point phase becomes fine-grained,

Figure 1.27. Comparison between (a) the equilibrium phase diagram of Fe–Ni alloys and (b) the phase/grain size of electroplated Fe–Ni alloy films.

44

Nano-Plating

whereas the lower melting point phase becomes large-grained. Figure 1.28 is a TEM micrograph showing the structure of an electrolytic film of Ag – Co, which is mutually insoluble. Spherically shaped Ag grains are surrounded by the fine-grained Co matrix. Cobalt has a higher melting point than silver metal and thus becomes fine-grained during rapid cooling. Conversely, silver crystals with a low melting point, can grow slowly and form crystals with a spherical shape, which is the configuration for the minimum surface energy. Figure 1.29 is a TEM micrograph for crystalline, crystalline/amorphous, and amorphous films of electrolytic chromium (Furuya et al., 1981; Furuya and Tanabe, 1982). Amorphous chromium films can be obtained under plating conditions of low current efficiency, whereby a large amount of hydrogen is incorporated into the deposits.

Figure 1.28. The microstructure of an electroplated Ag– Co alloy film and its electron diffraction patterns.

Microstructure Control Theory of Plated Film Figure 1.29. TEM images/electron diffraction patterns for (a) crystalline Cr, (b) partially crystalline Cr–H, and (c) amorphous Cr–H films obtained by electrodeposition.

45

46

Nano-Plating

Hydrogen is known to refine the grain size; plated chromium films containing hydrogen are generally fine-grained and sometimes become amorphous in extreme cases (see Figure 1.29(c)). This is due to the strong Cr –H bond, which suppressed the growth of Cr crystals. This phenomenon is similar to the formation of amorphous Si –H. There is a report (Loebichir et al., 1980) that plated nickel films, known as hard nickel, can become amorphous by alloying with hydrogen. Alloying of plated films with hydrogen may be an important subject in future (Raub, 1930). We can summarize the relationship between the grain size and the surface morphology of plated films in Figure 1.30. Pure metal deposits tend to grow in columns. It should be remembered that even large-grained films do not necessarily have rough surfaces but can have smooth or wavy surface, depending on the type of anions and additives. In the case of high melting point metals and alloys, even fine grains can produce either rough or smooth surfaces. Therefore, special care should be taken if one discusses the grain size from the surface morphology alone. 1.4.4 Preferred orientation (texture) (film normal/film plane direction and assembled structure) As described previously, plated films are an assembly of fine crystals. These crystals arrange themselves in an organized manner along the films crystallographic normal direction (see Figure 1.31(a)). This phenomenon is called a preferred orientation (or a texture). In Figure 1.31(b), we illustrate the case of a non-textured film. There are studies that involve the development of highly functional films by manipulating the texture (Mackinnon et al., 1987; Ng and Ling, 1990; Ye et al., 1992; Fujisawa et al., 1993; Kang et al., 1995; Tajiri et al., 1998; Tajiri, 2000). Pangarov (1965) developed a theory of preferred orientation in various crystal systems in 1965 and claimed that the texture is controlled by the overpotential alone. His theory has been universally accepted in the thin films community. As discussed later, however, the film texture cannot be theorized as easily as has been done by Pangarov (1965) because it depends not only on the overpotential (current density), but also on the type of anions, plating temperature, and even film thickness. Films obtained by vapor deposition or sputtering exhibit a texture phenomenon as well, but its formation mechanism is not well understood. In general, the surface energy varies with the type of crystallographic planes, which are expressed in terms of the Miller indices. For example, in fcc crystals, the surface energy decreases in the order of 110 . 100 . 111; and for bcc crystals, 111 . 100 . 110 (Figure 1.32). A plane with the lowest surface energy generally becomes the final growth facet seen on the top surface of the film (Reddy, 1963). It is, however, important to note that the lowest surface-energy plane does not necessarily become the plane of the texture. Figure 1.33(a) is a SEM micrograph of an electrolytic gold film showing a number of pyramidal crystals. All


Figure 1.30. Relationship between crystal size and surface morphology in plated films.

47

48

Nano-Plating

Figure 1.31. A schematic diagram showing (a) non-textured and (b) ,111 . -textured plated films grown on amorphous substrates.

the triangular faces bounding the pyramids are the {111} planes, which are the lowest surface-energy plane in fcc crystals. The film normal direction (the texture), however, is the , 100 . direction (see Figure 1.33(b)). In other words, the gold crystal has grown along the , 100 . direction by the combined growth of 4 {111} planes. In this case, the preferred growth plane and direction of a gold crystal are the {111} and , 111 . , respectively, but the direction (preferred orientation) normal to the substrate is , 100 . . This example


49

Figure 1.32. The cross-sectional structure of an electrolytic Ni –P alloy deposit. (a) The cross-section TEM micrograph, (b) the electron diffraction pattern, and (c) the X-ray diffraction pattern.

Figure 1.33. Surface morphology in an electrolytic {100} gold film. (a) The SEM micrograph of the surface and (b) a schematic diagram showing a cross-sectional view of the growth direction and the facets.

50

Nano-Plating

demonstrates that the texture of a crystal may be different from the preferred growth direction. The texture also changes with various plating conditions. We will describe the definition of film texture and experimental methods for determining the texture in the Section 6.5 of Chapter 6 (Willson and Rogers, 1964). 1.4.4.1 Experimental methods for determining the texture of plated films. Here, we will discuss our experimental results that demonstrate the relationship between the texture of various pure metal films and plating conditions. To investigate the cause of texture development in pure metal deposits, it is necessary to simplify the plating conditions. For example, a simple addition of additives to a plating solution will complicate the texture formation process of plated films, especially when the films becomes an alloy by incorporating some atomic elements from additives. Such an alloying is known to change the film texture. Therefore, the main purpose of the present study is to understand how the texture of pure metal deposits is affected by plating conditions. The surface roughness or the strain state of these films will not be considered here. Crystalline substrates should not be used for studying the texture phenomenon because the films generally grow epitaxially. On such substrates, it is difficult to obtain the intrinsic texture of plated films free of the substrate effect. For this reason, we used amorphous alloy foils as substrates, which will not affect the initial structure of plated films. The amorphous substrates used in this experiment are either electrolytic/electroless Ni– 25 at.% P alloy films grown on rolled copper sheets or splat-quenched Fe –Ni –Si –B alloy foils (Japan Amorphous Metal Co., Ltd.). 1.4.4.2 Relationship between the texture of various plated films and plating conditions. We have studied the relationship between the texture of plated films and plating conditions. These metals include Cu (Watanabe et al., to be submitted), Ni (Watanabe et al., to be submitted), Ag (Watanabe et al., 2002), Au (Inoue et al., 2002), and Pt (Terakado et al., 2001) for fcc crystals, Fe (Inoue et al., 2001) and Cr (Yoshioka et al., 1999) for bcc crystals, Cd (Miyazaki et al., 1999), Zn (Sasaki et al., 2001), Co (Koide et al., 2000), and Ru (Teshigawara et al., 2001) for hcp crystals and tetragonal Sn (Karayannis and Patermarakis, 1995). From these results, the factors affecting the texture of plated films were found to be: (1) the crystal system (the type of a metal); (2) the type of metallic salts (anions); (3) the current density (overpotential); (4) the film thickness; (5) the type of a buffer;1 (6) additives such as brighteners; and (7) a cation originating from a reagent added as the supporting electrolyte.2 1 The solution pH, i.e. the concentration of hydrogen ions, does not affect the texture directly but the type of anions from a reagent used as a buffer may influence the texture. 2 Most recently, a reagent has been added to a plating solution as a supporting electrolyte and was found to influence the film texture.


51

We have conducted a number of experiments on the texture of plated films as a function of plating conditions. Due to limited space here, we will present only a part of the results. For further details, Chapter 7 should be consulted. Figure 1.34(a)– (d) describe a change in the texture index of electrolytic copper films obtained from a pyrophosphate bath as a function of film thickness at the current densities of 100, 300, 500, and 1000 A/m2, respectively (Watanabe et al., to be submitted). All the thin copper films show the 111 texture for all current densities. The copper films obtained at low current densities also maintain the 111 texture, even with increasing current densities and film thickness. As the current density exceeds 300 A/m2, the texture changes to the 100 texture with increasing film thickness. Figure 1.35 shows the texture index for electrolytic copper films (obtained at two current densities (500 and 1000 A/m2) from a sulfate bath) plotted against film thickness. At low current densities, the films are seen to

Figure 1.34. A texture change with increasing film thickness and current density in electrolytic copper films grown on amorphous alloy substrates from a pyrophosphate bath. The solution temperature is 55 8C and pH is 8.7.

52

Nano-Plating

Figure 1.35. A texture change with increasing film thickness in electrolytic copper films grown on amorphous alloy substrates at the current densities of (a) 500 and (b) 1000 A/m2 from a sulfate bath. The solution temperature is 60 8C and pH is 2.0.

display the 110 texture. At high current density, no texture was initially present but the 110 texture developed with increasing film thickness. The texture of plated films thus changes with the current density and the type of a bath (the type of anions present). The effect of the solution temperature on the texture of the electrolytic copper films is seen in Figure 1.36. Similar to the case of Figure 1.35, two current densities (500 and 1000 A/m2) were used in this experiment. The 111 texture developed for both current densities but changed to the 110 texture with increasing solution temperature. Therefore, the solution temperature also affects the texture.

Figure 1.36. A texture change with increasing solution temperature in electrolytic copper films grown on amorphous alloy substrates at the current densities of (a) 500 and (b) 1000 A/m2 from a sulfate bath. The film thickness is 5 mm and pH is 2.0.


53

Graphs in Figure 1.37 show a change in the texture of electrolytic nickel films with increasing film thickness and current density (Watanabe et al., to be submitted). If the films are thin, they have either no texture or the 111 texture for all current densities. With increasing film thickness, however, the texture turns to 100 at low current densities and then to 111 at current densities larger than 500 A/m2. A change in the texture of electrolytic nickel films obtained from various plating solutions is plotted in Figure 1.38 as a function of current density. Four types of baths were used to deposit nickel films: (a) a sulfate bath; (b) a chloride bath; (c) a Watts bath (9 parts sulfate bath and 1 part chloride bath); and (d) a sulfamate bath. Nickel films from sulfate and chloride baths yielded the 110 and 311 texture, respectively at high current densities. The mixed sulfate– chloride bath, i.e. the Watts bath, yielded the 110 texture, consistent with the fact that the high mixing ratio of a sulfate electrolyte determined the resulting texture. Contrary to the case of the sulfate or chloride bath, nickel films from a sulfamate bath have the 100 texture.

Figure 1.37. A texture change with increasing film thickness in electrolytic nickel films grown on amorphous alloy substrates at the current densities of (a) 100, (b) 500, (c) 1000, and (d) 2000 A/m2 from a sulfate bath. The solution temperature is 40 8C and pH is 2.8.

54

Nano-Plating

Figure 1.38. A texture change with increasing current density in electrolytic nickel films grown on amorphous alloy substrates from (a) sulfate, (b) chloride, (c) Watts, and (d) sulfamate baths. The solution temperature is 40 8C and pH is 2.8.

Similar to the copper deposits described above, the texture of electrolytic nickel films also changes with the type of anions used. The effect of solution temperature on the texture of electrolytic nickel films from sulfate, Watts, and chloride baths is shown in Figure 1.39. The current density used for this experiment was 1500 A/m2. The texture of the films from a sulfate bath is 110 at low temperatures but changed to 100 at high temperatures. However, the texture from a chloride bath, is 311 at low temperatures, but changed to 110 at high temperatures. The mixed sulfate– chloride bath of Watts type starts with the 110 texture at low temperatures, followed by the 100 texture at the temperature range of 40 – 80 8C, and finally came back to the 110 at high temperatures. The observed texture change in nickel films with solution temperature is the same as that of copper deposits described above. Figure 1.40 shows the change in the texture of electrolytic cobalt films with current density (Sasaki et al., 2001). Figure 1.40(A) and (B) represents cobalt films from sulfate and chloride baths, respectively. The texture from both baths was 100 at low current


Figure 1.39. A texture change with increasing solution temperature in electrolytic nickel films grown on amorphous alloy substrates from (A) sulfate, (B) Watts, and (C) chloride baths. The current density is 1500 A/m2 and pH is 2.8.

55

56

Nano-Plating

Figure 1.40. A texture change with increasing current density in 5 mm-thick electrolytic cobalt films grown on amorphous alloy substrates from (A) sulfate and (B) chloride baths. The solution temperature is 30 8C and pH is 2.0.

densities and 110 at high current densities, except that the 100 texture at low temperatures is stronger for the chloride bath than that from the sulfate bath. In Figure 1.41, we show a change in the texture of nickel films as a function of solution temperature and the type of a bath (sulfate- and chloride-type). For both baths, the solution temperature affected the texture markedly; the texture was 110 at low temperatures and 100 at high temperatures. Similar to the case of electrolytic copper and nickel films, the texture of electrolytic cobalt films changes with current density and solution temperature, but not with the type of a bath (anion type) used.

Figure 1.41. A texture change with increasing solution temperature in 5 mm-thick electrolytic cobalt films grown on amorphous alloy substrates from (A) sulfate and (B) chloride baths. pH is 2.0.


57

A change in the texture of electrolytic gold films obtained from a cyanide bath is seen as a function of film thickness in Figure 1.42 (Inoue et al., 2002). The effect of current density and solution temperature on the texture is also included in Figure 1.42. For all plating conditions, the films initially showed no or 111 texture, which increased with increasing film thickness. The thick films eventually exhibit the 111 texture for all plating conditions. The texture of gold films obtained from a sulfite bath is seen in Figure 1.43. The plating condition were the same as those of the cyanide bath (see Figure 1.42). The initial films from a sulfite bath take the 111 texture similar to that from a cyanide bath. The texture change, however, becomes very complex with increasing film thickness and no systematic trend can be detected. An increase in the film thickness causes the texture to become 111, 100, or 110. In conclusion, the texture of electrolytic gold films varies with the type of anions and the solution temperature. 1.4.4.3 Effect of plating conditions on the film texture. We have described experimental results on the texture changes occurring under various plating conditions for 12 metals (Miyazaki et al., 1999; Yoshioka et al., 1999; Koide et al., 2000; Inoue et al., 2001, 2002; Terakado et al., 2001; Sasaki et al., 2001; Teshigawara et al., 2001; Watanabe et al., 2002; Watanabe et al., to be submitted). These results are summarized in Tables 1.1 –1.3. Here, we list the results according to the crystal system of plated films in the increasing order of melting point. For hcp crystals, although accidental, the c/a ratio is seen to decrease with increasing melting point. The table also shows a qualitative trend as to how the texture changes with the type of bath, film thickness, solution temperature, and current density. A horizontal arrow indicates a gradual change in the texture with increasing film thickness, solution temperature, or current density. A texture value containing no arrow means no texture change. To experimentally elucidate a systematic trend in the texture of plated films, it is important to compare all the films using identical anions. The problem is that not all metals can be plated with the same type of metal salts (or anions). For this reason, the underlined bath in the table indicates a bath containing other types of metal salt. The texture of fcc metals (Ag, Au, Cu, Cu, Ni, and Pt) varies with the type of a bath and plating conditions in a complicated manner, and thus it is difficult to identify a systematic trend in the texture of fcc metals. For bcc metals (Fe and Cr), the texture of iron films is the same regardless of the type of bath used, but changes in a complex fashion in chromium films. It is interesting to note that for hcp metals (Cd, Zn, Co, and Ru) the texture type is clearly separated by the value of the c/a ratio. If the shape of an atom in the hcp structure is perfectly spherical, the c/a value is 1.633. The c/a values for cadmium and zinc are 1.979 and 1.856, respectively. These metals are a hexagonal crystal that is stretched along the , 001 . direction (the c-direction). The shape of these atoms is considered to be elliptical with its long axis along the c-direction. The texture of these metals is , 001 . . The c/a

58 Nano-Plating Figure 1.42. A texture change with increasing film thickness, solution temperature, and current density in electrolytic gold films grown on amorphous alloy substrates from a cyanide bath.


59

Figure 1.43. A texture change with increasing film thickness, solution temperature, and current density in electrolytic gold films grown on amorphous alloy substrates from a sulfite bath.

60

Table 1.2. Epitaxial relationships between various substrate materials and plated films grown by different plating methods. Plated film/substrate

Crystallographic relationship

Displacement plating

Cu(fcc)/Fe(bcc)


Au(fcc)/Fe(fcc)

Electroplating

Au(fcc)/Fe(bcc)

Electroless plating (NaPH2O2) Contact plating

Au(fcc)/Fe(bcc) Au(fcc)/Fe(bec)


Ag(fcc)/Cu(fcc)


Au(fcc)/Cu(fcc)

Electroless plating (NaBH4)

Ni(fcc)/Cu(fcc)

(001)Cu//(001)Fe, (101)Cu//(101)Fe, (111)Cu//(111)Fe [100]Cu//[110]Fe, [011]Cu///[010]Fe, [011]Cu//[121]Fe (001)Au//(001)Fe, (110)Au//(110)Fe [100]Au//[110]Fe, [110]Au//[001]Fe (001)Au//(001)Fe, (110)Au//(110)Fe [010]Au//[110]Fe, [110]Au//[001]Fe (001)Au//(001)Fe, (110)Au//(110)Fe [010]Au//[110]Fe, [110]Au//[001]Fe (001)Au//(001)Fe, (110)Au//(110)Fe [010]Au//[110]Fe, [110]Au//[001]Fe (001)Ag//(001)Cu, (110)Ag//(110)Cu, (111)Ag//(111)Cu [100]Ag//[100]Cu, [110]Ag//[110]Cu, [110]Ag//[110]Cu (001)Au//(001)Cu, (011)Au//(011)Cu, (111)Au//(111)Cu [100]Au//[100]Cu, [011]Au//[011]Cu, [011]Au//[011]Cu and/or (122)Au//(111)Cu, (112)Au//(111)Cu, (112)Au//(111)Cu [100]Au//[101]Cu, [111]Au//[110]Cu, [111]Au//[011]Cu and/or (011)Au//(111)Cu [100]Au//[121]Cu (100)Ni//(100)Cu, (110)Ni//(110)Cu, (111)Ni//(111)Cu [020]Ni//[020]Cu, [020]Ni//[020]Cu, [220]Ni//[220]Cu

Nano-Plating

Plating method

˚ Au Pd(fcc) alloy/Cu(fcc) at a ¼ 4.05A

Electroless plating (NaBH4) Electroplating

Ni(fcc)/Fe(bcc) Cr(bcc)/Cu(fcc)

Electroplating

Fe(bcc)/Au(fcc)

Electroplating

Cu(fcc)Au(fcc)

Electroplating

Co(hcp)Au(fcc)

(001) Au-Pd//(001)Cu, (011)Au-Pd//(111)Cu, (111)Au-Pd//(111)Cu [100]Au-Pd//[100]Cu, [100]Au-Pd//[100]Cu, [011]Au-Pd//[011]Cu and/or (112)Au-Pd//(111)Cu, (121)Au-Pd//(111)Cu, (211) Au-Pd//(111)Cu [110]Au-Pd//[110]Cu, [101]Au-Pd//[101]Cu, [011]Au-pd[011]Cu (101)Ni//(001)Fe [111]Ni//[110]Fe, or [111]Ni//[110]Fe, (misfit twin crystal) (110)Cr//(001)Cu [112]Cr//[220]Cu, or [112]Cr//[220]Cu (misfit twin crystal) (100)Fe//(100)Au, (110)Fe//(110)Au, (110)Fe//(111)Au [011]Fe//[001]Au, [110]Fe//[001]Au, [001]Fe//[110]Au (100)Cu//(100)Au [001] Cu/[001]Au (2110)Cu//(100)Au [0002]Cu//[002]Au


Electroplating

61

62 Table 1.3. Texture and plating conditions for various pure metal electrodeposits (Raub, 1930; Pangarov, 1965; Sard et al., 1966; Eichkorn and Fischer, 1967; Watanabe, 1989a,b, 1990, 1992, 1995, 1998, 2001; Watanabe and Masumoto, 1990). Increase of plating condition

Ag (M.P. 960 8C)

Sulfate bath

Film thickness Temperature Current density

Depend on current density No ) 110 110 ) 111

Film thickness

Hydentoic bath 111 ) 100

Cyanide bath 111

111 ) 100

110 ) 111 Higher temp. 100 Lower temp. 110 ) 111

111 111

No ) 100 111 ) 100

Perchloric acid bath

Temperature Current density

Chloride bath


Sulfamate bath


Au (M.P. 1063 8C)

Cu (M.P. 1083 8C)

Ni (M.P. 1453 8C)

Sulfite bath No ) 110

110

Lower CD 111 ) 100 Higher CD No ) 110 110 ) 100 Lower Temp. 100 ) 110 Higher Temp. 110 ) 100

110

No ) 110 110 ) No

Pt (M.P. 1759 8C)

Alkaline bath Pt(NH3)2(NO2)2 100 100

Lower CD 111 ) 110 Higher CD 111 ) 311 311 ) 110 Lower Temp. No ) 311 Higher Temp. 110 Lower CD 100 Higher CD No ) 100 110 ) 100 Lower Temp. 100 ) 110 Higher Temp. 100

Acid bath (H2PtCl5) 111 111

Nano-Plating

fcc

Increase of plating condition

Fe (1537 8C)

Cr (1875 8C)

Sulfate bath

Film thickness Temperature Current density Film thickness Temperature Current density Film thickness Temperature Current density

No ) 211

No ) 211 110 ) 111 111

hcp

Increase of plating condition

Cd (M.P. 321 8C) c/a ¼ 1.979

Sulfate bath

Film thickness Temperature

No ) 001 Higher C.D. 001 Lower C.D. 001 ) 102 No ) 001 No ) 001 Higher C.D. ? Lower C.D. 001 ) 102 No ) 001

Chloride bath

Sulfamate bath

Chloride bath

Nitric acid bath

Acetic acid bath

Current density Film thickness Temperature Current density Film thickness Temperature Current density Film thickness Temperature Current density

211 No ) 211 211 No ) 211 211

Zn (M.P. 420 8C) c/a ¼ 1.856

Co (M.P. 1495 8C) c/a ¼ 1.6228

Ru (M.P. 2500 8C) c/a ¼ 1.615

100 ) 110 110 ) 100 No ) 001 001

100 ) 110 100 ) 110 110 ) 100

100 ) No 100 ) No

No ) 001 001

100 ) 110

100


bcc

No 001

63

64

Nano-Plating

values for cobalt and ruthenium are 1.622 and 1.615, respectively, making them a hexagonal crystal compressed along the c-direction. These metals (Co and Ru) take the 100 or 110 texture, in which the c-axis is parallel to the substrate plane. A plane having the wide spacing parallel to the substrate surface appears to become the texture in hcp metals. In the case of fcc metals, initially formed polycrystalline films display no texture or develop the 111 texture, indicating that the {111} plane, the most stable plane in fcc crystals, is parallel to the substrate surface. The 111 texture sometimes changes with increasing film thickness either by the preferential growth of non-111 grains or by the formation of new nuclei. It is not clear at the present time why the texture of plated films changes with the type of bath and solution temperature. It is possible that the adsorption/desorption of various chemical species present in a plating solution on the film surface is connected to the texture formation (Willson and Rogers, 1964; Nyung Lee and Ye, 1981; Yan-Ling et al., 1991; Karayannis and Patermarakis, 1995). Since the degree of adsorption/desorption is strongly affected by solution temperature, the observed dependency of the texture on the solution temperature can be explained. Furthermore, the effectiveness of chemical species on the texture can be determined by the degree of their preferential adsorption/desorption on the specific planes of plated metals. Karayannis and Patermarakis (1995) reported that with increasing current density the texture of electrolytic nickel films obtained from a sulfate bath varied from 110, then to 100, and finally to 211, whereas that from a chloride bath changed from 110 to 211. Although their result is quite different from ours, they described that the texture changed not only with the overpotential (current density), but also with the type of anions present in the solution. Nyung Lee and Ye (1981) and Shu et al. (1991) also reported that a change in the texture is due to adsorption of additives. We studied the effect of additives on the texture of plated films and obtained the following results. The effect of additives on the texture of electrolytic zinc films is shown in Figure 1.44 (Watanabe and Minami, 2000). A change in the X-ray diffraction patterns with increasing film thickness is also shown. An X-ray diffraction pattern in Figure 1.44(0) is taken from randomly oriented zinc powder. This pattern is called the standard diffraction pattern. A diffraction pattern in Figure 1.44(A), which displays a strong 001 texture, was taken from an electrolytic zinc film deposited in a sulfate bath. The addition of a brightener, cresol sulfonic acid, to the bath changed the texture to 110 (cf. Figure 1.44(B)), while the addition of citric acid produced a film with a texture of 103 (cf. Figure 1.44(C)). The corresponding surface morphology is seen in Figure 1.45 and also changes with increasing film thickness. The surface of zinc films from an additive-free sulfate bath is seen to consist of many characteristic hexagonal crystals with stair steps, whose morphology was modified by the presence of many twins. The diffraction pattern of these films indicated the 001 texture. An inspection of the surface morphology in Figure 1.45(A), however, does not reveal


65

Figure 1.44. A texture change with a change in film thickness and in the type of additives for electrolytic zinc films grown at the current density of 1000 A/m2 from (A) sulfate, (B) sulfate þ cresol sulfonic acid, and (C) sulfate þ citric acid baths.

66

Nano-Plating

Figure 1.45. A surface morphology change with film thickness and the type of additives used in electrolytic zinc films grown on amorphous substrates from (A) sulfate, (B) sulfate þ 24 g/l cresol sulfonic acid, and (C) sulfate þ 120 g/l citric acid baths.

a hexagonal or triangular pattern indicative of 001 symmetry. A sulfate bath containing cresol sulfonic acid produced a surface morphology with fine surface irregularities, whereas that containing citric acid yielded a surface morphology similar to Figure 1.45(C). It should be remembered that the texture, 001, in Figure 1.45(C) is different from that, 103, in Figure 1.45(A). It is clear from these results that the texture does not necessarily agree with the surface morphology. There are some reports claiming that the texture of pulse-plated films can be changed by varying the shape of the pulse (Cavalotti et al., 1988; Ng and Ling, 1990; Tajiri et al., 1998). 1.4.4.4 Texture of electroless films. Electroless plating often uses NaH2PO2, NaBH4, or DMAB as a reducing agent. In such a plating solution, non-metallic elements, such as phosphorus or boron, are incorporated in the film. In this case, we have to treat them as binary alloy films of metal –P or metal –B. Here, we simplified our experiment by choosing


67

electroless plating systems, which produce pure metal deposits without incorporating non-metallic elements. Figure 1.46 shows the texture change of electroless gold films obtained from a cyanide bath containing NaBH4 as a reducing agent (Makino et al., 2001). In this figure, the texture

Figure 1.46. A texture change with increasing film thickness and solution temperature in electroless gold films on amorphous alloy substrates. The reducing agent ¼ NaBH4.

68

Nano-Plating

is plotted as a function of film thickness and solution temperature. Electroless gold films from a cyanide bath initially showed the strong 111 texture for all temperatures. With increasing film thickness, however, the 111 texture deteriorated and no texture was observed for films having a thickness of more than 7– 8 mm. The texture change of electroless copper films from a sulfate bath, which uses formaldehyde (HCHO) as a reducing agent, is presented as a function of film thickness and solution temperature in Figure 1.47 (Hosoiri and Watanabe, 2001). Similar to the case of electroless gold films presented above, electroless copper films also initially take the 111 texture but lose the texture at the thickness of several 10 mm. We are currently conducting experiments to determine the texture of electroless nickel films obtained from a chloride bath containing hydrazine chloride (N2H5Cl) as a reducing agent (Yoshihara et al., 2002). Like copper and gold, these nickel films initially exhibit a strong texture but lose this with increasing film thickness. From the above results, it can be concluded that electroless films with a fcc structure initially have a texture but lose this texture with increasing film thickness. The reason for the texture loss in the thick films is not well understood at the present time. 1.4.4.5 Summary. In order to clarify an underlying mechanism of texture formation in plated films, we have conducted a wide range of experiments on the texture of various plated metals. Although some systematic trends are found in electrolytic hcp and electroless fcc metals, no unified theory for the texture formation can be formulated. The initiation of the texture and its changes may be connected to the surface energy of metals and the adsorption/desorption of various chemical species (various anions and cations) on the growth planes. It is not clear, however, whether or not a detailed study on the adsorption/desorption phenomenon in one plated metal system can be applied to all other metals. It is obvious that more experiments have to be done to obtain better statistics by further improving the matrix listed in Table 1.3. With the help of such a table, one might be able to discover a general trend for the texture of all metals, finally leading to the establishment of a control theory for texture. 1.4.5 Bonding with substrate and crystallographic matching (epitaxy) (Watanabe, 1986, 1994) Figure 1.48 shows the surface morphology of electrolytic nickel films grown on polycrystalline copper substrates. The effect of two current densities (100 and 1000 A/m2) and three film thickness values (0.33, 5, and 12 mm) on the surface morphology is seen in this figure. The geometrical patterns of the polycrystalline copper substrates are copied on the nickel films and appear to persist, even to a thickness of 12 mm. The effect of the orientation and structure of substrates on the structure of electrolytic gold films is shown in Figure 1.49. Here, the , 100 . -, , 110 . - and , 111 . -oriented


69

Figure 1.47. A texture change with increasing film thickness and solution temperature in electroless gold films grown on amorphous alloy substrates. The reducing agent ¼ HCHO.

single crystals of copper and amorphous film were used as the substrates (Imai et al., 1983). Characteristic crystallographic surface morphologies are observed on the single-crystal substrates, whereas uniformly distributed fine surface irregularities are seen on the amorphous substrate. Electron diffraction patterns shown in Figure 1.49(B) are taken from

70

Nano-Plating

Figure 1.48. A surface morphology change with film thickness in electrolytic nickel films grown on polycrystalline copper substrates from a Watts bath at current densities of (a) 100 and (b) 1000 A/m2.

the film (gold) – substrate (copper) composites. We found the following crystallographic relationship. {100}Au//{100}Cu, , 100 . Au//, 100 . Cu {110}Au//{110}Cu, , 110 . Au//, 110 . Cu {111}Au//{111}Cu, , 111 . Au//, 111 . Cu


71

Figure 1.49. Electrolytic gold films grown on (a) single-crystal {100} copper, (b) {110} copper, (c) {111} copper, and (d) amorphous substrates. (A) The surface morphology of plated gold films, (B) electron diffraction patterns from the gold film/substrate composites, and (C) the indexed diffraction patterns.

Single-crystal gold films are seen to form by lattice-matching to the structure of the substrate. Based on these results, atom arrangements at the interface for both Au and Cu atoms are drawn in Figure 1.50. The top surface of the Au films reflects the symmetry of single-crystal orientation (see Figure 1.49(A)). Since the final exposed growth plane is {111}, the surface morphology can be constructed using the {111} facets. The gold film grown on the amorphous substrate is polycrystalline and does not exhibit any characteristic surface morphology. We have plated various metals on a number of metal substrates using electrolytic, electroless, displacement, and contact-plating techniques and studied the lattice-matching relationship (Watanabe, 1994). These results are listed in Table 1.2, where all the metals exhibit some structural relationship to substrate crystals. The driving force for such a lattice matching is to minimize the deposit – substrate interfacial energy. If the crystal system of the substrate is different from that of the deposit, the deposit is rotated and misfit dislocations (Gaigher et al., 1976) or misfit twins (Watanabe and Tanabe, 1976; Watanabe et al., 1984) are introduced at the interface. Figure 1.51 is a plan-view TEM micrograph of an iron film plated on electrontransparent {100} copper substrate. Two orthogonal sets of misfit dislocations are seen. Misfit dislocations appear to relieve the interfacial mismatch between the deposit and substrate.

72

Nano-Plating

Figure 1.50. Epitaxial relationship between a plated gold film and a single-crystal copper substrate (determined from Figure 1.39).

Figure 1.51. TEM observations of misfit dislocations formed at the electrolytic Fe/Au substrate interface.


73

A similar plan-view micrograph is shown in Figure 1.52 for a nickel film plated on a {100} iron substrate (Willson and Rogers, 1964). Contrary to many epitaxial systems, numerous micro-twins appeared. From an analysis of electron diffraction patterns, these

Figure 1.52. TEM observations of misfit twins formed at the electrolytic Ni/Fe substrate interface. (a) The TEM image, (b) its electron diffraction pattern from the Ni/Fe composite film, and (c) the indexed diffraction pattern.

74

Nano-Plating

micro-twins were formed to relieve the interfacial misfit and are thus called “misfit twins”. Similar misfit twins were found to form for plated chromium films grown on copper substrates (Watanabe et al., 1984). Metal films plated on metal substrates tend to match the crystallography of the substrate to minimize the interfacial energy. In films obtained by dry processes, such as vapor deposition, the film copies the structure of the underlying substrate up to some film thickness, although the structure of the film is different from that of the substrate. This phenomenon, called pseudomorphism, has not been observed in plated film systems. As described above, there are three mechanisms whereby thin crystalline films accommodate an interfacial mismatch between deposit and substrate. These are: (a) pseudomorphism; (b) misfit dislocations; and (c) misfit twins (see Figure 1.53). Lattice-matched crystalline films grow in thickness with time up to a critical thickness, whereupon the formation of grain boundaries causes the lattice matching to cease (Figure 1.54(a)). The critical thickness at which plated films lose the lattice coherency depends on the type of deposit – substrate metal combinations and plating conditions. Figure 1.55 shows an example of alloy metal substrates containing different phases. Gold films were deposited for 5 s (Kamasaki et al., 1974) on Fe-0.8 wt.% C steel substrates. Gold film is seen to deposit on the a-phase (ferrite phase) but does not deposit on the cementite phase (marked with Ce). After 30 s of plating, a thin gold deposit appears on the Ce phase. Ce is an intermetallic compound with high resistivity. Due to this high resistivity, the current density was low, making the gold growth rate low. The thickness difference of the gold film between the a-phase and Ce is clearly seen. 1.4.6 Residual stress (compressive, tensile stresses, cracks) The stress of plated films can be divided into two types: (1) a stress generated during plating and (2) a residual stress remaining in the films after plating. The former can become the latter, but all the stresses do not necessarily remain in the film as residual stress (Konishi, 1970; Yonetani, 1986; Kaneo, 1992). In an industrial sense, the residual stress is more important. High internal stresses in plated films often lead to film bending or cracking and thus they impose a serious issue to the plating industry. The production of a stress-free film has been often discussed by developing a stress-reducing addition agent. A number of studies have previously been devoted to the understanding of internal stresses in plated films. The internal stress changes with the current density (overpotential), the type of anions/additives, and the solution temperature. Even under identical plating conditions, plated films exhibit different internal stresses if the substrate is different (Yoshioka et al., 1999). Weil (1971) has made the following classifications for the origin of internal stresses in plated films: (a) a crystallographic mismatch between deposit and substrate; (b) a


75

Figure 1.53. Three types of interfacial phenomena occurring between a plated film and a single-crystal substrate during epitaxial growth. (a) Pseudomorphism, (b) the formation of misfit dislocations, and (c) the formation of misfit twins.

compressive force generated by an attractive force between three-dimensional crystallites during crystal growth; (c) a residual stress due to high overpotentials; (d) a stress generated by the incorporation of impurities such as hydroxides; and (e) a stress by hydrogen inclusions. From our previous microstructural studies on plated films, we consider that mechanism (a) is not acceptable. The reason being that once grain boundaries are formed between the

76

Nano-Plating

Figure 1.54. Typical growth forms in plated films grown on (a) polycrystalline and (b) amorphous substrates.

deposited film and the substrate, the mismatch stress does not affect the subsequent film. The origin of the second stress in (b) is equivalent to the stress generated by the fusion of two spherical particles to form one large sphere to reduce the surface energy. As described previously, deposited atoms are in a high-temperature state but will be cooled by the solution or the substrate while diffusing to form a crystal. When these heated atoms are finally cooled at the substrate, they are considered to develop a thermal mismatch due to the difference in the thermal expansion coefficient. The stress in (c) can be interpreted as originating from a difference in the thermal expansion coefficient between the deposit and the substrate. We do not agree with the development of internal stresses due to the incorporation of impurities in (d). The last mechanism (e) by hydrogen incorporation is indeed one of the most serious causes for the internal stress. The incorporation of hydrogen during plating processes causes the development of internal stresses in the plated films through film volume expansion or shrinkage. The behavior of hydrogen in the films depends on the susceptibility of metals to hydrogen. The state of hydrogen in metals is complex, for example, hydrogen could be interstitial or substitutional in metals or occluded as molecular hydrogen (H2). The diffusivity of hydrogen in metals is rapid and thus can diffuse out readily to form H2. The out-diffusion of hydrogen may cause a tensile

Microstructure Control Theory of Plated Film Figure 1.55. The structure of electrolytic gold films grown for (a) 15 and (b) 30 s on the heterogeneous surface of a steel substrate containing two different phases (Ce: cementite and a: ferrite).

77

78

Nano-Plating

stress in the film by shrinkage. It is important to consider the in-diffusion of hydrogen (hydrogen permeation) into the substrate region. The residual stress is an elastic strain in metal crystals and the expansion or contraction of the lattice spacing is a lattice strain. A force that generates the lattice strain is an elastic force. In comparison with crystals having ionic or covalent bonds, the lattice of materials with metallic bonds can be readily expanded or compressed. For this reason, metallic materials can be deformed elastically. The expansion or contraction of the lattice spacing left inside a crystal means that there is a residual stress in this crystal. In one material, both compressive (expressed as a minus sign (2 )) and tensile (expressed as a plus sign (þ )) stresses always coexist and the total stress is zero. In the former, the plated film becomes flattened after being stripped off the substrate. In the latter, the stripped plated film will curl up. In cases that there is an expansion in a plated film, a large compressive stress will be generated in the film as long as the rigidity of the substrate is high. If a tensile stress in a plated film becomes too high, a crack may be generated. These cracks are generally formed at the grain boundaries but not inside the grains. Figure 1.56 shows the warping of metal strips plated with Ni –P alloy films as a function of phosphorus concentration. This warping is due to a residual stress in plated films and a large warping indicates a high residual stress. The films having no or high phosphorus content appear stress-free but those with 5 , 6 wt.% P exhibit high residual stresses. The films with no phosphorus were crystalline deposits, obtained from a bath

Figure 1.56. A texture change with increasing film thickness and solution temperature in electroless gold films on amorphous alloy substrates. The reducing agent ¼ NaBH4.


79

of high current efficiency (low hydrogen evolution), and thus had a low residual stress. The þ films with high P content did not warp because they are amorphous deposits, which cannot be deformed elastically. Finally, the films containing 5 , 6 wt.% P were crystalline deposits obtained from a solution of low current efficiency (high hydrogen evolution). Hydrogen included in these films was responsible for the presence of a residual stress. 1.4.7 Anomalous morphology 1.4.7.1 Nodules. A nodule is a semi-spherical protrusion appearing on the surface of plated films. An example is shown in Figure 1.57(a). These nodules are formed on the surface of an electroless Ni – P film grown on a rolled sheet of polycrystalline copper (Tashiro et al., 1996). Figure 1.57(b) is the cross-sectional view of the film shown in Figure 1.57(a). The internal region of the nodule is crystallographically the same as the surrounding matrix region. Since there is no interfacial bond present between the nodule and the matrix, the film cracks readily along the interface if a bending stress is applied to the film (see Figure 1.57(b)). Some types of impurities, like hydrogen or phosphorus, are believed to segregate at the interface and thus weaken the interfacial bond strength. A nodule tends to grow in a shape of a cannonball from the base of the plated film and does not nucleate from the middle of the film. Once it is formed, it continues to grow and cannot be stopped unless the plating is terminated. The formation mechanism of the nodules is complex, and they appear in both electrolytic and electroless deposits. One of the causes of nodule formation is pre-existing fine protrusions on the surface of substrates, e.g. surface irregularities which are generally produced during etching prior to plating. In addition, two parallel nodules are seen to have nucleated at two bank regions of a V-shaped surface scratch. We have conducted an experiment that demonstrates the effect of etching time on the formation of nodules. Figure 1.58(A) shows how the density of nodules in electroless Ni– P films changes with increasing etching time. It is seen that as the etching time increases, more protrusions are produced, thereby increasing the number of nodules. Figure 1.58(B) shows how the surface morphology of electroless Ni– P films grown on an etched sheet of polycrystalline copper changes with increasing film thickness. It is clear that nodules form at the protrusions present on the substrate surface. In the case of electroless plating, nodule formation may be connected to the flow rate of solution; the electrolyte flows faster at the protrusions than at the surrounding area. The fast flow rate increases the deposition rate at the protrusions. In electroless plating, there are other factors that favor the formation of nodules. These factors include plating solutions and conditions. For example, we deposited electroless Ni –P films using hypophosphate as a reducing agent together with various chelates.

80 Nano-Plating Figure 1.57. Nodules formed on the surface of electroless Ni –P alloy films grown on polycrystalline copper substrates. SEM micrographs showing morphologies from (a) the top surface and (b) the top/fracture surfaces.


81

Figure 1.58. Nodules on the surface of electroless Ni–P alloy films grown on polycrystalline copper substrates. A change in the nodule structure (A) with copper substrate etching time and (B) with plating time in the films grown on deeply etched copper substrates.

82

Nano-Plating

Except for a small number of exceptions, more nodules were formed at some particular phosphorus concentration in electroless Ni – P films. This effect was attributed to the type of chelates used (Miyazaki et al., 1999). It should be noted that our discussion on nodule formation pertains to the case of electroless nickel and care should be taken if one wishes to extend our results to other plated films. 1.4.7.2 Pits (Nakayama, 1991). The surface of plated films often contain fine pores, called “pits”, which should be avoided in industrial plating. One of the main causes is the adsorption of contaminants on the substrate surface. At the sites, where contaminants are adsorbed, no plating takes place, thus leaving a hole. Hydrogen is also responsible for the formation of pits. Hydrogen included during electrodeposition diffuses through the film and agglomerates to form hydrogen gas bubbles at locations that eventually become preferential sites for pits to form. These pits act like a chimney for the bubbles to escape. Catalytic materials that are capable of converting atomic hydrogen into the molecular form might be adsorbed at the bottom of the pits. One example is shown in Figures 1.59 and 1.60. In Figure 1.60(a), the pit is seen to be mortar-shaped and its surface is very smooth, containing a hole at the bottom. Although no micrograph is shown here, foreign material was found at the bottom. We believe that the foreign material has a catalytic action allowing it to convert atomic hydrogen into hydrogen gas and thus facilitates the evolution of hydrogen gas bubbles. When the bubbles are small, they remain adsorbed at the bottom of the pits, where no plating takes place. The fact that the surface of the mortar bottom is smooth indicates that it was in contact with a hydrogen gas bubble. When the bubbles reach a critical size, they will be desorbed from the surface followed by further plating at such sites. Surface roughness will start developing inside the pits and at the same time the pit bottom will be filled (see Figure 1.60(b) and (c)). 1.4.7.3 Cracks. Surface cracking phenomenon in plated films has been described in conjunction with the internal stress. The cracking occurs if the tensile stress exceeds the tensile strength of the film. The origin of such a tensile stress was discussed previously but is not well understood. For this reason, no theories, which describe how to prevent cracking, are available. Surface cracks are generated during or after electrodeposition to reduce the internal stress. For example, a tensile stress can be developed in a plated film as supersaturated hydrogen diffuses out of the film. When such a tensile stress reaches a critical value, cracks may be generated in the film. The magnitude of internal stresses and the direction of cracking change with the shape of substrate materials, e.g. consider plating on the outer surface of a cylinder. Stresses along the axial direction are different from those along the radial direction. Consequently, cracks may develop along the cylinder axis or along the radial direction.


83

Figure 1.59. Pits formed on electrolytic nickel films.

1.4.7.4 Formation of layer structure. Various fracture surfaces of nickel films obtained from a non-agitated Watts bath by varying current densities are shown in Figure 1.61(A) (Watanabe and Kanayama, 1994). The thickness of each film is kept constant by keeping the total quantity of current the same. Cross-section SEM samples were prepared by first embedding in a resin, followed by mechanical polishing and finally by etching in hydrofluoric acid (see Figure 1.61(B)). Figure 1.61(B) (a), (d), and (f) shows the etched cross-section of electrolytic nickel films plated at the current densities of 100, 750, and 1500 A/m2, respectively. The structure obtained at 100 A/m2 is uniform along the thickness direction (see Figure 1.61(B) (a)),

84

Nano-Plating

Figure 1.60. Various forms of pits observed on electrolytic nickel films.


85

Figure 1.61. Various cross-sectional structures obtained by changing the current density in electrolytic nickel films grown from a non-agitated Watts bath. SEM images of (A) the as-fractured surfaces and (B) the cross-sections, which were polished mechanically with an Emery paper and buffing, followed by the final HF etching.

whereas at 750 A/m2 starts showing the layer structure near the surface region (see Figure 1.61(B) (d)). The film plated at 1500 A/m2 shows the layer structure throughout the thickness (cf. Figure 1.61(B) (f)). The fracture surface also indicates the presence of layers in these films. The fracture surface of nickel films deposited at 100 A/m2 shows a ductile fracture behavior. The films obtained at high current densities (750 and 1500 A/m2), which exhibited a layer structure, failed in a brittle mode. To understand the fracture mode, the cross-section of these samples was examined by TEM (see Figures 1.62 and 1.63). The layer structure actually consisted of large- and

86

Nano-Plating

Figure 1.62. A cross-section TEM image from the plated nickel film shown in Figure 1.61(A)(d) and its electron diffraction patterns taken from the four locations ((b)–(e)).

fine-grained regions. A large-grained layer exhibits ductile fracture, whereas a fine-grained layer fractures in a brittle manner. The difference in the fracture mode shows up as a layer structure (see Figure 1.61(A) (d) – (f)). The formation of a layer structure in nickel films obtained from Watts-type baths has been reported previously and its cause was attributed to impurities or to a change in the plating conditions. The reason for the periodic change in the grain size was believed to be periodic agitation by hydrogen gas evolving during plating.


87

Figure 1.63. An enlarged micrograph of fine- and large-grained layers seen in Figure 1.62.

This agitation changed the amount of hydrogen incorporation into the film periodically (Figures 1.62 and 1.63). Figure 1.64 is a TEM micrograph showing the cross-section of a Ni– Sn alloy film obtained from a pyrophosphate bath (Watanabe et al., 2000) without buffer or agitation.

88

Nano-Plating

Figure 1.64. Cross-section TEM micrographs of Ni–Sn alloy films plated for (a) 5 and (b) 10 min from a nonagitated solution containing (a) no buffer and (b) glycine as a buffer.

The structure in Figure 1.64(a) consists of three layers. A large-grained layer appears next to the substrate as an initial layer. The second layer is fine-grained and uniform, and the final layer is large-grained. The formation of this type of layer structure can be explained as follows. Firstly, the plating was done without a buffer or solution agitation. In addition, the pyrophosphate complex of Sn was reported to break up at pH ¼ 9:3 , 9:3: Although the pH was 8 at the beginning of the plating, it increased above 9.3 during the plating. Consequently, Sn was stripped from the pyrophosphate complex and became Sn ions. The plating proceeded with bare Sn ions and formed Sn-rich deposits instead of the desired composition of alloy films. The layer structure observed in the Ni– Sn alloy films, therefore, consists of the initial alloy layer followed by the Sn-rich alloy layer. The addition of a buffer, glycine, suppressed the formation of the layer structure as seen in Figure 1.64(b). 1.4.7.5 Initial layer (Watanabe, 1990). A layer at the beginning of electrocrystallization may take a different structural form from that of the subsequently grown thick layer. Such a layer can be often found in alloy deposits and is called the initial layer. The thickness of the initial layer depends on plating conditions and the type of alloy deposits, ranging from few nm to several hundred nm. The formation of the initial layer can be reduced by solution agitation during plating.


89

The first type is a layer grown epitaxially on a substrate crystal (see Figures 1.4 and 1.38). A polycrystalline layer will follow this layer. The second type often appears in alloy deposits as a layer with different alloy composition. In this case, the composition of alloy films containing more noble metal appears as the initial layer. With increasing thickness, alloy films whose composition is suitable for plating conditions will be formed. This example is shown in Figure 1.64(a). Here a Ni –Sn alloy film with high (51.5 at.%) Sn concentration initially forms, followed by a thick and uniform Ni-14.5 at.% Sn layer. Figure 1.62 is an example of a pure nickel electrodeposit. The nickel grains initially grow large but eventually become fine (Watanabe and Kanayama, 1994). In this case, the concentration of nickel ions on the substrate surface is initially high and thus the evolution of hydrogen is small. Since the nickel deposit contained less hydrogen, it had large grains. With further deposition, hydrogen starts evolving and an inclusion of evolved hydrogen caused the grains to be refined. Figure 1.65 is a cross-section TEM micrograph of an electrolytic Ni– P alloy film grown on a copper substrate and its diffraction patterns taken at the three locations, together with their corresponding phosphorus concentrations (Watanabe and Kanayama, 1994). A 60-nm thick fine-grained layer having the composition of 8.7 at.% P forms as an

Figure 1.65. A cross-section TEM micrograph for an electrolytic Ni–P alloy film obtained from a non-agitated solution. The electron diffraction patterns, (a)– (c), correspond to an amorphous layer, a fine-grained layer and the copper substrate, respectively.

90 Nano-Plating

Figure 1.66. Whiskers formed on electrolytic zinc films (Courtesy of Mr. Ryoichi Murai, Sambix Co., Ltd.).


91

initial layer. Beyond this initial layer, a uniform amorphous layer of Ni-25 at.% P formed. 1.4.7.6 Whisker (Takemura et al., 1986; Nagai et al., 1989; Lee and Lee, 1998; Ogasawara and Muroi, 1998, 2000). A whisker is a hair-shaped crystal of a few mm in diameter and a few mm in length, and grows on the surface of plated films a few hours, or a few days, after the plating is completed. The whisker is defect-free and is a nearly perfect single crystal. The growth occurs not from the whisker tip but from its base. A number of whisker growth mechanisms have been proposed, including internal stress, carbon codeposition, and hydrogen. Although the origin of whiskers can be explained using a dislocation model or a recrystallization model, its explanation is still not satisfactory. Whiskers generally appear in low-melting point metals (Sn, Cd, Zn, and Al) and their formation is considered to be due to compressive residual stresses present in the films. The appearance of whiskers in stress-free tin deposits is also reported. If copper or zinc is used as the substrate, tin films can react with the substrate at ambient temperatures and form their intermetallic compound. During the interdiffusion, an additional stress can be generated which affects the nucleation or growth of whiskers. Whiskers can even form on the oxidized film surface by breaking the oxide film. Figure 1.66 shows whiskers on the surface of a zinc deposit. It is clear from the micrographs that whiskers do not appear to originate from any specific sites, such as the interior regions of the grains or the grain boundaries. Most whiskers grow linearly but some start bending (Figure 1.66).

REFERENCES Aotani, K. (1950) J. Jpn Inst. Metals, B14, 55. Aotani, K. (1951) J. Jpn Inst. Metals, B15, 52. Aotani, K. (1953) J. Electrochem. Soc. Jpn, 21, 180. Arai, S. & Watanabe, T. (1996) J. Jpn Inst. Metals, 60, 1149. Bockris, J.O.M. (1964) The mechanism of electrodeposition of metals, Modern Aspects of Electrochemistry, vol. 3, Butterworths, London, p. 224. Bockris, J.O.M. & Razmneey, G.A. (1967) Fundamental Aspects of Electrocrystallization, Plenum Publ. Co., pp. 155– 185. Cavalotti, P.L., Colombo, D., Galbiati, E., Piotti, A. & Kruger, F. (1988) Plating Surf. Finish., 78. Eichkorn, G. & Fischer, H. (1967) Z. Phys. Chem. N. F., 53, 29. Enomoto, H., Fujiwara, Y., Isaki, M. & Ono, H. (1982) J. Metal Finish. Soc. Jpn, 33, 369. Fischer, H. (1960) Electrochim. Acta, 2, 50. Fujisawa, Y., Tsuji, M., Narishige, T. & Machida, K. (1993) J. Jpn Inst. Metals, 32, 247. Fukumuro, N., Chikazawa, M. & Watanabe, T. (1996) J. Surf. Finish. Soc. Jpn, 47, 461. Fukumuro, N., Imano, M., Chikazawa, M. & Watanabe, T. (1998) J. Magn. Soc. Jpn, 22, 1268.

92

Nano-Plating

Furuya, H. & Tanabe, Y. (1982) J. Jpn Inst. Metals, 46, 1042. Furuya, H., Misaki, Y. & Tanabe, Y. (1981) J. Metal Finish. Soc. Jpn, 32, 631. Gaigher, H.L. & van der Berg, N.G. (1976) Electrochim. Acta, 21, 45. Gerischer, H. (1960) Electrochim. Acta, 2, 1. Girin, O.B. (1995) Electron. Mater., 24, 947. Haruyama, S. (1963) J. Electrochem. Soc. Jpn, 31, 478. Hishino, S. (1976) J. Metal Finish. Soc. Jpn, 27, 145. Hishino, S. & Ro, B. (1973) J. Metal Finish. Soc. Jpn, 24, 567. Hishino, S. & Ro, B. (1974) J. Metal Finish. Soc. Jpn, 25, 504. Hosoiri, K. & Watanabe, T. (2001) Abstract at the 129th Meeting of Japan Inst. Metals, p. 416. Idemoto, M. (1999) J. Surf. Finish. Soc. Jpn, 50, 412. Imai, H., Watanabe, T. & Tanabe, Y. (1983) J. Metal Finish. Soc. Jpn, 34, 129. Inoue, K., Nakata, T. & Watanabe, T. (2001) J. Jpn Inst. Metals, 65, 229. Inoue, K., Nakata, T., Shindo, Y. & Watanabe, T. (2002) J. Jpn Inst. Metals, 66, 400. Isaki, M., Enomoto, H. & Omi, T. (1987) J. Metal Finish. Soc. Jpn, 38, 189. Jerkiewicz, G., Feliu, J.M. & Popov, B.N (2000) Hydrogen at surface and interfaces, Phys. Electrochem. Div., Electrochem. Soc., Proc. vol. 2000-16. Kamasaki, S., Tanabe, Y. & Matsumoto, Y. (1974) J. Metal Finish. Soc. Jpn, 25, 647. Kaneo, Y. (1992) J. Surf. Finish. Soc. Jpn, 43, 667. Kang, S., Yang, J.-S. & Lee, D.N. (1995) Plating Surf. Finish., 82(October), 67. Karayannis, H.S. & Patermarakis, G. (1995) Electrochim. Acta, 40, 1079. Koide, K., Kotuji, H., Teshigawara, T., Terakado, T., Nakata, T., Sone, T. & Watanabe, T (2000) Abstract at the 127th Meeting of the J. Japan Inst. Metals, p. 70. Konishi, S. (1970) J. Metal Finish. Soc. Jpn, 21, 470. Kossel, W. (1927) Nachr. Gea. Go¨ttingen, Math-Physik. K1, 135. Lee, B.Z. & Lee, D.N. (1998) Acta Met., 46, 3701. Liang, H., Chikazawa, M. & Watanabe, T (1995) Abstract at the 91st Meeting of Surface Finish. Soc. Japan, p. 73. Liang, H., Chikazawa, M. & Watanabe, T. (1999) J. Jpn Inst. Metals, 63, 474. Loebichir, O., Murakami, T., Raub, Ch. J. (1980) Extended Abstract, Fall Meeting of the Electrochem. Soc., Oct. 5 – 10, Hollywood, Florida, p. 1028. Mackinnon, D.J., Brannen, J.M. & Fenn, P.L. (1987) J. Appl. Electrochem., 17, 1129. Makino, Y., Inoue, K., Nakata, T., Sone, T., Watanabe, T (2001) Abstract at the 129th Meeting of Japan Inst. Metals, p. 416. Minami, T. & Mayanma, S.M. (1993) J. Electrochem. Soc., 140, 984. Miyazaki, S., Chikazawa, M. & Watanabe, T. (1999) Abstract at the 99th Meeting of the Surf. Finish Soc. Japan, p. 12. Mizushima, I., Chikazawa, M. & Watanabe, T. (1996) J. Electrochem. Soc., 143, 1978. Nagai, T., Natori, K. & Furusawa, T. (1989) J. Jpn Inst. Metals, 53, 303. Nakayama, I. (1991) J. Surf. Finish. Soc. Jpn, 42, 303. Narita, A. & Watanabe, T. (1991) J. Surf. Finish. Soc. Jpn, 42, 559. Ng, S.L. & Ling, H.C. (1990) J. Electrochem. Soc., 137, 458. Nichols, R.J., Bach, C.E. & Meyer, H. (1993) Ber. Bunsenges. Phys. Chem., 97, 1012. Nyung Lee, D. & Ye, G.C. (1981) Plating Surf. Finish, 68, 46. Ohno, I. (1988) J. Metal Finish. Soc. Jpn, 39, 149. Ohno, I. & Haruyama, S. (1991) J. Jpn Inst. Metals, 30, 735. Ogasawara, K. & Muroi, R. (1998) J. Surf. Finish. Soc. Jpn, 49, 502.


93

Ogasawara, K. & Muroi, R. (2000) J. Surf. Finish. Soc. Jpn, 51, 729. Okinaka, Y. & Straschil, H.K. (1986) J. Electrochem. Soc., 133, 2608. Onical, T. & Muresan, L. (1991) J. Appl. Electrochem., 21, 565. Onoda, M., Tsuchiya, T., Ogawa, K. & Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 388. Pangarov, N.A. (1962) Electrochim. Acta, 7, 139. Pangarov, N.A. (1965a) J. Electrochim., 9, 70. Pangarov, N.A. (1965b) J. Electroanal. Chem., 9, 70. Pangarov, N.A., Nenov, I. & Christova, I. (1963) Bull. Inst. Phys. Chem., (Sofia), 3, 133. Povetkin, V.V. & Devyatkova, O.V. (1996) Trans. Inst. Metal Finish., 74, 177. Raub, Ch.J. (1930) Plating Surf. Finish., 30(September). Raub, E. (1953) Metalloberfla¨che, 7, 17. Raub, E. & Sautter, F. (1957) Metalloberfla¨che, 11, 249. Reddy, A.K.N. (1963) J. Electroanal. Chem., 6, 141. Sard, R., Schwarz, C.D. & Weil, R. (1966) J. Electrochem. Soc., 113, 424. Sasaki, N., Nakata, T. & Watanabe, T (2001) Abstract at the 129th Meeting of the J. Japan Inst. Metals, p. 416. Schmidt, W.V., Alkire, R.C. & Gewirth, A.A. (1996) J. Electrochem. Soc., 143, 3122. Seiter, H., Fischer, H. & Albert, L. (1960) Electrochim. Acta, 2, 97. Shimizu, Y. & Tanabe, Y. (1976a) J. Metal Finish. Soc. Jpn, 27, 574. Shimizu, Y. & Tanabe, Y. (1976b) J. Metal Finish. Soc. Jpn, 27, 20. Shimizu, Y. & Tanabe, Y. (1978) J. Metal Finish. Soc. Jpn, 29, 21. Shimizu, Y., Tanabe, Y., Tomita, I. & Kakegawa, M. (1978) J. Metal Finish. Soc. Jpn, 29, 131. Stranski, I.N. (1928) Z. Phys. Chem., 136, 259. Tajiri, K. (2000) Proceedings of the 59th Fine Plating Meeting, Surf. Finish. Soc. Japan, p. 13. Tajiri, K., Kamihata, N. & Kajima, K. (1998) J. Surf. Finish. Soc. Jpn, 49, 127. Takemura, T., Kobayashi, M. & Okutani, M. (1986) Jpn. J. Appl. Phys., 25(Pt. 1), 1439. Tanabe, Y., Kakegawa, M. & Shimizu, Y. (1976) J. Mater. Sci. Soc. Jpn, 13, 255. Tashiro, K., Chiba, K., Fukuda, Y., Nakao, H. & Watanabe, T. (1996) J. Surf. Finish. Soc. Jpn, 47, 349. Terakado, T., Inoue, K., Nakata, T., Sone, T. & Watanabe, T. (2001) Abstract at the 129th Meeting of Japan Inst. Metals, p. 417. Teshigawara, T., Nakata, T., Inoue, K. & Watanabe, T. (2001) Scr. Mater., 44, 2285. Tsuru, Y., Kohno, M. & Hosokawa, K. (1993) J. Surf. Finish. Soc. Jpn, 44, 161. Vetter, K.J. (1961) Electrochemische Kinetik. Springer Verlag, Berlin, p. 698. Watanabe, T. (1986) J. Metal Finish. Soc. Jpn, 37, 440. Watanabe, T. (1989a) J. Surf. Finish. Soc. Jpn, 40, 375. Watanabe, T. (1989b) J. Surf. Finish. Soc. Jpn, 40, 122. Watanabe, T. (1990a) ISIJ Int., 76, 1597. Watanabe, T. (1990b) J. Surf. Finish. Soc. Jpn, 41, 652. Watanabe, T. (1992) Current Topics in Amorphous Materials, Physics & Technology. North-Holland Publ., p. 137. Watanabe, T. (1992 – 1993) Abstract at the 85th Meeting of the Surface Finish. Soc. Japan, p. 224. Watanabe, T. (1994a) J. Surf. Sci. Soc. Jpn, 15, 637. Watanabe, T. (1994b) Mater. Sci. Engng, A179/A180, 193. Watanabe, T. (1995) Seminar Text on Thermodynamics of Novel Materials, The Metallurgical Soc. Japan, p. 33.

94

Nano-Plating

Watanabe, T. (1998) The 3rd Thin Film Basic Seminar (Plating), The Surface Science Soc. Japan, p. 115. Watanabe, T. (2001) Mater. Jpn, 40, 871. Watanabe, T. & Arai, K. (1991) Amorphous Plating, 21, 6. Watanabe, T. & Kanayama, T. (1994a) J. Jpn Inst. Metals, 58, 132. Watanabe, T. & Kanayama, T. (1994b) J. Jpn Inst. Metals, 58, 138. Watanabe, T. & Masumoto, T. (1990) Plating methods of amorphous metals/alloys and their applications, Nikkan Kogyo Shimbun-sha, p. 145. Watanabe, T. & Minami, S. (2000) J. Jpn Inst. Metals, 64, 67. Watanabe, T. & Tanabe, Y. (1976) Trans. Jpn Inst. Metals, 17, 655. Watanabe, T., Hasegawa, N. & Tanabe, Y. (1984) Trans. Jpn Inst. Metals, 25, 531. Watanabe, T., Naoe, T., Mituo, A. & Katsumata, A. (1989) J. Surf. Finish. Soc. Jpn, 40, 458. Watanabe, T., Arai, K., Hirose, T. & Chikazawa, M. (1999) J. Jpn Inst. Metals, 63, 489. Watanabe, T., Arai, K. & Hirose, T. (2000) J. Jpn Inst. Metals, 64, 242. Watanabe, T., Sawanobori, H. & Osada, M. (2002) J. Jpn Inst. Metals, 66, 614. Watanabe, T., Tomita, M. & Miura, Y. J. Jpn Inst. Metals, to be still submitted. Watanabe, T., Kitagawa, Y., Mizushima, I., Takasaka, T. & Igarashi, N. J. Jpn Inst. Metals, to be submitted. Weil, R. (1971) Plating, 58, 137. Willson, K.S. & Rogers, J.A. (1964) Tech. Proc. Am. Electroplaters Soc., 5, 92. Winand, R. (1994) Electrochim. Acta, 38, 1091. Yan-Ling, Z., Ichino, R., Okido, M. & Oki, T. (1991) J. Surf. Finish. Soc. Jpn, 42, 1245. Ye, X., Bonte, H.D., Celis, J.P. & Roos, J.R. (1992) J. Electrochem. Soc., 139, 1592. Yonetani, S. (1986) J. Metal Finish. Soc. Jpn, 37, 449. Yoshihara, T., Nakata, T. & Watanabe, T (2002) Abstract at the 130th Meeting of Japan Inst. Metals, p. 434. Yoshioka, S., Chikazawa, M. & Watanabe, T. (1999) Abstract at the 99th Meeting of the Surf. Finish Soc. Japan, p. 11.

Chapter 2

Film Formation Mechanism in Electrodeposition 2.1.

Formation of electrolytic films 2.1.1 The initial nucleation and growth stages of pure metal films 2.1.2 Structure of electrolytic binary alloy films 2.1.2.1 Formation of phase separation type 2.1.2.2 Formation of solid-solution alloy 2.1.2.3 Formation of intermetallic compounds 2.1.2.4 Meta-stable phase 2.1.2.5 Amorphous phase 2.1.2.6 Mixed phase 2.2. Electroless films 2.2.1 Electroless plating method 2.2.1.1 Electroless plating method and microstructure 2.2.1.2 Formation mechanism of electroless films 2.2.1.3 Bonding between electroless films and substrate 2.2.2 Immersion coating 2.2.2.1 Film formation mechanism in immersion (displacement) coating 2.2.2.2 Bonding between immersion coating and substrate 2.2.3 Contact plating References

97 97 106 107 108 109 109 113 119 121 121 121 122 127 132 132 136 136 138


Chapter 2

Film Formation Mechanism in Electrodeposition 2.1. FORMATION OF ELECTROLYTIC FILMS

2.1.1

The initial nucleation and growth stages of pure metal films

The formation mechanism of pure metal and alloy films by electrodeposition is illustrated in Figure 2.1 (Gerischer, 1960). Metal ions, which are generally hydrated or complexed in a bulk electrolyte, first diffuse toward a cathode (substrate) through the potential gradient. After these ions pass through a diffusion layer and an electrical double layer (Helmholtz layer) present on top of the cathode surface, they will be stripped from the hydrate or complex ions, and become bare metal ions. The bare metal ions are then discharged by combining with electrons on the cathode, and become neutral atoms (adatoms). The adatoms now start migrating over the substrate surface until they are adsorbed at active sites, finally forming a strong chemical bond with the substrate. Kinks or steps present on the substrate surface are believed to be active sites for adatoms to be finally adsorbed. It can be seen that the successive adsorptions of adatoms at such sites result in the continuous spreading of the mono-atomic layer over the substrate surface. Seiter et al. (1960) proposed that the crystal growth takes place primarily by spiral growth around the emerging sites of screw dislocations on the substrate surface. We can visualize such a crystal growth process on a theoretical level. However, we are interested in knowing what happens in a practical situation. Figure 2.2(a) –(d) shows TEM micrographs showing the early stages of electrocrystallization for electrolytic gold grown on a single-crystal iron substrate (Kamasaki, 1974). The initial stages of thin-film growth have been categorized into three growth modes, as illustrated in Figure 2.3. These modes include: (a) Volmer-Weber (V-W); (b) Frank-van der Merwe (F-M); and (c) Stranski-Krastanov (S-K) types. Nucleation and growth processes seen in Figure 2.2(a) start with the formation of threedimensional islands via the V-W mechanism, but do not appear to show the F-M mode, which involves a monolayer growth by adsorption of adatoms at kinks and step edges. In addition, if the gold islands nucleate at the kinks or step edges of the substrate, their locations should indicate some systematic pattern specific to the substrate structure. No systematic pattern, however, can be recognized in Figure 2.2(a). The absence of a systematic arrangement in the locations of gold islands strongly indicates that the surface-diffusing adatoms are not adsorbed at the kinks or steps of the substrate surface. It is most likely that these adatoms meet randomly with other diffusing adatoms, then agglomerate in groups of two, three, etc. and finally form a three-dimensional 97

98

Nano-Plating

Figure 2.1. The initial stages of electrocystallization for (a) pure metals and (b) alloys.

Figure 2.2. The early stages of gold electrodeposition on single-crystal iron substrates. The film thickness increases in the alphabetical order.

Film Formation Mechanism in Electrodeposition

99

Figure 2.3. Three modes for describing the initial stages of thin-film growth. (a) Volmer-Weber (V-W), (b) Frank-van der Merwe (F-M), and (c) Stranski-Krastanov (S-K) types.

island. Gold islands thus formed grow by coalescence with other neighboring islands and eventually cover the whole substrate surface (see Figure 2.2(b) and (c)). A careful inspection of Figure 2.2(c) indicates that the shape of the gold islands has well-defined crystallographic edges, which were induced by the epitaxial growth of gold atoms over the structure of the copper substrate. This epitaxial growth is also seen in the inserted electron diffraction patterns (cf. Figure 2.2(b) and (c)). Although many islands, visible in the bright-field images of Figure 2.2(b) and (c), appear randomly distributed, they are actually arranged crystallographically because the corresponding diffraction patterns show a simple single-crystal pattern containing various diffraction spots. In addition to the single-crystal spots, faint diffraction rings are also present. It is generally accepted that when the size of such islands is very small, the islands do not develop a good lattice match to the substrate. The observed diffraction rings are an indication of these small islands being randomly oriented. The ring pattern was indeed stronger for the thinner films, which contained more random islands (see Figure 2.2(a)). A schematic view of this type of epitaxial growth is illustrated in Figure 2.4 (Watanabe, 1990a). Initial crystalline islands should take a semi-spherical shape, like water droplets, but in fact form some crystallographic shape by developing their lowenergy facets. The formation of faceted islands results from an energy balance between the island – substrate interfacial energy, the surface energy of the island, and the surface energy of the substrate. An electron diffraction pattern indicates that the gold film covering the iron substrate seen in Figure 2.2 is a single-crystal film. Since a gold metal has a different crystal system and lattice constant from an iron metal, it cannot lattice match to iron substrates perfectly. Therefore, we have a gold single-crystal film being grown on top of

100

Nano-Plating

Figure 2.4. A schematic diagram showing experimental results shown in Figure 2.2.

the single-crystal iron substrate through a low-angle sub-grain boundary that accommodates the lattice mismatch by introducing periodically distributed misfit dislocations. We have discussed the surface morphology of plated films prior to forming a continuous film. In island growth, the thickness, at which the deposit becomes a continuous film, is the height of the islands being linked by island coalescence. With further growth, these films develop their own characteristic surface morphologies. This is shown in electrolytic gold films grown on the {100} surface of copper (see Figure 2.5(a) – (d)) (Watanabe, 1990a). Steps, kinks, and pyramids are present on the surface of the gold film (cf. Figure 2.5(a)).


101

Figure 2.5. Growth processes of gold films on the {100} surface of a single-crystal copper. The film thickness increases in the alphabetical order.

After the copper substrate surface is completely covered by the gold deposit, subsequent gold deposition becomes similar to the homo-epitaxial growth of gold on gold. The gold adatoms are then adsorbed at the kinks or steps of the underlying gold surface and the film thickens by single- or multi-layer growth. At some critical thickness, pyramids form (Figure 2.5(b)) and eventually nucleate all over the surface (Figure 2.5(c)). This growth process is considered to be of the S-K type shown in Figure 2.3(c). With further growth, dendrites are finally formed (cf. Figure 2.5(d)). The reason why the pyramids start forming is that with an increasing film thickness, the gold deposit can lower the surface energy by exposing the lowest surface-energy plane (or the most stable plane), {111}. It can be easily seen that pyramids formed on the {100} substrate maintain the 4-fold symmetry by developing the 4 {111} facets. Similar surface-energy minimization processes can be achieved on the {110} or {111} planes as seen in Figure 2.6(b) and (c). The development of these surface morphologies can be understood by considering the outermost planes as the {111} planes. The metal-ion discharge process will further modify the geometry of the pyramidal structure on the {100} substrate. It has been shown in Chapter 1 that the metal discharge

102

Nano-Plating

Figure 2.6. Growth processes of gold films on the (1) {100}, (2) {110}, and (3) {111} surfaces of a single-crystal copper. The film thickness increases in the alphabetical order.

process will occur preferentially at pointed sites (protrusions), in this case, the tip of the pyramids (see Figure 2.7). In other words, the pyramidal morphology promotes rapid growth at the tips and slow growth at the valley regions. The deposited adatoms at the tip diffuse down the sloped {111} planes until they solidify (see a schematic in Figure 2.8). The dimension of the {111} plane represents the limit of the diffusion distance for these adatoms. For this reason, an increase in the height of these pyramids occurs by the stacking of pyramids in the shape of a pagoda but not by the enlargement of an individual pyramid.


Figure 2.7. Pyramids formed on an electrolytic gold film, which was grown on a single-crystal {100} copper substrate.

103

104

Nano-Plating

Figure 2.8. A model for the formation of pagoda-shaped crystals on the gold film shown in Figure 2.7.

Surface morphologies observed in Figure 2.5(c) and (d) cause a random light scattering, which results in a change of appearance in the order of golden luster ! no luster ! yellow ! orange ! brown: Obviously, these gold films do not have any industrial use. The pagoda-shaped crystals or dendrites are mechanically very weak and break easily. In the plating industry, various bath modifications have been attempted to obtain gold films free of these surface irregularities or dendrites. In the above section, the initial growth stages of electrolytic gold on copper substrates were shown to be of the V-W type (cf. Figure 2.2(a)). Our studies indicate that most pure metal and alloy electrodeposits take similar growth morphologies (Shimizu, 1976). In Figure 2.9(A) (a) – (c), we illustrate three modes of atom deposition and subsequent film formation on an ideal surface free of adsorbates.


105

Figure 2.9. Thin-film formation processes on metal surfaces. Film formation (A) [(a) Affinity between substrate atom and deposited atom affinity among deposited atoms (Frank–van der Merwe type growth). (b) Affinity between substrate atom and deposited atom ¼ affinity among deposited atoms. (c) Affinity between substrate atom and deposited atom affinity among deposited atoms (Vomer– Waber type growth)]. on ideal surfaces free of adsorbates and (B) [(a) Affinity between substrate atom and deposited atom . affinity among deposited atoms (without adsorbed atoms) (Frank–van der Merwe type growth). (b) Affinity between substrate atom and deposited atom ¼ affinity among deposited atoms (soft adsorbed some chemical species) (Vomer –Waber type growth). (c) Affinity between substrate atom and deposited atom , affinity among deposited atoms (hardly adsorbed some chemical species for ensample oxide film)] on non-ideal surfaces, on which some chemical species are adsorbed.

106

Nano-Plating

Figure 2.9(A) (a) represents a monolayer growth of the F-M type shown in Figure 2.3(b). The mode in Figure 2.9(A) (c) is the V-W type (cf. Figure 2.3(a)), which undergoes a stage of island formation, whereas the mode in Figure 2.9(A) (b) is intermediate between the F-M and V-W types. We will describe these three modes in more detail below. For the case of Figure 2.9(A) (a), depositing atoms form strong chemical bonds (affinity) to substrate atoms and thus can spread easily over the substrate surface. This system represents good wettability. Figure 2.9(A) (c), on the other hand, is an instance where the attractive force between depositing atoms is stronger than that between the depositing atom and the substrate atom. Here, depositing atoms tend to coagulate together on the substrate surface and form an island crystal. This system is the case of poor wettability. Figure 2.9(A) (b) is the intermediate between the two cases shown in Figure 2.9(A) (a) and (A) (c). The initial growth morphology on adsorbate –free substrate surfaces can be determined by the relative magnitude of attractive forces existing in depositing atoms and substrate atoms. In the V-M type growth system, no combination of attractive forces can produce the F-M type monolayer growth. It is therefore impossible to fabricate a superlattice of various metal layers by monolayer-by-monolayer growth. Some information on the metal-to-metal atom affinity can be obtained from the equilibrium binary alloy phase diagram. Alloy systems forming simple eutectic or peritectic phases have almost no mutual atom affinities. Alloy systems forming intermetallic compounds, on the other hand, have a strong affinity, i.e. the higher the melting point of the compound, the stronger the chemical bond and the affinity. We now consider a real surface, on which some chemical species, such as H or OH, is adsorbed. On such a surface, even the F-M type monolayer growth (cf. Figure 2.9(B) (a)) cannot maintain its mode and tends to change to the V-W type (cf. Figure 2.9 (B) (b) or (B) (c)). If the bonding force of the adsorbate to the substrate is weak, depositing atoms will bond strongly to the substrate surface while expelling the adsorbate atoms or molecules. If adsorbates are strongly bonded to the substrate, depositing atoms cannot push the adsorbates away and consequently form an island on top of the adsorbate layer, resulting in the formation of the V-W type layer. This situation corresponds to film growth on top of passive films like oxides. Here, the bond strength between the deposit and the oxide film may be weak, although the oxide-to-substrate bond is expected to be strong.

2.1.2 Structure of electrolytic binary alloy films Five types of microstructure form when a mixture of two metals undergoes the standard metallurgical processes involving melting, cooling, and solidification (see Figure 1.3(b) – (f)). The same microstructures are present in plated binary alloy films. In Figure 2.1(a), we have shown schematically how pure metals are formed on a cathode surface. The case of binary alloy deposition is illustrated in Figure 2.1(b).


107

The alloy films also produce various microstructures, which depend on the type of metals and their composition. 2.1.2.1 Formation of phase separation type. Suppose that two kinds of atoms, A and B, deposit simultaneously and diffuse over the cathode surface. If an affinity between like atoms (A –A or B – B) is stronger than that between two different atoms (A –B), like atoms tend to agglomerate together and form their own crystalline phases. The microstructure thus formed is expected to be that shown in Figure 2.10 (see also Figure 1.3(b)). This is a formation mechanism for two A-rich and B-rich alloy phases during electrodeposition.

Figure 2.10. The microstructure of alloy electrodeposits with phase separation, eutectic, and peritectic types.

108

Nano-Plating

Figure 2.11. (a) The equilibrium phase diagram of Ag– Cu alloys (eutectic type) and (b) the lattice parameters of plated Ag– Cu alloy films versus the composition.

This type of microstructure often appears in alloy systems having eutectic/peritectic phases or phase separation (spinodal decomposition). Figure 1.27 in the previous chapter is a TEM micrograph for an electrodeposited Ag –Co alloy film. The Ag – Co system forms an immiscible alloy over the entire composition range. Therefore, Ag and Co crystallize separately. In Figure 2.11(a) and (b), we show the equilibrium binary alloy phase diagram of Ag/Cu and the lattice parameters of plated Ag –Cu alloy films as a function of alloy composition (Tanabe et al., 1986). Since the Ag-rich and Cu-rich phases crystallize separately, their lattice parameters change independently over the entire composition range. Figure 2.10 shows that the Ag-rich and Cu-rich phases did indeed crystallize independently, thus exhibiting separate lattice parameter values. It is also interesting to note that the lattice parameter of each phase are constant and do not change significantly over the whole composition range, implying that the composition of each phase must be close to pure silver or pure copper. One of the well-known eutectic alloys is the Pb –Sn system used as solders, which phase-separate into Pb- and Sn-rich phases and form the microstructure shown in Figures 1.3(b) and 2.10. The smallest grain size was obtained from the 50 at.% Pb – Sn (50:50 Pb –Sn) alloy but not from the eutectic alloy (26.1 at.% Pb – Sn). A deviation towards the Pb-rich side from the 50:50 composition yielded a small grain size, whereas the deviation toward the Sn-rich side gave a larger grain size. This grain size effect can be understood from the fact that the melting point of Pb (327 8C) is higher than that of Sn (232 8C). We have shown in Figure 1.27 of Chapter 1 that in alloy systems, the magnitude of melting point temperature of each metal affects the grain size. 2.1.2.2 Formation of solid-solution alloy. Suppose that two kinds of metal atoms, denoted as A and B, deposit on a substrate and the affinity of the A – A bond is the same as that of the B– B and A – B bonds. Then, the probability of forming the A –A bond is the


109

same as that of the B – B and A – B bonds. Furthermore, if the size of both metals and their atom coordination number are approximately the same, it is possible to obtain a uniform solid-solution alloy of substitutional type, whose atom positions are occupied randomly by two kinds of atoms. A typical example for this solid-solution type is the Au –Pd alloy system (Shimizu and Tanabe, 1976a). The equilibrium phase diagram of Au – Pd is of a solid-solution type over the entire compositional range (see Figure 2.12(a)). The lattice parameters of Au –Pd alloy films obtained from cyanide and chloride baths are plotted in Figure 2.12(b) and (c), respectively. In both cases, the lattice parameter changes linearly from the lattice constant of Au (0.4079 nm) to that of Pd (0.3890 nm) and follows Vegards law. These results confirm that the solid-solution structure of the plated films is in thermal equilibrium regardless of the type of the bath used (cyanide or chloride type) or plating conditions. This result further supports our previous statement that “ The structure of plated films is determined solely by the film composition regardless of the type of plating baths used.” This statement is applicable not only to solid-solution alloys in the whole composition range, but also to those formed in a limited composition range. Interstitial-type solid-solution alloys are generally formed with elements having small atomic radii, such as hydrogen, carbon, or nitrogen, which can occupy sites between the host atoms. Some non-metallic atoms, like phosphorus and boron, although having a small atomic radius, cannot be introduced interstitially. If phosphorus or boron atoms are introduced as an alloying element for plated films, they generally segregate at the grain boundaries and help refine the grains. Electrolytic or electroless Ni – P films are a typical example for this class of alloys. 2.1.2.3 Formation of intermetallic compounds. In alloy deposition, two kinds of atoms deposit simultaneously on the cathode and diffuse over the surface (see Figure 2.1(b)). The formation of intermetallic compounds becomes possible if these two atoms meet favorable conditions in terms of the atom size, the coordination number, and thermodynamic stability. Figure 2.13 shows the equilibrium phase diagram of Au –Sn alloys (cf. Figure 2.13(a)) and the phase diagrams of plated Au – Sn films obtained at 100 and 5000 A/m2 (cf. Figure 2.13(b) and (c)) (Tanabe et al., 1983). All the phases listed in the equilibrium phase diagram appear in plated films. An increase in the current density (or overpotential) from 100 to 5000 A/m2 did not change the phase formation significantly. 2.1.2.4 Meta-stable phase. The formation mechanism of meta-stable phases is essentially the same as that of intermetallic compounds. These phases are

110

Nano-Plating

Figure 2.12. (a) The equilibrium phase diagram of Au– Pd binary alloys, (b) the lattice parameters of Au–Pd alloy films plated from a cyanide bath, and (c) the lattice parameters of Au–Pd alloy films grown from a chloride bath.


111

Figure 2.13. (a) The equilibrium phase diagram of Au–Sn binary alloys, (b) the lattice parameters of Au–Sn alloy films plated at the current density of 100 A/m2, and (c) the lattice parameters of Au– Sn alloy films plated at 5000 A/m2.

thermodynamically unstable and are in a state of meta-stability. Meta-stable phases often appear in plated alloy films containing high melting point metal(s) and intermetallic compound(s). Figure 2.14(a) is the equilibrium phase diagram of Au – Ni alloys. The Au – Ni system forms solid-solution alloys over the entire composition range at high temperatures but generates phase-separated alloys at low temperatures. In Figure 2.14(b), we show the lattice parameters of several phases observed in plated Au –Ni alloy films (Shimizu and Tanabe, 1976b). In the composition range of 50 –95 at.% Ni, a new phase (denoted as Au – Ni (M)) not recorded in the equilibrium phase diagram appeared but phase-separated into stable a1 and a2 phases upon heat treatment. This new phase was apparently thermodynamically meta-stable. In addition, the new phase had a fcc structure, whose lattice parameter changed with the composition.

112

Nano-Plating

Figure 2.14. (a) The equilibrium phase diagram of Au–Ni binary alloys (a solid– solution alloy system with a miscibility gap) and (b) the lattice parameters of various phases appearing in plated Au–Ni alloy films.


113

Another example of the formation of meta-stable alloys in plated films is the appearance of supersaturated solid-solution alloys, which contain a solute at a concentration exceeding the solubility limit specified in the equilibrium phase diagram. The degree of supersaturation in a solid-solution alloy can be extended up to the maximum solubility achieved at high temperatures. As will be described later, plated Ni – Sn alloy films (Watanabe et al., 1999a,b) (cf. Figure 4.2) form not only supersaturated solid-solution alloys, but also two kinds of meta-stable phases. As discussed in Chapter 1, the appearance of meta-stable phases is common in plated films. The structural model shown in Figure 1.3(d) is only a schematic illustration and thus does not necessarily represent the real meta-stable crystal structure. In general, the structure of meta-stable phases is simpler than that of thermally stable intermetallic compounds. For example, in the structure of a meta-stable crystal shown in Figure 1.3(d), foreign metal atoms are inserted periodically into every other atom row. Such an atom arrangement can be realized only if the element of ionic bonds is present for building a meta-stable crystal structure. However, as seen in Figure 2.14(b), the lattice parameter of the meta-stable phase, i.e. Au – Ni (M) changes with the alloy composition. The reason for the change in lattice parameter is that the meta-stable phase is built not only by ionic bonds, but also by metallic bonds, which can adjust the lattice parameter by introducing the compositional change. 2.1.2.5 Amorphous phase. Amorphous films can be obtained by electrodeposition (Watanabe, 1989, 1993; Diokic, 1999). An amorphous Ni – P film is one well-known alloy system that can be grown by both electrolytic and electroless methods. These amorphous films contain no crystalline facets or grain boundaries and are uniform structurally and compositionally. For this reason, Ni –P alloy films have high brightness/corrosion resistance and are used for various decorative as well as protective coating applications. First, we will discuss the formation mechanism of amorphous films by electrodeposition. We consider two kinds of atoms, A and B, depositing on a cathode surface (see Figure 2.1(b)). Assume that the affinity between different atoms (A– B) is stronger than that of like atoms (A – A and B –B), and both the atomic radius and coordination number of atom A are different from those of atom B. Under these conditions, it is difficult to grow alloy films containing atoms arranged in an orderly fashion, thus resulting in the formation of a random structure shown in Figure 2.15 (see also Figure 1.3(e)). This is the basic formation mechanism of amorphous films by electrodeposition. This mechanism will be further elaborated using the binary phase diagram given in Figure 2.16. Figure 2.16(a) is a binary phase diagram for metals A and B. In this diagram, one intermetallic compound forms in the middle composition range, surrounded by two eutectic systems. The formation of the intermetallic compound in this diagram strongly suggests that A and B atoms have a strong chemical affinity to each other.

114

Nano-Plating

Figure 2.15. A structural model for plated amorphous films.

Now we consider an alloy whose composition is close to that of the intermetallic compound. Suppose that the alloy solidified without receiving sufficient activation energy for the formation of an intermetallic compound. Then, it is possible to ignore the existence of the intermetallic compound in the phase diagram and thus to hypothesize one large eutectic system by extending the liquidus lines as indicated with dotted lines (see Figure 2.16(a)). Near the eutectic point in the hypothetical eutectic system, there will be a liquid state (random structure) at a temperature, T1. In the case of alloy plating at this temperature, T1, a mixture of A and B atoms having the composition of the intermetallic composition will deposit and form a solid film on the substrate surface. However, if insufficient activation energy for the crystallization is

Figure 2.16. (a) A binary alloy phase diagram having one intermetallic compound in the middle composition range, surrounded by two eutectic systems. If a sufficient activation energy for the formation of the intermetallic compound is not given at the temperature T1, it is possible to produce a meta-stable amorphous phase. (b) The free energy diagram at the temperature T1, showing that the free energy of the amorphous phase is larger than that of the intermetallic compound.


115

provided, the mixture will solidify very rapidly at this temperature, and finally form an amorphous phase. In the previous chapter, therefore, we have emphasized that plated films are similar to solids obtained by rapid cooling from high temperatures (see Figure 1.4). Figure 2.16(b) shows the free energy diagram for the intermetallic compound and the amorphous phase at the temperature, T1. Although the free energy of the intermetallic compound is lower than that of the amorphous phase, the amorphous phase can still exist in a meta-stable state. In the above, we presented qualitatively a thermodynamic explanation for the formation mechanism of an amorphous film by electrodeposition. Suppose that the composition of plated alloy films in Figure 2.16(a) is changed from that of the intermetallic compound (50:50) to the A-rich side. The structure of the alloy films changes from Figure 2.17(a) to Figure 2.17(b) and (c). Here, the atoms A and

Figure 2.17. A schematic diagram showing how the structure changes with the composition of plated alloy systems, which form an amorphous phase.

116

Nano-Plating

B are shown as light and dark, respectively. It is seen that due to the strong affinity between A and B atoms in this alloy system, there is a strong tendency for A and B atoms to bond together preferentially. When the concentration ratio of the light atom (A atom) to the dark atom (B atom) is 80:20, then the area occupied by aggregated A atoms increases as seen in Figure 2.17(c). A further increase in the concentration ratio to 95:5 continues to expand the region of the A-atom aggregates. Such an aggregation assists these alloy films in becoming crystalline. This demonstrates that even under the favorable condition of amorphous film formation, the alloy film may become crystalline if the composition changes. The composition effect on the structure of electrolytic Ni– P films can be seen in X-ray diffraction patterns (see Figure 2.18) (Itoh et al., 2001). Here, the films are amorphous if the P concentration (B atom) is high, and become crystalline with decreasing P concentration. Additional examples are shown in Figure 2.19, where X-ray diffraction patterns from: (a) electrolytic Ni – B (Onoda et al., 1995, 1999); (b) electrolytic Ni – S (Narita and Watanabe, 1991); and (c) electrolytic Fe– Mo (Watanabe et al., 1989; Zhu and Watanabe, 1992)

Figure 2.18. X-ray diffraction patterns for electrolytic Ni –P alloy films grown at various current densities, which in turn produced the alloy films with different phosphorous concentrations.


117

Figure 2.19. X-ray diffraction patterns for electrolytic (a) Ni –B, (b) Ni –S, and (c) Fe– Mo alloy films having various B, S, and Mo concentrations, which were obtained by varying the current density.

118

Nano-Plating

are plotted as a function of current density and alloy composition. With increasing concentration of B, S, and Mo, the structure of these films changes from the crystalline to amorphous phase. The crystal size becomes larger at high current densities (overpotentials) and becomes smaller at low current densities. With further decreases in current densities, the grains become increasingly smaller and ultimately transform to an amorphous phase. This current density effect directly contradicts the overpotential theory (Vetter, 1961; Ohno and Haruyama, 1991; Winand, 1994), which states that “The grain size decreases with increasing current densities (overpotentials).” Applying heat treatment to plated alloy films often helps determine which alloy composition produces crystalline or amorphous phases (Suda et al., 1981a,b; Wang et al., 2000; Itoh et al., 2001). For example, if plated films are completely amorphous, they should crystallize at some well-defined crystallization temperature and nucleate crystals of the intermetallic compound. Such a crystallization behavior can be tested by heat treatment. Another useful example is to heat-treat plated films consisting of microscopically fine crystals, which could be mistakenly assumed amorphous from a broad X-ray diffraction pattern. In this case, heat treatment will reveal whether or not grain growth originated from pre-existing fine grains or from new grains nucleated out of an amorphous phase. More details on this subject will be described in Chapter 4 (Itoh et al., 2001) (cf. Figure 4.15). In the example of electrolytic Ni –P alloy films, we have observed that the P concentration is high at high current densities and low at low current densities (cf. Figure 2.18). This phenomenon can be understood using the concept of adsorption. We note that nickel atoms are ionized and thus incorporated electrochemically. Non-metallic elements such as P are not ionized, and thus their incorporation can only be accomplished by being adsorbed on the cathode surface (Onoda et al., 1995, 1999). As listed in Table 2.1, a large number of alloy films can be obtained by electrolytic and electroless methods (Watanabe, 1989, 1993). It is interesting to note that in most alloy systems, amorphous phases are produced around the composition of their intermetallic compounds. These amorphous phases are in a meta-stable state and thus transform to their stable intermetallic crystalline phases upon heat treatment. In summary, we can offer the following conclusions. In pure metal plating, amorphous phases stable at room temperature cannot be prepared. In alloy plating, the formation of amorphous phases requires that: (1) the alloy system must form intermetallic compound(s) and (2) the composition of the alloy deposits have to be near that of the intermetallics. No amorphous phase can be obtained from an alloy system consisting of simple solidsolution, eutectic, or peritectic types.

119

Film Formation Mechanism in Electrodeposition Table 2.1. Various amorphous alloys obtained by electrolytic and electroless methods. Electroplating Ni–H Pd– H Cr–H Cr–H Cr–W–H Cr–Mo –H Cr–Fe –H

Electroless plating Ni–P Co –P Ni–Co –P Ni–Fe –P Ni–Mo –P Ni–W–P Ni–Cu –P Pd– P–(H) Pd– Ni–P Ni–Re –P

Ni–P Fe–P Co –Ni–P Co –Zn–P Ni–S Co –S Cr–C Pd– (As) Ni–B Co –W–B Ni–Cr –P Ni–Fe –P

Ni –W Co –W Fe –W Ni –Mo Co –Mo Fe –Mo Co –Re Co –Ti Fe –Cr Fe –Cr–P Fe –Mu–Wd

Bi–S Bi–Se Cd –Te Cd –Se Cd–Se– S Si Si–C –F

Ni–B Co –B Ni–Co –B Co –W–B Ni–Mo –B Ni–W–B

2.1.2.6 Mixed phase. We have discussed a formation mechanism of various phases and microstructures in plated binary alloy films. Plated films may consist of a single phase, but often contain more than one phase. Figure 2.20 is a cross-section TEM micrograph for an electrolytic Ni – Sn alloy film (Watanabe et al., 1999a). The overall microstructure taken at a low magnification is seen in Figure 2.20(a) and a part of the micrograph is enlarged in (b). Electron diffraction patterns from three kinds of crystalline phases are shown in Figure 2.20(c), (c0 ), and (c00 ). From an electron diffraction analysis of these phases, we identified the phases, c, c0 , and c00 , to be Sn, a meta-stable Ni – Sn crystal, and Ni3Sn4 (dphase), respectively. The simultaneous presence of three phases in plated binary films is thermodynamically rare but not impossible. It is important to recognize that each phase has its individual electrical conductivity. For example, the presence of two phases having two markedly different conductivity values on the cathode surface is expected to affect the discharge process of depositing atoms. We show this example in Figure 2.21(a), where the degree of metal-ion discharge on the electrode surface (the film surface) is high at high-conductivity grains and is low at low-conductivity grains. As a result, the film develops a thickness variation, which could be reduced by the surface diffusion of depositing atoms.

120

Nano-Plating

Figure 2.20. (a) The cross-sectional structure of a plated Ni–Sn alloy film and (b) the enlarged micrograph showing three phases marked with (c), (c0 ), and (c00 ), whose electron diffraction patterns are displayed next to the micrograph.

Figure 2.21. A schematic diagram illustrating how (a) the current density distribution and (b) metal deposition/surface diffusion characteristics are affected by the heterogeneous nature of a substrate surface containing mixed grains with different electrical conductivities.


121

2.2. ELECTROLESS FILMS

An electroless plating method is a metal deposition technique without the application of an external power source. This method can be broadly divided into three types: (1) electroless; (2) displacement (immersion); and (3) contact platings. We will describe each plating method and its formation mechanism. 2.2.1 Electroless plating method 2.2.1.1 Electroless plating method and microstructure. The basic electroless plating solution consists of a metal salt and a reducing agent for converting metal ions to a metal. A solution containing only these two ingredients is highly unstable and a chemical reaction proceeds spontaneously, accompanied by the precipitation of the metal power. In practice, a complexing agent, a pH buffer/stabilizer, and other chemical reagents must be added to suppress such chemical reactions and to stabilize the solution. An immersion of substrate materials in this solution will initiate the metal reduction process on the substrate surface, which acts as a catalyst. Reduced metal atoms will be adsorbed (chemically bonded) on the substrate surface, followed by the formation of a metal film. After the metal film covers the whole substrate surface, it will act as an autocatalyst, which helps increase the thickness. Typical metals obtainable by an electroless plating method and their reducing agents are listed in Table 2.2. If phosphorus- or boron-based reducing agents are used, the top three metals (Ni, Co, and Pd) form alloys with phosphorus or boron (Ni –P, Ni –B, Co –P, Co – B, Pd –P, etc.). The quantity of these incorporated alloying elements depends on the concentration of reducing agent, pH, and solution temperature. Under conditions of fast reaction rates, the proportion of these elements in the deposits decreases. The crystal size changes with the amount of the alloying element. With increasing alloying element, the film becomes fine-grained and at some concentration, turns amorphous. Table 2.2. Reducible metals by electroless plating methods and their reducing agents. Metal

Reducing agent

Ni NaH2PO2, DMAB, NaBH4, KBH4, NH2NH2, HCHO Co NaH2PO2, DMAB, NaBH4, KBH4, NH2NH2, HCHO Pd NaH2PO2, NaHPO, NH2NH2 -------------------------------------------------------------------------------Cu HCHO, DMAB, KBH4 Ag DMAB, KBH4 Au DMAB, KBH4, NH2NH2, NaH2PO2, CH4N2S, C6H8O6 Pt NH2NH2, NaBH4 Pb SnCl2 Rh NH2NH2 Ru NaBH4

122

Nano-Plating

Figure 2.22. X-ray diffraction patterns for Ni –B alloy films with various boron concentrations, obtained by an electroless method using dimethylamine borane (DMAB) as a reducing agent.

Figure 2.22 is an X-ray diffraction pattern for electroless Ni– B films obtained from a bath containing dimethylamine borane (DMAB) as a reducing agent (Suda et al., 1981). It is seen from the continuous broadening of the X-ray peak with increasing B content that the crystal size became increasingly finer and finally transformed to an amorphous phase. In this experiment, the B content in the deposit was varied by altering the quantity of DMAB. Contrary to the case of DMAB, the use of formaldehyde or formalin as a reducing agent did not result in the incorporation of non-metallic elements in these metals (Ni, Co, and Pd). Consequently, the film was large-grained and the surface irregularity increased. 2.2.1.2 Formation mechanism of electroless films (Watanabe and Tanabe, 1974a; Watanabe, 1990b). A film formation mechanism by an electroless plating method will be described in this section. Contrary to the electrolytic method, an electroless method supplies no external electric field for attracting metal or complex ions toward the substrate


123

surface. In the electroless plating process, metal ions migrate toward the substrate surface via diffusion through the metal-ion gradient or drift with the aid of solution agitation. Here, we assume that the substrate surface acts as a catalyst for a reducing agent. Consequently, the reducing agent will be oxidized on the catalytic substrate surface and generate electrons. These electrons will combine with incoming metal ions and form neutral metal atoms. The neutral metal atoms become adatoms after undergoing a physical or chemical adsorption process on the substrate surface, followed by the formation of an electroless film. This process is shown in Figure 2.23 (Watanabe and Tanabe, 1974a; Watanabe, 1990b). As illustrated in Figure 2.23(a), an electroless plating process can be regarded as a surface reaction phenomenon on the substrate. For initiating metal deposition, the substrate surface has to possess catalytic activity. In other words, immersion of substrate materials having a catalytic activity, into an electroless solution results in spontaneous metal deposition. In some cases, however, such metal deposition may not occur readily, even on catalytic substrates. The lack of metal deposition can be attributed to the adsorption of chemical species, such as H, OH, or O, which generally make the substrate surface catalytically inactive. In this instance, the metal deposition can be initiated by passing

Figure 2.23. A schematic diagram showing a mechanism of electroless film formation. The film formation processes are illustrated: (a) at the edge of the metal island, (b) in the cross-sectional view, and (c) from the top view.

124

Nano-Plating

a current, e.g. by touching the substrate with a less noble metal like aluminum (contact plating). This type of practice is called an initiation treatment. After the metal deposits on the substrate by the initiation treatment, it will serve as an autocatalyst for subsequent metal deposition. It is important to remember that occasionally, metal deposition may not occur due to a reduction in catalytic activity caused by adsorbed chemical species. Here, this problem can be remedied with the initiation treatment. A metal layer first forms on a catalytic substrate surface in the form of a flat island. We call this metal layer the 1st layer. As opposed to the top surface region, the edge of the metal island is highly activated due to heat evolution and violent solution movement caused by the reaction (see Figure 2.24(a)). Therefore, the edge region will grow preferentially and expand laterally in a two-dimensional manner over the substrate surface. When the island height reaches a certain critical thickness (t1), a layer of low metal-ion concentration is generated at the solution –island interface in much the same way as the formation of the metal-ion deficient layer (MIDL) in the electrolytic case. Consequently, the film cannot grow beyond this thickness (t1) until the layer of low metal-ion concentration is destroyed by solution agitation and replaced with a fresh solution of high metal-ion concentration (HMIC) from the bulk electrolyte. When the HMIC solution comes into contact with the 1st layer, the formation of the 2nd layer becomes possible. In the same manner as the 1st layer, the 2nd layer will grow up to the thickness of t1 on top of the 1st layer. The same growth process will be repeated for the 3rd layer and so on (see Figure 2.23(b) and (c)). It is clear that a layer having the thickness of t1 is the basic building unit for the layered microstructure of electroless films. The thickness, t1, depends on the metal-ion concentration and temperature of the solution. We estimate t1 to be less than 10 nm, which will not be discernible microscopically. An increase in the metal-ion concentration and solution temperature is expected to increase the thickness, t1, and vice versa. Accordingly, the growth rate of electroless films is controlled by the thickness, t1, agitation speed, and solution temperature. In principle, the film cannot grow without solution agitation. In practice, however, there will be a certain amount of solution convection and thus the film can grow without agitation. Microscopic examination of the cross-section of electroless films often reveals the presence of a layer structure. This layer structure is not the same as the one described here. These layers represent a trace of compositional modulation induced by a change in solution pH and temperature. Figure 2.24 is a plan-view TEM micrograph showing the formation of the 1st, 2nd, and 3rd layers in electroless Ni – B films. In Figure 2.24(a), the 1st layer (denoted with (1)) is seen in the form of islands. The region marked with (0) is the substrate, on which no metal islands have yet grown. In Figure 2.24(b), the overlapped 1st (marked with (1)), 2nd ((2)), and 3rd ((3)) layers are visible. The example in Figure 2.24(c) is the case where the 1st


125

Figure 2.24. TEM micrographs showing the initial film formation stages of electroless Ni –B alloy films. Symbols denoted by (0), (1), (2), and (3) in the micrographs indicate the number of overlapped layers. Growth of (a) the 1st layer, (b) the triple (1st, 2nd, and 3rd) layer, and (c) the 2nd layer on top of the 1st layer.

126

Nano-Plating

Figure 2.25. A cross-section TEM micrograph showing an electroless Ni –P alloy film grown on an etched copper substrate (Courtesy of Mr. Katsuhiko Tashiro).

layer is porous. In this example, the 2nd layer is seen to grow in such a way as to fill the pores left in the 1st layer. In electroless plating processes, the film thickens by layer growth. In contrast to the electrolytic plating case, the phenomenon of current density concentration at protrusions is absent. For this reason, the thickness variation is relatively small in electroless films. Figure 2.25 shows the cross-section TEM micrograph of a 28-nm thick electroless Ni– P film grown on an etched pure copper substrate. In this range, the film thickness appears to be very uniform regardless of any surface irregularity on the copper substrate. With increasing film thickness nodules may nucleate at the protrusions, where the flow rate of the solution becomes very rapid (Tashiro et al., 1996). It should be remembered that the protrusions described here are already present on the substrate surface and are not the ones generated during electroless plating.


127

2.2.1.3 Bonding between electroless films and substrate 2.2.1.3.1 Metal substrates. Like electrolytic films, electroless films can grow epitaxially on metal substrates by forming a metal – metal bond (Watanabe and Tanabe, 1974b, 1975, 1976). The bond strength represents the maximum adhesive force of the film to the substrate. The lattice matching relationship of various combinations of metals prepared by an electroless method were presented in Table 1.2. It has been shown that alloying with phosphorus or boron refines the grains of electroless films. Even in this situation, epitaxial growth will be initiated and maintained up to some thickness. An electroless plating system that uses hydrazine or formalin as a reducing agent does not incorporate phosphorus or boron, thus yielding a high-purity metal film with large surface irregularities. In this case, epitaxial growth can be easily achieved. Figure 2.26 is a micrograph showing the surface morphology and internal structure of electroless nickel films grown on a polycrystalline copper sheet using hydrazine as a reducing agent. The surface morphology captured by SEM is shown in Figure 2.26(a), which is magnified in Figure 2.26(a0 ). Two TEM micrographs and their electron diffraction

Figure 2.26. The surface morphology and internal structure of electroless nickel films grown on a polycrystalline copper sheet using hydrazine (NH2NH2) as a reducing agent. (a) and (a0 ) are the low- and high-magnification SEM micrographs, whereas (b) and (b0 ) are the low- and high-magnification TEM micrographs with electron diffraction patterns.

128

Nano-Plating

patterns are shown in Figure 2.26(b) and (b0 ). From the SEM micrographs, the nickel film is seen to have copied the surface structure of the copper substrate and grown preferentially thicker at the grain boundaries and twin boundaries. It is clear from the single-crystal electron diffraction pattern that the film has grown epitaxially on the copper substrate. 2.2.1.3.2 Non-metal substrates. The advantage of using an electroless plating method is in its ability to deposit a metal film on non-metallic (non-conductive) substrates, such as plastic or ceramic materials. The crucial step for successful metallizing of non-metallic substrates is in the proper preparative surface treatment of the substrate prior to metal deposition. The substrate surface is first cleaned and then etched creating microscopic pores on the substrate surface. The etched substrate is then dipped into a tin chloride solution, followed by immersion in a palladium chloride solution. After each immersion, the substrate has to be rinsed with water to remove the solution residue. The substrate is finally immersed in an electroless plating solution for metal deposition. Substrate etching is performed in order to induce surface pores, where the tin chloride solution can be retained. Palladium chloride is subsequently reduced by the retained tin chloride on the pore walls, forming fine palladium particles (colloidal particles (Osaka et al., 1982)), which will be left adsorbed on the pore walls. Figure 2.27 is a schematic diagram illustrating how an electroless film grows on a nonmetallic substrate. Catalytic palladium particles are first formed on the wall of pores induced by etching (cf. Figure 2.27(1)). An electroless metal reduction process initiates on these catalytic palladium particles (cf. Figure 2.27(2)). The electroless metal film initially fills the inner wall space of the pores by nucleating on the palladium particles, then grows out of the pores, and finally covers the whole substrate surface (cf. Figure 2.27(3)). Figure 2.28 (Ono et al., 1998) shows a cross-section TEM micrograph showing the initial stage of electroless copper deposition on an acrylonitrile butadiene styrene (ABS) resin substrate, which was etched and activated with palladium particles following the steps described above. From Figure 2.28(a), palladium particles are observed adsorbed on the wall of pores, which are uniformly covered by electroless copper. Two separate layers of palladium particles and electroless copper can be distinguished in Figure 2.28(b), where electron diffraction patterns from (A) the copper and (B) palladium are also seen. The size of the palladium particles and electroless copper grains can be measured from the highresolution lattice image (see Figure 2.28(b)). These results clearly demonstrate how an electroless film nucleates and grows on non-metallic substrates. We have discussed an electroless metal deposition mechanism on non-metallic substrates. Contrary to the catalytic metal substrates, no chemical bonds exist between an electroless metal film and a non-metallic plastic/ceramic substrate. In principle, therefore, there is no adhesion between an electroless metal film and a non-metallic substrate. It is


129

Figure 2.27. Film formation processes of electroless metals on an etched/activated non-metallic substrate.

equivalent to saying that palladium particles have no chemical bonds to the substrate. As described above, the initial deposit is imbedded and anchored inside the pores. Therefore, a finite amount of force, which exceeds the anchoring strength, is required to separate the film from the substrate. This anchoring force is mechanical in nature and represents the adhesive strength of electroless films to non-metallic substrates. The anchoring force should depend on the shape, depth and density of these pores. It is clear that a high density of pores has to be created on the substrate surface in order to increase the adhesive strength. Control of the pore geometry and density by etching requires the proper selection of the substrate material. The most important condition for this selection is that the substrate material must be structurally non-uniform. In other words, if the substrate material is microscopically uniform, it may dissolve uniformly without forming etch-related pits or pores. Figure 2.29 shows a SEM micrograph showing electroless Ni –B alloy films grown on an etched poly-propene (PP) resin (Watanabe and Tanabe, 1974a). The etched surface of the PP resin is shown in Figure 2.29(0). In Figure 2.29(A), the growth sequence of crystalline Ni –B films containing low B content in the order of (1) – (3) is shown. A similar sequence for amorphous Ni –B (high B content) films is shown in Figure 2.29(B) (1) – (3).

130

Nano-Plating

The top surface covered by the crystalline Ni – B film appears to be rough, whereas that by the amorphous film is smooth. The etched surface of a PP resin and its schematic three-dimensional view are shown in Figure 2.30(a) and (b), respectively. In Figure 2.30(c), the back-side of the Ni – B film grown on the PP resin is seen to contain a number of protrusions or anchors, which are essentially a replica of the pore structure created by etching on the resin surface. Figure 2.31 is an SEM micrograph showing the back-side of an electroless copper film grown on an alumina substrate. In contrast to the back-side in Figure 2.30(c), the anchors of the metal film are sheared off and left on the substrate surface, indicating a strong film adhesion to the substrate. In addition to the etching treatment, various attempts have been made to improve the film adhesion to non-metallic substrates using novel surface treatment techniques. For example, these surface treatments include plasma etching on plastic substrates (Carbonnier et al., 1996), doping on silicon substrates (Gorostiza et al., 1997), or silane coupling (Wasserman et al., 1989).

Figure 2.28. Cross-section TEM micrographs taken at three increasing magnifications in the order of (a) –(c), which show the cross-sectional view of an electroless copper film nucleated on palladium particles adsorbed on an ABS resin (Courtesy of Prof. Sachiko Ono, Kogakuin University).


Figure 2.28 (continued )

131

132

Nano-Plating

Figure 2.29. Growth processes of an electroless Ni –B alloy film on top of an etched poly-propene (PP) resin (marked with (0)). (A) Growth of a crystalline Ni–B film containing a low boron concentration and (B) growth of an amorphous Ni –B film with a high boron content.

2.2.2

Immersion coating

2.2.2.1 Film formation mechanism in immersion (displacement) coating. The immersion method is the simplest coating method, and has increased in popularity in recent years, especially in the microelectronics industry. For example, an immersion zinc coating known as a zincate (Yasuzumi et al., 1996, 1997; Nakata et al., 1997) has


133

Figure 2.30. (a) The etched surface and (b) the schematic view of a PP resin. (c) The backside of an electroless Ni–B film grown on the PP resin.

134

Nano-Plating

Figure 2.31. An SEM micrograph showing the backside of a deliberately peeled copper film grown by an electroless method on an alumina substrate.

been adopted for the manufacture of hard disks. An immersion gold coating has also been applied as a solder joint for electronic devices, such as a condenser, a resistor, or a coil. We will present a brief description of the principle of the immersion coating. Assume that a metal, B, is less noble than a metal, A. If a metal substrate, B, is immersed into a solution containing metal ions, A, the metal, B, will dissolve into the solution and become metal ions. This dissolution process will generate electrons, which will reduce the metal ions, A, in the solution and form adatoms, A, on the substrate surface. In this way, a metal film, A, will be formed on the substrate metal, B. In immersion coating, therefore, the substrate surface must contain both cathodic and anodic sites, which serve for metal deposition and dissolution, respectively. By this technique, the maximum attainable film thickness is limited by the availability of the sites for metal dissolution. In other words, the film growth ceases when the anodic sites are completely covered by the film. Figure 2.32 shows a TEM micrograph showing the growth sequence of immersion gold films grown on pure copper substrates. The gold starts as an island, then forms a network of islands over time, and finally covers the whole surface uniformly. The electron diffraction rings in Figure 2.32(a) clearly indicate that the initial small islands are randomly oriented. It is possible that these small islands are adsorbed physically without strong chemical bonds. For this reason, the small islands are mobile and can take randomly oriented positions without being affected by the substrate structure. As the islands grow in size, they form chemical bonds to the substrate and initiate epitaxial growth. The effect of island size on epitaxial growth was discussed in Section 2.1.1. Epitaxial relationships observed in immersion coatings are listed in Table 1.2. In addition to epitaxial growth, an increase in the film thickness requires a continuous dissolution of the substrate metal. It is of interest to find out which part of the substrate is


Figure 2.32. TEM micrographs showing the growth sequence of immersion gold films grown on pure copper substrates.

135

136

Nano-Plating

dissolving. SEM examinations of these films, however, have been unsuccessful in revealing such sites for metal dissolution. Figure 2.33(A) and (B) shows the surface morphology of thick silver films obtained by an immersion method. This morphology clearly reveals thickness and surface roughness variations on the surface of the silver films (see Figure 2.33(A)). Such variation can be attributed to a difference in the displacement rate, which in turn depends on the orientation of copper grains. It is interesting to note in Figure 2.33(B) that a number of holes are present inside the grain, suggesting possible sites for the metal dissolution. The grains seen in Figure 2.33(A) do not show such holes. 2.2.2.2 Bonding between immersion coating and substrate (Tanabe and Ishibashi, 1967; Tanabe and Urao, 1968; Tanabe et al., 1968a,b; Tanabe and Kamasaki, 1970). To the author’s knowledge, there is no reported work in the literature, which addresses the adhesion mechanism of immersion coatings. As is clear from the film formation mechanism, the growth of immersion coatings requires the dissolution of the substrate. In other words, part of the substrate becomes sites for metal dissolution and the other part becomes sites for metal deposition. Under these conditions, it is difficult to imagine how one can grow an adherent film. The problem is that no information is available at moment as to which part of the substrate dissolves. Without such information, it is difficult to estimate the adhesive strength of immersion coatings. 2.2.3 Contact plating In the electroless plating technique, an immersion of a metal substrate into a plating solution, containing metal ions of interest, spontaneously initiates metal deposition. In a contact plating solution metal deposition does not occur upon substrate immersion, only commencing if the substrate is contacted by an electrochemically less noble metal. Upon contact, a local battery circuit is created between the substrate and the contacting metal. As the contacting metal dissolves into the solution, noble metal ions in the solution, proportional to the concentration of dissolved metal ions, will deposit on the substrate. Recently, based on reports by Wangyu and Bangwei (1991) and Fujita et al. (1998), we successfully grew amorphous Fe –B films on copper substrates using aluminum as a contacting metal and were able to study the formation mechanism together with their magnetic properties. The principle of contact plating is different from that of immersion plating. One major difference is that in the contact plating method, the substrate does not dissolve into a plating solution, and thus provides an excellent surface for film adhesion. Contrary to immersion coatings, there is no thickness limitation during film growth. It should be remembered that the contacting metal itself continues to dissolve during metal deposition and increases its ion concentration in the solution. The contacting metal ions, however, are


137

Figure 2.33. The surface morphology of thick silver films obtained by an immersion method. (A) Plating conditions: a 24-h immersion in a solution containing 7.5 g/l AgCN þ 15.0 g/l NaCN and (B) a 5-day immersion in a solution containing 7.5 g/l AgCN þ 7.5 g/l NaCN.

138

Nano-Plating

not expected to be significantly included into the deposit because they are electrochemically less noble.

REFERENCES Carbonnier, M., Alami, M. & Romand, M. (1996) J. Electrochem. Soc., 143, 472. Diokic, S.S. (1999) J. Electrochem. Soc., 146, 1824. Fujita, N., Lim, P.B., Makino, E., Inoue, M., Arai, K. & Fujii, Y. (1998) J. Surf. Finish. Soc. Jpn, 49, 769. Gerischer, H. (1960) Electrochim. Acta, 2, 1. Gorostiza, P., Diaz, R., Servat, J., Sabz, F. & Morante, J.R. (1997) J. Electrochem. Soc., 144, 909. Itoh, K., Wang, F. & Watanabe, T. (2001) J. Jpn Inst. Metals, 65, 495. Kamasaki, S. (1974) A TEM study on the crystal growth and fine structure of electrolytic gold films, Ph.D. Dissertation, Tokyo Metropolitan University. Nakata, T., Wada, I., Imai, H., Ikejima, K., Inoue, K. & Watanabe, T. (1997) J. Surf. Finish. Soc. Jpn, 48, 820. Narita, A. & Watanabe, T. (1991) J. Surf. Finish. Soc. Jpn, 42, 559. Ohno, I. & Haruyama, S. (1991) J. Jpn Inst. Metals, 55, 736. Ono, S., Osaka, T., Naito, K. & Nakagishi, Y. (1998) J. Surf. Finish. Soc. Jpn, 49, 625. Onoda, M., Shimizu, K., Tuchiya, T. & Watanabe, T. (1995) Mater. Trans. JIM, 36, 1104. Onoda, M., Shimizu, K., Tuchiya, T. & Watanabe, T. (1999) Trans. IMF, 77, 44. Osaka, T., Nihei, K. & Goto, F. (1982) Denki Kagaku, 50, 418. Seiter, H., Fischer, H. & Albert, L. (1960) Electrochim. Acta, 2, 97. Shimizu, Y (1976) A TEM study on the crystal growth, fine structure, and crystalline phases of electrolytic binary alloy films, Ph.D. Dissertation, Tokyo Metropolitan University. Shimizu, Y. & Tanabe, Y. (1976a) J. Metal Finish. Soc. Jpn, 27, 574. Shimizu, Y. & Tanabe, Y. (1976b) J. Metal Finish. Soc. Jpn, 27, 20. Suda, H., Watanabe, T., Misaki, Y. & Tanabe, Y. (1981a) J. Jpn Inst. Metals, 45, 5. Suda, H., Watanabe, T., Misaki, Y. & Tanabe, Y. (1981b) J. Jpn Inst. Metals, 45, 118. Tanabe, Y. & Ishibashi, K. (1967) J. Metal Finish. Soc. Jpn, 18, 41. Tanabe, Y. & Kamasaki, S. (1970) J. Metal Finish. Soc. Jpn, 21, 281. Tanabe, Y. & Urao, K. (1968) J. Metal Finish. Soc. Jpn, 19, 265. Tanabe, Y., Urao, K. & Ogura, T. (1968a) J. Metal Finish. Soc. Jpn, 19, 2611. Tanabe, Y., Urao, K., Ogura, T. & Kurihara, T. (1968b) J. Metal Finish. Soc. Jpn, 19, 217. Tanabe, Y., Hasegawa, N. & Kodaka, M. (1983) J. Metal Finish. Soc. Jpn, 34, 452. Tanabe, Y., Kakegawa, M. & Shimizu, Y. (1986) J. Mater. Sci. Soc. Jpn, 13, 255. Tashiro, Y., Chiba, K., Fukuda, Y., Nakao, H. & Watanabe, T. (1996) J. Surf. Finish. Soc. Jpn, 47, 349. Vetter, K.J. (1961) Electrochemische Kinetik. Springer Verlag, Berlin, p. 698. Wang, F., Itoh, K. & Watanabe, T. (2000) J. Jpn Inst. Metals, 64, 1133. Wangyu, H. & Bangwei, Z. (1991) Physica, B175, 396. Wasserman, S.R., Tao, Y.T. & Whitasides, G.M. (1989) Langmuir, 5, 1074. Watanabe, T. (1989) J. Surf. Finish. Soc. Jpn, 40, 375. Watanabe, T. (1990a) J. Surf. Finish. Soc. Jpn, 41, 652. Watanabe, T. (1990b) J. Surf. Finish. Soc. Jpn, 41, 349.


139

Watanabe, T. (1993) Current Topics in Amorphous Materials, Physics & Technology. NorthHolland, p. 137. Watanabe, T. & Tanabe, Y. (1974a) J. Metal Finish. Soc. Jpn, 25, 87. Watanabe, T. & Tanabe, Y. (1974b) J. Metal Finish. Soc. Jpn, 25, 36. Watanabe, T. & Tanabe, Y. (1975) J. Jpn Inst. Metals, 39, 1. Watanabe, T. & Tanabe, Y. (1976) Trans. JIM, 17, 655. Watanabe, T., Naoe, T. & Katsumata, A. (1989) J. Surf. Finish. Soc. Jpn, 40, 458. Watanabe, T., Arai, K., Hirose, T. & Chikazawa, M. (1999a) J. Japan Inst. Metals, 63, 496. Watanabe, T., Hirose, T., Arai, K. & Chikazawa, M. (1999b) J. Jpn Inst. Metals, 63, 489. Winand, R. (1994) Electrochim. Acta, 39, 1091. Yasuzumi, K., Fujishige, Y., Seo, M., Nanis, L., Nakao, H. & Tashiro, T. (1996) J. Surf. Finish. Soc. Jpn, 47, 802. Yasuzumi, K., Fujishige, Y., Seo, M., Saeki, I., Nanis, L., Nakao, H. & Tashiro, T. (1997) J. Surf. Finish. Soc. Jpn, 48, 1019. Zhu, L. & Watanabe, T. (1992) J. Jpn Inst. Metals, 56, 664.


Chapter 3

Plating in Organic Electrolyte 3.1. 3.2. 3.3.

Introduction Experimental methods Selection of plating electrolyte 3.3.1 Surface morphologies of plated films 3.3.2 Crystal structure of plated films 3.4. Cobalt deposition from various organic solvent solutions 3.4.1 Grain-refining effect 3.4.2 Inclusions in plated films 3.4.2.1 H2O bath 3.4.2.2 DMSO bath 3.4.2.3 FA bath 3.4.3 Relationship between solvent decomposition product and grain size 3.4.4 Surface morphology of plated cobalt films 3.4.5 Crystal orientation and form References

143 144 146 146 150 152 161 162 162 163 165 166 168 169 170


Chapter 3

Plating in Organic Electrolyte 3.1. INTRODUCTION

Most industrial electroplating processes are performed in aqueous solutions. The types of metals that can be electrodeposited from an aqueous electrolyte, however, are limited (Ishibashi et al., 1971; Takei, 1998). Metals plated from an aqueous solution may contain hydrogen, which affects their mechanical properties adversely. Hydrogen inclusion is known to cause the development of high internal stresses and cracks in plated films. It is therefore ideal to conduct plating operations in hydrogen- or oxygen-free electrolytes. As an alternative to aqueous solutions, plating in non-aqueous electrolytes (organic media) or in fused salts has been previously investigated. A number of review articles on plating in non-aqueous electrolytes have appeared in the literature, typically papers by Brenner et al. (Brenner, 1956, 1959; Conner and Brenner, 1956; Reid et al., 1957; Wood and Brenner, 1957) and by Japanese investigators (Hayashi and Ishida, 1962; Takahashi and Nomura, 1962; Hayashi, 1965; Hayashi et al., 1965; Baba, 1974; Yoshio et al., 1976; Takei, 1990). Metals described in these reviews represent metals that cannot be electrodeposited in aqueous solutions, i.e. rare-earth metals (Usuzuka et al., 1988; Matsuda et al., 1992; Sato, 1995; Sato et al., 1995), Al (Biallazor and Lisowska-Oleksiak, 1990; Yoshio, 1986; Matsuda et al., 1992; Takei, 1998), and Mo (Turu et al., 1995), and metals that readily absorb hydrogen, i.e. Ni (Hamashima et al., 1970; Takei, 1974a –c) and Cr (Tajima et al., 1970; Turu et al., 1977; Turu et al., 1980; Turu et al., 1982a –c; Turu et al., 1983a – c). Baba et al. (1971) studied 10 different metals (Cd, Ni, Pd, etc.) obtained from a dimethyl formamide (DMF) solution. Other studies involve Cd, Sn, and Pb deposits from a DMF (Baba et al., 1993) solution, Cd –Hg – Te ternary alloy deposits from a propylene carbonate solution (Colyer and Cocivera, 1992), Nd –Li alloys from a nonaqueous solution (Nakamura et al., 1992), Cu (Takei, 1973, 1974d), Pb (Kimura et al., 1976), and Au (Kimura et al., 1995) deposits from a methanol solution. These papers primarily concern reactions occurring at the electrode, while very little emphasis has been placed on the film composition and its crystallographic structure. A large number of metal systems have been deposited in non-aqueous solutions. A thorough review of all these plating systems is beyond the scope of this book. Therefore, we (Suzaki and Watanabe, 2000) will only present a detailed account of selected plating systems, specifically iron-group metals (Fe, Co, and Ni). Iron-group metals, whose mechanical properties are very sensitive to hydrogen inclusions, are generally plated in aqueous solutions. We will explore what kind of organic media can be used to electrodeposit these metals. 143

144

Nano-Plating

Organic solvents used in this study are conveniently divided into 4 types. The addition of a water-based solution to these solvent systems produces 13 different electrolytes, which were used to test the feasibility of depositing Fe, Co, and Ni metals. We will first discuss the possibility of plating these metals in organic media, and then present the fundamental principles of the plating method along with detailed data on the morphology and structure of these films.

3.2. EXPERIMENTAL METHODS

The solvents used in this study are listed in Table 3.1. The 13 organic solvents are categorized into 4 types (protic neutral, protic base, aprotic neutral, and aprotic base) and their properties (dielectric constant and water content) are also tabulated. The water content is measured after first subjecting the commercial solvents to treatment with a molecular sieve (Molecular Sieve 4A, manufactured by nakalai tesque). The molecular sieve was first dehydrated by heating at 703 K (430 8C) in vacuum (0.133 Pa) for 2 h. A 4-g molecular sieve was placed into a beaker containing 500 ml solvent, and stored in a shaker for several days. The solvent’s residual water content is detailed in Table 3.1. Water content was measured using a water content micro-analyzer, Model AQ-6 (Hiranuma Sangyo Co., Ltd.). Rolled sheets of polycrystalline copper were used as the substrates. They were annealed in vacuum (0.133 Pa) at 873 K (600 8C) for 1 h, followed by furnace cooling. To define the area for metal deposition, the copper sheet was masked except for the plating area

Table 3.1. Organic solvents used in this experiment. Classification of solvent

Solvent

Symbol

Relative dielectric constant

Amount of residual water

Protic and amphoteric

Water Ethanol N-methylformamide Formamide Acetone Ethyl acetate Benzene Dimethylsulfoxide N,N-dimethylforamide Acetonitrile Pyridine Tetrahydrofuran Di-n-butyl ether

H2O EtOH N-MFA FA Ac EA Be DMSO DMF AN Py THF DBE

78.5 24.55 182.4 111 20.72 6.02 2.28 48.9 36.7 37.5 12.3 7.58 3.06

– 0.1% 150 ppm 600 ppm 0.3% 500 ppm 10 ppm 0.1% 500 ppm 250 ppm 10 ppm 700 ppm 100 ppm

Protic and basic Uprotic and neutral

Aprotic and basic

145

Plating in Organic Electrolyte

ð12 £ 20 mm2 Þ and a terminal region. Here, special care was taken to prevent masking materials from dissolving in organic media. Seven types of masking materials were evaluated as listed in Table 3.2. Except for a commercial sealant called Bath-Caulk (Cemedine Co., Ltd.), all masking materials dissolved in all five solvents studied. This particular sealant, Bath-Caulk, was highly viscous and thus it was necessary to dilute it with toluene for easy application. After the application of the Bath-Caulk sealant to the copper sheet substrate, a deformed surface layer left by rolling on the exposed plating area was removed and then mirror-finished by electropolishing. A metal film was deposited on the polished part of the copper sheet using a platinum sheet as an anode. Electrodes and a thermometer were attached to a Pyrex glass-plating tank, equipped with lid. To remove dissolved oxygen from the plating solution, nitrogen gas was bubbled through the solution during electrodeposition. The whole plating tank was kept in a dry box filled with nitrogen gas. In this way, the entrance of water and oxygen from air into organic solvent solutions was minimized. The solution temperature was 298 K ð25 8CÞ ^ 1:5 K: Anhydrous metal salts FeCl2, CoCl2, and NiCl2 were dried in silica-gel desiccators and used without any further dehydration treatment or refining. In this experiment, a small amount of water is expected to come from water contained in the metal salt as well as the residual water left in the solvent. The concentration of the salt was 0.1 mol/l and the bath volume was 200 ml. No other chemicals were added to the solution to minimize the introduction of impurities, which in turn affect the morphology and structure of plated metals. The plating was conducted at a constant current density with the total amount of electric charge kept at 18 C. After plating the films were washed in ethanol (EtOH) and dried by blowing with pressured nitrogen gas. The surface and cross-section of the plated films were examined using SEM. For this observation, cracks were deliberately introduced by bending the films by 1808 while attached to the substrate. In this way, the crack surface (cross-section) and the top surface can be imaged simultaneously. For a structural analysis, an X-ray diffraction method was adopted using the cobalt a-line and an iron filter. Table 3.2. Suitability of masking materials in various organic solvents. Masking agent

N-MFA

DMSO

DMF

Py

THF

Masking sol Sony bond for wood Sony bond vinyl Sony bond leather Sony bond clear Bond for styrene foam Bascoke

K K W W £ £ W

£ £ W W £ £ W

K W £ £ £ £ W

£ W £ £ £ £ W

K £ £ £ £ £ W

W: no change; D: peal off; £ : dissolve.

146

Nano-Plating

3.3. SELECTION OF PLATING ELECTROLYTE

The 13 organic solvents were evaluated for their ability to plate Fe, Co, and Ni metals (see Table 3.1). The solvent’s capability to plate metals was determined using the following three criteria: (1) A sufficient amount of metal salt must dissolve in the solvent. (2) The solution must be stable over an extended period of time. (3) Metal deposition must occur upon the passing of a current. In Table 3.3 (a) – (c), we show plating conditions for metal salts, FeCl2, CoCl2, and NiCl2, respectively. In this table, we list the type of solvents, their dielectric constant, the amount of the residual water, the solubility of the metal salt, the solution color, the feasibility of electrodeposition, the limiting current density, the condition of the deposit, the color of the deposit, and the presence/absence of gas evolution. In this experiment, metal salts (FeCl2, CoCl2, and NiCl2) were dissolved in 200 ml solvent to make a concentration of 0.1 mol/l. The plating was conducted at current densities of 25 and 250 A/m2. No unusual heat evolution or reaction was observed when these salts were dissolved in the various solvents listed in the table, except for the case of addition of FeCl2 to the pyridine (Py) solvent, which caused the appearance of white smoke. From Table 3.3, we note that the kind of platable metals depends on the type of the solvent used. For example, an iron metal can be deposited in all solutions except ethyl acetate (EA), Py, and tetrahydrofuran (THF) solvent baths. Cobalt and nickel metals, on the other hand, can be deposited in all solutions, except EA and THF baths. 3.3.1 Surface morphologies of plated films We first conducted feasibility experiments to determine which deposit – solvent combination(s) produce an acceptable electrodeposit and then studied the surface/crosssection of the deposits using SEM. We have selected one set of the deposit –solvent combinations and shown their SEM micrographs in Figure 3.1. The surface morphology and cross-section of plated films were found to change with the deposit –solvent combination. All three metals (Fe, Co, and Ni) obtained from the EtOH bath showed a rough surface and their fracture surfaces were non-uniform with high porosity. The surfaces of cobalt and nickel films obtained from the formamide (FA) bath were very smooth and bright. Iron and nickel films from dimethyl sulfoxide (DMSO) also showed a smooth surface. The surface of the cobalt film from the DMSO bath was very rough and contained

Table 3.3. Summary of the feasibility of plating iron-group metals in various organic solvent baths. Classification Solvent of solvent

(b) CoCl2 Protic and amphoteric Protic and basic Uprotic and neutral Aprotic and basic

Probability of Limit of current Form of electrodeposition density for plating deposit

Color of deposit

Evolution of gas

Luster film Lusterless film Lusterless film Luster film Lusterless film Powder Luster film Luster film Powder A very small Powder

Silver Black Silver Black Black Brown Silver Black Brown Yellow

W £ W £ £ £ £ £ £ £

Black

£

Dendrite Lusterless film Lusterless film Luster film Luster film … … Lusterless film Powder A very small Lusterless film

Black Gray Gray Gray Gray-brown … … Black Brown Yellow

£ £ W W £ £ £ £ £ £

Gray-light blue

£

Water 78.5 EtOH 24.55 FA 111 N-MFA 182.4 Ac 20.7 EA 6.02 DMSO 48.9 DMF 36.7 AN 37.5 Py 12.3

… 0.1% 600 ppm 150 ppm 0.3% 500 ppm 0.1 % 500 ppm 250 ppm 10 ppm

K W K W W K K W K K

Light yellow Red brown Brown Deep brown Deep brown Brown Red brown Deep brown Light yellow Yellow

Possible Possible Possible Possible Possible Possible? Possible Possible Possible Possible?

… About 300 A/m2 … … About 700 A/m2 About 20 A/m2 … About 400 A/m2 About 200 A/m2 About 3 A/m2

THF

7.58

700 ppm

W

Deep brown

Possible

About 100 A/m2

Water 78.5 EtOH 24.55 FA 111 N-MFA 182.4 Ac 20.7 EA 6.02 Be 2.28 DMSO 48.9 DMF 36.7 Py 12.3

… 0.1% 600 ppm 150 ppm 0.3% 500 ppm 10 ppm 0.1 % 500 ppm 10 ppm

K K W W W K £ K K K

Light red Deep blue Red purple Deep blue Deep blue Light blue Transparency Deep blue Deep blue Red purple

Possible Possible Possible Possible Possible Impossible Impossible Possible Possible Possible

… 300–400 A/m2 … … 400–500 A/m2 … … … … 2 A/m2

THF

700 ppm

K

Deep blue

Possible

About 10–20 A/m2

7.58

147

(continued on next page)


(a) FeCl2 Protic and amphoteric Protic and basic Uprotic and neutral Aprotic and basic

Relative Amount of Solubility Caller of dielectric residual of FeCl2 solution constant water

148

Table 3.3. (continued ) Classification Solvent of solvent

DBE

Probability of Limit of current Form of electrodeposition density for plating deposit

Color of deposit

Evolution of gas

3.06

100 ppm

£

Transparency Impossible

…

…

…

£

Water 78.5 EtOH 24.55 FA 111 N-MFA 182.4 Ac 20.7 EA 6.02 DMSO 48.9 DMF 36.7 AN 37.5 Py 12.3 THF 7.58

… 0.1% 600 ppm 150 ppm 0.3% 500 ppm 0.1% 500 ppm 250 ppm 10 ppm 700 ppm

K K W K K £ W K K K K

Green Yellow Green Green yellow Transparency Transparency Green Deep brown Light blue Green Light orange

… 150–250 A/m2 … … About 0.6 A/m2 0.01 A/m2 … About 300 A/m2 10 A/m2 About 0.5 A/m2 About 0.15 A/m2

Lusterless film Powder No bright film No bright film Powder? … Lusterless film Lusterless film Powder … …

Black Black Gray-black Black White … Black Black Yellow … …

W £ W W £ £ £ £ £ £ £

Possible Possible Possible Possible Possible? Impossible Possible Possible Possible? Impossible Imossible

Solubility of (a) FeCl2/(b) CoCl2/(c) NiCl2 and their platability are discussed in terms of solubility of metal salts, solution color, feasibility of metal deposition, limiting current density, deposit color, and gas evolution. Solubility: K, quickly soluble; W, soluble after long time; K, some part remains; £ , not soluble. Mark “…” in limit of current density means that the limit is over 500 A/m2.

Nano-Plating

(c) NiCl2 Protic and amphoteric Protic and basic Uprotic and neutral Aprotic and basic

Relative Amount of Solubility Caller of dielectric residual of FeCl2 solution constant water


Figure 3.1. Surface morphology and fracture surface for Fe, Co, and Ni films plated from three kinds of organic solvent baths.

149

150

Nano-Plating

characteristic irregularities/lumps, which appeared to have induced a cleavage-type crack on the fracture surface. From the above results, the surface morphology and cross-section of plated iron-group metals were shown to change with the deposit – solvent combination. The surface of the metal films from all the tested organic solvent baths, except the H2O and EtOH baths, were generally smooth with only minor surface irregularities.

3.3.2

Crystal structure of plated films

Figure 3.2 shows the corresponding X-ray diffraction patterns of the plated films seen in Figure 3.1. Note that all the patterns are superimposed with diffraction peaks from the copper substrate. Although these superimposed peaks made the analysis somewhat awkward, the following results were obtained.

Figure 3.2. X-ray diffraction patterns for Fe, Co, and Ni films plated at 25 A/m2 from three kinds of organic solvent baths.


151

The intensity of X-ray diffraction peaks from iron films increases in the order of EtOH, FA and DMSO baths, indicating that the corresponding grain size also increases in this order. The background intensity of these iron films appears to become higher toward the low-angle region. This high background intensity, which made the diffraction peaks less visible, is due to the oxidation of the iron films. The oxidation was especially prominent for the films grown from the H2O and EtOH baths. The rough surface of these films may have made the films more susceptible to oxidation due to an increase in the surface area. X-ray diffraction patterns of iron films from the FA and DMSO baths did not exhibit high background intensity. In some of these oxidation-free iron films, we noted line broadening or the absence of diffraction peaks. These iron deposits were thick enough to produce a strong diffraction pattern as noted in Figure 3.1. Therefore, the absence of diffraction peaks cannot be attributed to an insufficient deposit. We believe that the diffraction peaks were not absent but rather their intensity had become diminishingly weak as the grain size decreased. This grain-size decrease is indeed consistent with line broadening. A careful examination of these diffraction patterns also revealed that the exact location of the broadened peaks did not shift from the peak of the bcc iron crystal. This evidence confirms that the films maintained the bcc structure of the iron crystal, while their grain size decreased. This behavior is different from that of cobalt or nickel films, and will be discussed later. X-ray diffraction patterns of cobalt films from the EtOH, FA, and DMSO baths showed the hcp structure, although the angles of the peaks did not necessarily coincide with that of the hcp structure. In particular, the {002} peak of cobalt films from the FA bath shifted toward the lower angle side and approached the angle of the {111} in the fcc structure. However, no {200} peak for the fcc structure appeared at 55.88. From these results, we conclude that these cobalt films do not contain the fcc crystals but contain the hcp crystals with distorted {002} planes. Both the {100} and {101} diffraction peaks of the cobalt films increasingly broadened and their intensity decreased in the order of the EtOH, FA, and DMSO baths. The grain size is also considered to decrease in this order. All the diffraction lines of nickel films broadened, suggesting that these films are finegrained or close to an amorphous structure. In Table 3.4, we summarize crystal data for iron-group metals (Fe, Co, and Ni) grown from all the solvents used in this study. Except for cobalt films from the H2O bath, all films retained their equilibrium structure (Fe: bcc, Co: hcp, and Ni: fcc) at ambient temperatures. Cobalt films from the H2O bath, however, took not only the hcp structure at ambient temperatures, but also the fcc structure stable at high temperatures. Concerning the presence of the fcc high-temperature phase, Nakahara and Mahajan (1980) concluded that Hþ present in the H2O bath is reduced to atomic hydrogen, which was incorporated into the cobalt lattice, causing the expansion of the {001} lattice. Cobalt films from the organic solvent bath, FA, however, did not show the presence of the fcc phase.

152

Nano-Plating

Table 3.4. The structure of metals obtained from various solvent baths. Classification of solvent

Solvent (symbol)

Solution of FeCl2

Solution of CoCl2

Solution of NiCl2


Water EtOH N-MFA FA Ac EA DMSO MeCN DMF Py THF

bcc Amorphous? Amorphous bcc (F.G.) bcc(F.G.) þ ? £ bcc (F.G.) Amorphous Amorphous £ ?

hcp þ fcc hcp hcp hcp hcp (F.G.) – hcp (F.G.) Amorphous hcp £ Amorphous

fcc Amorphous fcc (F.G.) fcc (F.G.) £ – fcc (F.G.) £ fcc (F.G.) – –

Protic and basic Aprotic and neutral Aprotic and basic

F.G.: fine crystal; £ : a little deposition; –: no deposition; ?: unknown because diffraction is very weak.

As seen in Figure 3.2, the {002} peak of cobalt films from the FA bath shifts toward the lower angle side and moves close to the {111} peak of the fcc structure. However, the {200} peak of the fcc structure did not appear. Based on these results, it can be concluded that the cobalt films from the FA organic solvent bath do not contain the fcc crystals. All iron-group metal films obtained from the organic solvent baths used in this study resulted in the formation of fine grains. The grain-refining phenomenon was especially prominent for the films grown from the DMSO and DMF baths. No useful films were obtained from the aprotic neutral solvent. For example, cobalt films from the acetone (Ac) bath and iron films from the EA bath were brown-colored nonmetallic deposits. Deposits from the Py bath were yellow and those from the THF bath were blue. These non-metallic-colored deposits were probably their hydroxides or solventrelated compounds. Deposits from all other solvents had a metallic luster, although their surfaces were very rough. In Table 3.3, we summarize the above results in terms of the feasibility of electrodepositing 3 iron-group metals from 11 solvents. The surface morphology and crystal structure are summarized in Table 3.4. In terms of our solvent classification, we found no systematic trend for the type of platable metals and their structure. Nevertheless, the grain size was refined in all the solvent baths tested here.

3.4. COBALT DEPOSITION FROM VARIOUS ORGANIC SOLVENT SOLUTIONS

We have discussed the feasibility of depositing iron-group metals (Fe, Ni, and Co) in various organic solvent solutions. We will further describe the structure of cobalt films grown from six electrolytes including a water-based solution. One of the unique properties

153


of cobalt metal is that it undergoes an hcp ! fcc phase transformation at 723 K (450 8C). In addition, the high-temperature fcc phase is known to form in cobalt electrodeposits from a low-pH aqueous solution. This anomaly was shown to originate from hydrogen inclusions (Nakahara and Mahajan, 1980). The structure of cobalt metal is apparently very sensitive to hydrogen. Baba et al. (1993) electrodeposited cobalt films using DMF as a solvent and described the dendritic growth. Electrolytes used in this experiment include five types of organic solvents plus water (see Table 3.5). As described in Section 3.3.2, these organic solvent solutions can be further grouped into three types (protic neutral, protic base, and aprotic base). Although aprotic neutral media (Ac and EA) were also evaluated for depositing cobalt films, they did not yield useful deposits. With the exception of distilled water, we used commercial superhigh-grade reagents for the solvents, which were dehydrated using a commercial molecular sieve (Molecular Sieve 4A, manufactured by nakalai tesque). We could not dehydrate the solvents completely and list the residual water in Table 3.5. Anhydrous CoCl2 was dried over silica gel and added as a metal salt to the solvents. The metal salt concentration was 100 mol/l in a 200 ml solution. No other chemicals were added to the solution. Figure 3.3(A)– (F) are (a) X-ray diffraction patterns and (b) SEM images showing the surface morphologies/cross-section of cobalt films obtained from the H2O, EtOH, FA, N-methylformamide (N-MFA), DMF, and DMSO baths, respectively. These patterns are displayed as a function of current density and their corresponding SEM images are indicated with arrows. To avoid complication in the analysis of X-ray diffraction patterns, we eliminated the diffraction peaks of the copper substrate from all the patterns. It is clear from the X-ray diffraction patterns in Figure 3.3 that cobalt films grown from each organic bath take the hcp structure. X-ray diffraction patterns of cobalt films from the H2O bath contained the characteristic {200} diffraction peak of the fcc crystal over all the current densities used (see Figure 3.3(A) (a)).

Table 3.5. The type of solvent baths used for plating cobalt films. Classification of solvent

Solvent

Symbol

Relative dielectric constant

Amount of residual water


Water Ethanol Formamide N-methylformamide N,N-dimethylformamide Dimethylsulfoxide

(A) H2O (B) EtOH (C) FA (D) N-MFA (E) DMF (F) DMSO

78.5 24.55 111 182.4 36.7 48.9

– 0.1% 600 ppm 150 ppm 500 ppm 0.1%

Protic and basic Aprotic and basic

154

Nano-Plating

Figure 3.3. (A) Cobalt films grown from the H2O bath: (a) X-ray diffraction patterns and (b) SEM images for the films plated at various current densities.


155

Figure 3.3. (B) Cobalt films grown from the EtOH bath: (a) X-ray diffraction patterns and (b) SEM images for the films plated at various current densities.

156

Nano-Plating

Figure 3.3. (C) Cobalt films grown from the FA bath: (a) X-ray diffraction patterns and (b) SEM images for the films plated at various current densities.

In Figure 3.4, we plotted the current efficiency of cobalt deposition for all the baths as a function of current density. We calculated the current efficiency from the weight of the deposit by assuming that the cobalt ions were charged as Co2þ. The current efficiency is high at low current densities but decreases with increasing current density. The maximum


157

Figure 3.3. (D) Cobalt films grown from the N-NFA bath: (a) X-ray diffraction patterns and (b) SEM images for the films plated at various current densities.

158

Nano-Plating

Figure 3.3. (E) Cobalt films grown from the DMF bath: (a) X-ray diffraction patterns and (b) SEM images for the films plated at various current densities.


159

Figure 3.3. (F) Cobalt films grown from the DMSO bath: (a) X-ray diffraction patterns and (b) SEM images for the films plated at various current densities.

160

Nano-Plating

Figure 3.4. A change in the current efficiency with increasing current density during cobalt electrodeposition in various organic solvent baths.

current efficiency for all the baths was generally obtained at a current density less than 70 A/m2. In Figure 3.5, we plotted the grain size determined from the full width at half maximum (FWHM) of X-ray peaks shown in Figure 3.3. The grain size can be calculated using Scherrer’s equation. In this calculation, we used the {100} peak of the hcp structure and the {111} peak of the fcc crystal. For some hcp crystals, we used the {101} peak. According to the SEM images in part (b) of Figure 3.3(A), the film thickness of the cobalt film from the H2O bath is thick at low current densities and thin at high current densities. A similar trend can be observed for the films from the EtOH bath (see part (b) of Figure 3.3(B)). The surface of the film obtained at high current efficiencies was smooth, whereas those at low current efficiencies were rough with leaf-shaped surface morphologies. As shown in Figure 3.5, both the hcp and fcc phases appeared in the cobalt deposits from the H2O bath. An increase in the hcp grain size appears to coincide with a decrease in the fcc grain size. Above a current density of 100 A/m2, the grain size became constant (, 25 nm). The grain size of cobalt films grown from all the organic baths decreased with decreasing current efficiency. These films may have grown as an amorphous phase, an assembly of extremely small grains close to the amorphous phase, or a mixture of the two.


161

Figure 3.5. A change in the grain size of electrolytic cobalt films with increasing current density from various organic solvent baths.

3.4.1 Grain-refining effect It is generally believed that plated films can become fine grained under conditions of high overpotential (high current density), whereas they grow coarse-grained at low overpotentials (Seizan and Ito, 1974; Jokoh et al., 1988; Ohno, 1988; Ohno and Haruyama, 1991; Yamashita, 1993). The grain size of cobalt films obtained from the H2O bath, however, did not become smaller with increasing current density but reached a constant value of 25 nm. In the case of organic baths, the cobalt grain size decreased with increasing current density, which is in agreement with predictions made by the overpotential theory. The observed grain refining in the cobalt films from the organic baths may be connected to the decreased current efficiency with increasing current density. In other words, the decomposition of solvent molecules and the subsequent incorporation of the solvent decomposition products (SDPs) into the deposit must have taken place. Consequently, the SDPs helped refine the cobalt grains. This idea is based on our theory that, The structure of plated films is determined by the type of included elements and their composition. For example, the grains of Ni –P (Watanabe and Kanayama, 1989), Ni – S (Narita and Watanabe, 1991), and Ni– B (Onoda et al., 1990, 1999) are increasingly refined as the

162

Nano-Plating

Figure 3.6. Current–potential curves taken in the (a) H2O and (b) the FA baths containing 100 mol/l CoCl2.

concentration of the incorporated elements, P, S, and B, is increased. In this study, the decreased current efficiency is due to the increased incorporation of SDPs, which contributed to the grain-refining effect. Cyclic voltammetry was applied to understand the decomposition of solvent molecules during electrodeposition. Figure 3.6 is a current –potential curve for the H2O and FA baths containing 100 mol/l CoCl2. The reduction potential of both cobalt and copper (substrate) can be seen in both curves. From these curves there is no indication of solvent decomposition at a particular potential, instead it is seen to occur across the full range of potentials. In addition, when we measured the current – potential characteristics for the other solvent baths we obtained a similar curve. 3.4.2 Inclusions in plated films 3.4.2.1 H2O bath. The fcc phase, which is unstable at ambient temperatures but stable at high temperatures, only appeared in cobalt deposits from the H2O bath. The presence of the fcc high-temperature phase was previously attributed to the formation of the {111} fcc lattice by the expansion of the {001} hcp lattice, which was assisted by the incorporated atomic hydrogen (Nakahara and Mahajan, 1980). In this study, the fcc phase was also seen in the H2O bath containing a high concentration of Hþ ions. The hydrogen ions were reduced to atomic hydrogen and subsequently incorporated into cobalt deposits. The size of the fcc crystals were small from the high-efficiency bath, in which hydrogen evolution was low, but were large from the low-efficiency bath. In search of hydrogen inclusions, we also examined X-ray diffraction patterns of cobalt films from all the organic solvent baths. Although the peaks of the hcp phase


163

were present, the fcc’s characteristic {200} peak located at the angle of 60.68 was not present. In conclusion, the presence of hydrogen inclusions in cobalt films from the H2O bath was confirmed but not detected in the films from all the organic solvent baths used here. 3.4.2.2 DMSO bath. From a consideration of the observed grain-refining effect, it is possible that some kinds of chemical species could be included in the cobalt films. We applied an X-ray photoelectron spectroscopy (XPS) technique to detect such inclusions in cobalt films grown from the DMSO bath. In Figure 3.7, we show the XPS spectra of sulfur atoms in cobalt films as a function of current density. XPS measurements on these cobalt specimens were performed after argon etching in the XPS chamber for 5 min. Thus the spectra for the sulfur atom did not originate from the film surface but came from the internal region. The peak height is seen to increase as the current density is raised from 15 to 37.5 A/m2, while the current efficiency decreased. Under conditions of high current density and low current efficiency, more decomposition products are formed and subsequently incorporated into the deposits. This trend supports the idea that the inclusions helped refine the cobalt grains. A further increase in the current density to 50 A/m2 was accompanied by a decrease in the intensity of the sulfur peak. Here, the decreased sulfur intensity is not due to a decrease in the amount of the inclusions, but the emission efficiency of the photoelectrons decreased because of the development of surface roughness in the film. The sulfur peak of the CH3 –SO –CH3 (DMSO) solvent molecule is shifted toward the high-energy side (see Figure 3.7). This peak coincides with the energy position (Baker and Betteridge, 1975) of sulfur in the molecule of CH3 – S– CH3, which is the reduced form of the solvent molecule. An additional peak is present between this reduced molecule and the solvent molecule. This peak roughly corresponds to the value (164.0 eV) (Chastain et al., 1992) of the binding energy of a sulfur molecule (S –S bond). From these results, we conclude that both CH3 – S –CH3 and sulfur molecules are included in the cobalt films. To further investigate the binding state of sulfur in the cobalt films, we annealed the films at 873 K (600 8C) for 1 h in vacuum. This annealing not only reduced the peak, but also shifted it towards the high-energy side. It can be shown that the quantity of sulfur atoms was decreased by annealing and, at the same time, the sulfur atoms were further oxidized. The peak of the oxidized sulfur corresponds to the S– S bond energy (Chastain et al., 1992). From these results it can be concluded that annealing oxidizes sulfur molecules in the cobalt films, and the majority of them are expelled outside the film. The remaining sulfur

164

Nano-Plating

Figure 3.7. XPS spectra taken from cobalt films plated from the DMSO bath at various current densities.

segregates as a sulfur molecule, while the grain size of the cobalt films is increased by the annealing. An X-ray or electron diffraction analysis of the cobalt films often revealed one unknown peak (see a peak marked with a question mark “?” in Figure 3.3(F)). To identify this peak, the binding energy of cobalt was measured using an XPS technique (see Figure 3.8(F)). The cobalt peak was located at the same place as that from the H2O bath and did not show a peak shift. The cobalt atom retained a pure metallic state (780.4 eV) (Chastain et al., 1992) without binding other atomic elements. Although the unknown diffraction peak remains unidentified, it does not originate from cobaltrelated compounds, such as cobalt oxides.


165

Figure 3.8. XPS spectra taken from cobalt films plated from (A) the H2O, (C) the FA, and (F) the DMSO baths.

3.4.2.3 FA bath. A wavelength dispersive spectrometer (WDS) attached to a SEM was used to map the distribution of nitrogen in cobalt films obtained from the FA bath (see Figure 3.9). Figure 3.9(a) and (b) are the SEM image and its corresponding nitrogen distribution map, respectively. The film was deliberately bent to produce cracks. Nitrogen atoms are distributed uniformly inside the plated film but do not exist in the cracked part, where the copper substrate is exposed. A chemical shift from these films was also studied using an XPS method as seen in Figure 3.8(C). The cobalt peak coincides with the energy position (Chastain and King, 1992) of cobalt metal and thus is not associated with any other atoms besides cobalt atoms. Again, nitrogen in the cobalt is considered to be present in a molecular form or in SDPs.

166

Nano-Plating

Figure 3.9. Cobalt films grown from the FA bath. (a) SEM image and (b) the corresponding nitrogen atom concentration map.

3.4.3 Relationship between solvent decomposition product and grain size So far, we have considered inclusions in cobalt films obtained from the DMSO and FA baths. For the other organic solvent baths we expect that SDPs, molecular sulfur, or molecular nitrogen are probably included in the deposits without forming chemical bonds to the cobalt atoms. Possible SDPs are listed in Figure 3.10. In Figure 3.11, we plotted a bar graph of the grain size taken at the current efficiency of 70% in Figure 3.4 for various organic solvents. The molecular weight of the SDPs is also included in a curly bracket (see Figure 3.11). From this bar graph, we found an interesting trend that The grain size of cobalt films grown from an organic solvent bath, whose decomposition products have a small molecular weight, is large and vice versa.


167

Figure 3.10. The molecular structures of decomposition products from various organic solvents. (A) H2O, (B) EtOH, (C) FAA, (D) N-MFA, (E) DMF, and (F) DMSO.

This result indicates that the grain size is affected not only by the type of inclusions, but also by their molecular weight. These SDPs lying at the grain boundaries help refine the grains by promoting more nucleations.

168

Nano-Plating

Figure 3.11. A bar graph showing the grain size of cobalt films (horizontal axis) for various molecules of decomposition products, whose molecular number is indicated inside a parenthesis.

3.4.4 Surface morphology of plated cobalt films Grain size is often thought to be a reflection of surface roughness. The present experimental results appear to contradict such an idea. The grain size of cobalt films from the H2O bath shown in Figure 3.5 is constant (25 nm) at current densities above 100 A/m2, whereas the surface roughness changes dramatically in this range as seen in part (b) of Figure 3.3(A). In this case, the grain size is not related to the surface roughness. In comparing Figure 3.3(b) with Figure 3.5(A), we can draw similar conclusions for the cobalt films obtained from the other organic solvent baths. In the previous chapters, we have discussed a control method for the surface morphologies of pure electrodeposits grown from an aqueous electrolyte containing no additives such as brighteners. According to theory, the surface morphology is primarily controlled by the current density (overpotential), by the plating temperature, the type of anions, and solution agitation. In the present experiment, chloride anions (Cl2) may contribute to the surface morphology of cobalt deposits, but adsorbed decomposition products from the solvent may also affect the surface morphology. Especially with increasing current density, cobalt films from the FA, N-MFA, and DMSO baths show a number of bubble-shaped protrusions on top of the initial uniform thin film.


169

A cause for the formation of this surface morphology was discussed by Cui et al. (1992), who attributed it to impurity inclusions formed as a result of electrolysis of DMF. They thought that the adsorption of impurities and their reduction at the electrode surface inhibited the metal deposition processes, resulting in the formation of the rough surface. In addition, they discussed the effect of dissolved oxygen, but the effect of dissolved oxygen in our experiment was negligibly small. At the high current density side, cobalt films initially form a continuous film and then develop bubble-shaped protrusions as the whole surface is progressively covered by highly adsorptive/viscous chemical species with a large molecular weight. Discharge processes at high overpotentials are expected to take place preferentially at a less-covered area, where the protrusions will initiate. This process resembles the growth mechanism of dendrites, although the lack of dendrite formation is probably due to the presence of viscous chemical species adsorbed on the surface. 3.4.5 Crystal orientation and form As described above, all the iron-group metal films grown from organic solvents consist of fine crystals. The maximum crystal size of 50 nm was obtained from the EtOH bath. The grain size of all other deposits was generally smaller than 50 nm. The texture of these films was calculated from the intensity ratio of the X-ray diffraction patterns using Willson’s (Willson and Rogers, 1964) equation. The films from the FA bath had the , 001 . texture as indicated by the strong {002} peak. The films from all the other baths had the , 100 . texture. The grain size of cobalt films from the DMF and DMSO baths was also calculated using diffraction peaks from both the {100} texture and the {101} non-texture directions (see Figure 3.5). In the low current density regions, the grain size from the {100} peak is about twice that from the {101} peak. This indicates that the grains were elliptical with the long axis lying along the film normal direction. In summary, we attempted to electrodeposit cobalt films from three types of organic solvent baths (protic neutral, protic base, and aprotic base) and discussed their plating characteristics and the surface morphology/structure of the deposits. The following conclusions were reached: (1) The structure of the plated cobalt films did not show any systematic trend within the organic solvent classification (protic and aprotic). The current efficiency was high at low current densities but low at high current densities. This trend indicates that the quantity of decomposition products from the solvent during electrodeposition increases with increasing current density. (2) The decomposition products are adsorbed on the surface of the cobalt films and affect the surface morphology.

170

Nano-Plating

(3) The decomposition products are incorporated into the deposits. The quantity is low at low current densities and high at high current densities. (4) Inclusions in the deposits include SDPs, nitrogen molecules, and sulfur molecules. These inclusions can decompose upon heat treatment and escape outside the film. (5) Inclusions segregate at the grain boundaries of the cobalt films and help refine the grains. The degree of grain refining depends not only on the quantity of inclusions, but also on their molecular weight. The larger the molecular weight, the stronger the grain-refining effect. (6) No hydrogen inclusions were found in the cobalt films from any of the solvent baths used here. (7) The films from the FA bath had the , 001 . texture and those from all other baths had the , 100 . texture. REFERENCES Baba, N. (1974) Hyomen, 12, 660. Baba, N., Takeuchi, Y. & Morisaki, S. (1971) J. Metal Finish. Soc. Jpn, 22, 175. Baba, N., Takeuchi, Y. & Yoshio, M. (1993) J. Surf. Finish. Soc. Jpn, 44, 2157. Baker, A.D. & Betteridge, D. (1975) New Chemical Series 59, Photo Electron Spectroscopy, Fundamental & Applications, Gendai Kagaku Doujin, p. 1164. Biallazor, S. & Lisowska-Oleksiak, A. (1990) J. Appl. Electrochem., 20, 590. Brenner, A. (1956) J. Electrochem. Soc., 103, 652. Brenner, A. (1959) Electrolysis Org. Solvents, 106, 148. Chastain, J. & King, R.C., Jr. (1992a) Handbook of X-Ray Photoelectron Spectroscopy, Physical Electronics, Inc., p. 60. Chastain, J. & King, R.C., Jr. (1992b) Handbook of X-Ray Photoelectron Spectroscopy, Physical Electronics, Inc., p. 82. Colyer, C.L. & Cocivera, M. (1992) J. Electrochem. Soc., 139, 406. Conner, J.H. & Brenner, A. (1956) J. Electrochem. Soc., 103, 657. Cui, C.Q., Jiang, S.P. & Tseung, A.C.C. (1992) J. Electrochem. Soc., 139, 1535. Hamashima, I., Nishijima, T. & Kojima, R. (1970) Denki Kagaku, 38, 184. Hayashi, T. (1965) Denki Kagaku, 33, 535. Hayashi, T. & Ishida, T. (1962) Denki Kagaku, 30, 552. Hayashi, T., Kuwa, Y., Yoshida, M. & Kikuchi, N. (1965) Denki Kagaku, 33, 584. Ishibashi, N., Hanamura, T., Yoshio, M., Waki, H. & Kiyoyama, T. (1971) Asahi Glass Ind. Promot. Res. Rep., 44, 599. Jokoh, M., Ohno, I. & Haruyama, S. (1988) J. Jpn Inst. Metals, 52, 95. Kimura, T., Kobayashi, S. & Inui, T. (1976) J. Metal Finish. Soc. Jpn, 27, 230. Kimura, T., Kobayashi, S. & Inui, T. (1995) J. Surf. Finish. Soc. Jpn, 46, 1182. Matsuda, Y., Fujii, T., Yoshimoto, N., Morita, M. & Yoshiga, M. (1992) J. Surf. Finish. Soc. Jpn, 43, 37. Nakahara, S. & Mahajan, S. (1980) J. Electrochem. Soc., 127, 283. Nakamura, Y., Uchiyama, M. & Yoshio, M (1992) Proc. 85th Meeting of the Surf. Finish. Soc. Japan, p. 39.


171

Narita, A. & Watanabe, T. (1991) J. Surf. Finish. Soc. Jpn, 42, 559. Ohno, I. (1988) J. Surf. Finish. Soc. Jpn, 39, 149. Ohno, I. & Haruyama, S. (1991) J. Surf. Finish. Soc. Jpn, 30, 735. Onoda, M., Tsuchiya, T., Ogawa, M. & Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 388. Onoda, M., Shimizu, K., Tateishi, Y. & Watanabe, T. (1999) Trans. Inst. Metal Finish., 77, 44. Reid, W.E., Bish, J.M. & Brenner, A. (1957) J. Electrochem. Soc., 104, 21. Sato, K. (1995) J. Surface Finish. Soc. Jpn, 46, 1094. Sato, Y., Ibi, K. & Kobayakawa, K. (1995) Proc. 3rd International Conf. on Rare Earth Development & Applications, Aug. 21 – 25, 1995, p. 143. Seizan, T. & Ito, S. (1974) Denki Kagaku, 42, 206. Suzaki, A. & Watanabe, T. (2000) J. Metal. Soc. Jpn, 64, 869. Tajima, S., Baba, N. & Morisaki, S. (1970) Electrochim. Acta, 17, 184. Takahashi, T. & Nomura, K. (1962) J. Metal Finish. Soc. Jpn, 14, 58. Takei, T. (1973) Nippon Kagaku Kaishi, 73, 1661. Takei, T. (1974a) J. Metal Finish. Soc. Jpn, 25, 343. Takei, T. (1974b) Bull. Chem. Soc. Jpn, 2, 249. Takei, T. (1974c) Bull. Chem. Soc. Jpn, 2, 257. Takei, T. (1974d) Nippon Kagaku Kaishi, 74, 1403. Takei, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 1278. Takei, T (1998) Surface Finishing Handbook, edited by Metal Finishing Soc. Japan, Nikkan Kogyo Shimbun Co., pp. 283– 295. Turu, T., Takenaka, I., Kobayashi, S. & Inui, T. (1977) J. Metal Finish. Soc. Jpn, 28, 85. Turu, T., Kobayashi, S., Nose, T. & Inui, T. (1980) J. Metal Finish. Soc. Jpn, 31, 249. Turu, T., Kobayashi, S., Fujiyama, J. & Inui, T. (1982a) J. Metal Finish. Soc. Jpn, 33, 140. Turu, T., Kobayashi, S., Fujiyama, J. & Inui, T. (1982b) J. Metal Finish. Soc. Jpn, 33, 402. Turu, T., Kobayashi, S. & Inui, T. (1982c) J. Metal Finish. Soc. Jpn, 33, 608. Turu, T., Kobayashi, S., Kusuhara, K. & Inui, T. (1983a) J. Metal Finish. Soc. Jpn, 34, 12. Turu, T., Takenaka, I., Kusuhara, K. & Inui, T. (1983b) J. Meatl Finish. Soc. Jpn, 34, 114. Turu, T., Takenaka, I., Kusuhara, K. & Inui, T. (1983c) J. Metal Finish. Soc. Jpn, 34, 157. Turu, T., Takagi, T. & Kobayashi, S. (1995) J. Surface Finish. Soc. Jpn, 46, 1130. Usuzuka, N., Yamaguchi, H. & Watanabe, T. (1988) Mater. Sci. Engng, 99, 105. Watanabe, T. & Kanayama, T. (1989) J. Surf. Finish. Soc. Jpn, 40, 425. Willson, K.S. & Rogers, J.A. (1964) Tech. Proc. Am. Electroplaters Soc., 51, 92. Wood, G.B. & Brenner, A. (1957) J. Electrochem. Soc., 104, 21. Yamashita, T. (1993) Surface Technology Handbook. Japan Inst. Material Technology, Bikousha Publ. Co., Ltd., p. 38. Yoshio, M. (1986) J. Metal Finish. Soc. Jpn, 37, 367. Yoshio, M., Yamakawa, K. & Ishibashi, N. (1976) Denki Kagaku, 44, 462.


Chapter 4

Microstructural Changes in Plated Films During Heat Treatment 4.1. Introduction 4.2. A change of meta-stable and non-equilibrium phases to stable crystals 4.3. Transformation of amorphous phase to equilibrium phase References

175 176 182 192


Chapter 4

Microstructural Changes in Plated Films During Heat Treatment 4.1. INTRODUCTION

As described in Chapter 1, the structure of a plated film resembles that of a solid obtained by rapid quenching from a high temperature. For this reason, plated films form either an equilibrium crystalline structure or a thermodynamically unstable phase. The latter can be in the form of a meta-stable or amorphous solid. In the case of an equilibrium crystalline structure, heat treatment cannot change the basic crystalline structure, except for metallurgical changes. These changes include the removal of occluded gas (notably hydrogen), a reduction of internal stress, and a microstructural change through recrystallization/grain growth (Dill et al., 1998). For non-equilibrium phases (meta-stable and amorphous solids), heat treatment produces a more stable structure. A structural analysis of meta-stable or non-equilibrium phases can be difficult because they are not found in the standard equilibrium phase diagram or in the JCPDS X-ray diffraction card. We will now describe the difference between meta-stable and non-equilibrium phases. As seen in Figure 4.1, neither phase is thermodynamically stable. (a) Non-equilibrium phases are solids that, with time, change toward a more stable form without the application of an external energy such as heating. In addition, no welldefined phase transformation appears in these phases. (b) Meta-stable phases, on the other hand, can remain in the same state semipermanently without undergoing any phase transformation, as long as no external energy is supplied. Differential scanning calorimetry (DSC) allows us to determine a difference between the two phases in terms of their thermal properties. For meta-stable phases, no sharp exothermic peaks appear at any temperatures, although heat is generated continuously during testing. In contrast, non-equilibrium phases do not exhibit any exothermic reactions at low temperatures; however, a sharp exothermic peak does appear at high temperatures. Meta-stable solids appear in the form of meta-stable or amorphous phases. The formation of a special crystal structure or a supersaturated solid solution is characteristic of non-equilibrium phases. Fine crystallites, with a nanometer dimension, are also called nonequilibrium phases because they tend to grow larger in order to reduce their surface energy. 175

176

Nano-Plating

Figure 4.1. A schematic view showing a difference in energy state between (a) non-equilibrium and (b) meta-stable phases.

4.2. A CHANGE OF META-STABLE AND NON-EQUILIBRIUM PHASES TO STABLE CRYSTALS

As described above, both meta-stable and non-equilibrium phases, which are not included in the equilibrium phase diagram, are known to form in plated films. These phases can be transformed by heat treatment into a more stable form without changing the overall composition. Using Ni –Sn alloy films as an example, we will describe how meta-stable and nonequilibrium solids change into stable phases after heat treatment. Figure 4.2(a) is the binary phase diagram of a Ni– Sn alloy. On the basis of X-ray diffraction results (see Figure 4.3), we constructed a one-dimensional structure diagram of various plated Ni– Sn alloy films over the entire composition range, Figure 4.2(b). According to the equilibrium phase diagram, there is no Sn solubility in Ni at ambient temperatures, but Sn can dissolve up to 11 at.% at 1130 8C. Plated Ni – Sn alloy films can form a supersaturated solid solution containing at least 12.2 at.% Sn at ambient temperatures. Therefore, this solid solution is a non-equilibrium phase. If this film is heated, supersaturated Sn atoms will be rejected from the Ni phase and will segregate at the Ni grain boundaries or form its b phase (Ni3Sn) inside the Ni grains. As a result, the microstructure becomes a two-phase alloy containing Ni and b-phase grains. In plated Ni – Sn alloy films, we discovered a phase called M1 (see Figure 4.2(b)). This phase is meta-stable and is not listed in the standard equilibrium phase diagram. This phase was reported to have a NiAl-type hexagonal structure by Dutta and Clarke (1968), Clarke and Dutta (1971), and other investigators (Wynne et al., 1971; Wilson, 1972; Schubert, 1973; Tanabe and Shimizu, 1975; Bennett and Tompkins, 1976; Tamura and Hosokawa, 1977; Augis and Bennett, 1978; Lo, 1980; Enomoto et al., 1982) (see Figure 4.4). The M1 phase forms in the Sn concentration range of 18 – 60 at.% where b-Ni3Sn, gNi3Sn2, and d-Ni3Sn4 intermetallic compounds are the equilibrium phases. The M1 phase appears in the composition range 30 – 55 at.% Sn. Although the composition of this phase is close to that of the g phase, it shows a wide Sn solubility range, which can be explained

Microstructural Changes in Plated Films During Heat Treatment

177

Figure 4.2. (a) The equilibrium phase diagram of Ni–Sn binary alloys and (b) a phase diagram of electrolytic Ni– Sn alloys.

Figure 4.3. X-ray diffraction patterns from electrolytic Ni–Sn alloy films with various Sn compositions.

178

Nano-Plating

Figure 4.4. Structure of meta-stable phase (M1) found in electroplated Ni –Sn alloy films.

by the lattice occupancy probability of Ni atoms (denoted with a black circle in Figure 4.4) in the M1 phase. In the M1 lattice, Ni atoms can be easily added or removed, allowing the Ni concentration to vary significantly within the M1 phase. Thus, it is clear that the lattice occupancy probability of Ni atoms is 1/2. This compositional variation is also accompanied by a change in the lattice constant. Figure 4.5 is a DSC spectrum of a plated Ni –Sn alloy film having about 41.0 at.% Sn. Here the meta-stable phase appears as a single phase. While heating, the 41.0 at.% Sn film releases heat but does not show any characteristic exothermic peaks. An X-ray analysis of this film before and after annealing at 500 8C, is shown in Figure 4.6. Here only the M1 phase was present in the as-deposited film but after the annealing, only diffraction peaks from the equilibrium g phase remained. Since the M1 phase has a crystal structure similar to that of the g phase, the M1 ! g phase transformation must have taken place without significant diffusion and nucleation processes. Thus, the M1 phase is regarded as the non-equilibrium form of the g phase. A DSC curve of the 50.3 at.% Sn film exhibited a characteristic exothermic peak around 370 8C as seen in Figure 4.5(b). Before heating, only the meta-stable M1 phase was present


179

Figure 4.5. A DSC curve showing the appearance of meta-stable phase (M1) in a plated Ni –Sn alloy film.

Figure 4.6. A structural change by heat treatment in a meta-stable phase (M1) having the composition of Ni–41.0 at% Sn as observed in X-ray diffraction patterns.

180

Nano-Plating

in this film. As seen in Figure 4.7, an X-ray analysis of the film heated at 350 8C just below the exothermic peak temperature indicates the appearance of prominent diffraction peaks from g and d phases. In contrast, the intensity of diffraction peaks from the M1 and unknown phases (denoted with ?) have weakened. Here, the film became a two-phase alloy containing g and d phases. The formation of these phases is consistent with the equilibrium phase diagram shown in Figure 4.2(a). Upon further heating to 390 8C, near the tail end of the exothermic peak, both g and d phases appeared. It can be shown that the meta-stable M1 phase first transformed into the g phase and then temporarily formed the uniform film of the g phase. When heating reached the crystallization temperature of the d phase, the d phase nucleated. From the results above, we can offer the following conclusion. Both 41.0 and 50.3 at.% Sn films initially contained the M1 phase uniformly, but after heat treatment their phase structure changed to their equilibrium form (Watanabe et al., 1999a,b).

Figure 4.7. A structural change by heat treatment in a meta-stable phase (M1) having the composition of Ni–50.3 at% Sn as observed in X-ray diffraction patterns.


181

Near the 85 at.% Sn composition in Ni –Sn alloy films, we found an M2 phase as indicated by the X-ray diffraction pattern in Figure 4.8. This phase was previously unreported and thus its structure was not known. Using an X-ray diffraction technique, we followed time-dependent structural changes occurring in the 87.1 at.% Sn film during annealing at 100 8C (see Figure 4.9). The intensity of X-ray diffraction peaks from the M2 phase gradually decreased, accompanied by the appearance of lines from Sn crystals and d phase (Ni3Sn4). An analysis of intensity variations in the X-ray diffraction lines revealed that the M2 phase has a structure (superlattice) that can be described as the lattice of Sn supersaturated with Ni. While heating, the M2 phase rejected Ni atoms and became Sn grains. The rejected Ni atoms were then absorbed into the d phase and finally the d phase grew at the expense of the Ni atoms. Since this phase transformation does not involve nucleation processes, the M2 phase is considered a non-equilibrium phase. From the above examples, we learned that plated alloy films could take various complex structures in the form of meta-stable and non-equilibrium phases. After heat treatment, however, these phases will eventually go back to their stable crystal form as expected from the equilibrium phase diagram.

Figure 4.8. X-ray diffraction patterns from Sn-rich Ni–Sn alloy films, showing the appearance of a meta-stable M2 phase.

182

Nano-Plating

Figure 4.9. Effect of annealing on the structure of the meta-stable M2 phase. These specimens were obtained by annealing Ni–87.1 at.% Sn alloy films for various time.

4.3. TRANSFORMATION OF AMORPHOUS PHASE TO EQUILIBRIUM PHASE

Amorphous solids are meta-stable phases in a thermodynamic sense. As described in Chapter 2, amorphous films can be obtained in two ways; these are the electrolytic and electroless deposition techniques (see Table 2.1). Alloy systems containing intermetallic compounds are capable of forming amorphous phases. Additionally, the composition of amorphous phases has to be close to the composition of the intermetallic compounds (Watanabe, 1989, 1993). The thermal stability of amorphous phases is high if the melting point of the corresponding intermetallic compound is high. If the alloy composition is lowered from the corresponding intermetallic compound composition, it forms a mixture of amorphous and fine-grained phases. Further decreases in the composition produces alloy films with a single phase of fine grains. Section 2.1.2.5 explains the composition range bounding the region between an amorphous and a mixed amorphous/fine-grained phase. Next, we will describe electrodeposited Ni– B alloy films covering various composition ranges and discuss how the films reach their equilibrium state through annealing (Onoda et al., 1990, 1992). Figure 4.10(a)– (e) are X-ray diffraction patterns showing structural changes in electrodeposited Ni– B alloy films after annealing. Boron concentration in the films is varied in the horizontal axis while the annealing temperature is given in the vertical axis.


Figure 4.10. Effect of annealing on the structure of electrolytic Ni –B alloy films having various boron concentrations.

183

184

Nano-Plating

For Ni– B alloy films containing boron at less than 11.2 at.% (Figure 4.10(a) and (b)), a {111} diffraction peak of Ni crystal appears and becomes sharper with increasing annealing temperature. This indicates the growth of Ni grains. The as-deposited 19.3 at.% B Ni – B film (Figure 4.10(c)) displayed a broad diffraction line, which could be taken as a sign of an amorphous phase. A rise in the annealing temperature to 280 8C increased the intensity of the {111} Ni peak. Finally, at an annealing temperature of 300 8C, Ni3B crystals were formed. For the Ni –B alloy films (Figure 4.10(d) and (e)), whose boron concentration is near the stoichiometric Ni3B intermetallic (25 at.% B), their diffraction pattern remained broad with annealing at temperatures below 280 8C, but X-ray lines of Ni crystals were absent. Ni3B crystals, however, suddenly appeared at a temperature of 280 8C. The 19.3 at.% B film contained fine Ni grains, although the X-ray diffraction pattern was broad and appeared amorphous. The fine-grained Ni grew during annealing and Ni3B crystals nucleated when the annealing temperature reached the crystallization temperature of Ni3B. The as-deposited 24.4 and 26.3 at.% B films, were initially amorphous in the absence of fine-grained Ni, but Ni3B crystals nucleated in the amorphous matrix when the annealing temperature reached the crystallization temperature of Ni3B. The annealing and crystallization processes of Ni –B alloy films obtained by electroless plating were investigated using a transmission electron microscope (TEM) (Watanabe and Tanabe, 1975; Suda et al., 1981) (see Figures 4.11 and 4.12). Figure 4.11(a) is a TEM micrograph showing the amorphous structure of as-deposited 25 at.% B Ni –B alloy films. After annealing, Ni3B crystals nucleated inside the amorphous matrix (Figure 4.11(b)). Figure 4.10(d) and (e) are diffraction patterns from electrolytic Ni– B alloy films, which correspond to the electroless Ni– B films in Figure 4.11(a) and (b). Figure 4.13(a) is a bright-field (BF) micrograph showing spherulite-like Ni3B crystals. The corresponding dark-field (DF) image is shown in Figure 4.13(b). The DF image of Ni crystals taken from Ni diffraction spots shows the uniform distribution of fine Ni crystals (Figure 4.13(c)). It is interesting to note that the fine Ni crystals are distributed uniformly even inside the Ni3B crystals. During low-temperature annealing at 280 8C and at 300 8C, Ni3B crystals must have nucleated and grown radially while passing around the existing fine Ni crystals. Therefore, for this composition (25 at.% B), it is reasonable to assume that these fine Ni crystals were already present in the as-deposited film. In other words, the asdeposited film was crystalline. A similar TEM study on the annealing behavior of electroless Ni –B films was reported by Masui et al. (1985). Figure 4.14 displays the TEM observations of crystallization processes in electroless Ni –B amorphous films. First, the film forms Ni3B nuclei during annealing and eventually becomes covered by this intermetallic compound as seen in Figure 4.14(b). If the boron concentration in the film is less than that (25 at.% boron) of the corresponding Ni3B


185

Figure 4.11. Structure of an electroless Ni –B film before and after heat treatment. (a) Before heat treatment and (b) an initiation of crystallization after heat treatment.

compound, excess dissolved Ni atoms will nucleate as fine Ni crystals inside the Ni3B grains or at their grain boundaries (Figure 4.14(d) and (d0 )). The microstructural changes described above are schematically summarized in Figure 4.15 (Ito et al., 2001). As-deposited alloy films, having a composition similar to the intermetallic compound, are amorphous (Figure 4.15(1)). In addition, the meta-stability of the amorphous structure is maintained up to the crystallization temperature of the intermetallic compound (Figure 4.15(2)). When the annealing temperature reached the crystallization temperature of Ni3B, the Ni3B nucleates (Figure 4.15(3)). With further increase in the annealing temperature, the intermetallic grains swept throughout the film, accompanied by the disappearance of the amorphous phase (Figure 4.15(4)). The structural evolution illustrated in Figure 4.15(b) is based on detailed experimental results in electrolytic Ni– P films (Ito et al., 2001). At the crystallization temperature of the intermetallic compound, crystallization occurred, as expected, and at the same time fine Ni crystals appeared. This phenomenon appears to occur when the film composition is slightly lower than that of the intermetallic compound. In this partially Ni-rich film, Ni crystals nucleated at the Ni3B grain boundaries amidst increasing temperature while Ni3B continued to crystallize.

186

Nano-Plating

Figure 4.12. Crystallization processes in an amorphous Ni–B alloy film during annealing. (a) BF image. (b) An electron diffraction pattern from an amorphous region. (c) An electron diffraction pattern from the crystallized region.

From X-ray diffraction patterns, it is difficult to ascertain whether or not an amorphous phase underwent any structural changes until the crystallization temperature of the associated intermetallic compound is reached. For amorphous Ni –B or Ni – P alloy films whose boron or phosphorus composition is slightly less than that of the intermetallic Ni3B or Ni3P phase, we found (Ito et al., 2001) lattice images indicative of several nm crystallites in the amorphous phase (see Figure 4.16(c)). Figure 4.16 consists of high-resolution TEM micrographs, which show how the structure of electrolytic Ni –P alloy films changes with P content. The low (12.4 at.% P) P film is apparently crystalline since the lattice image covers the whole film (Figure 4.16(a)). The18.5 at.% P Ni –P film contains fine crystallites with a size less than 10 nm distributed

Microstructural Changes in Plated Films During Heat Treatment Figure 4.13. Structural changes in Ni –B alloy films, which contained fine Ni crystallites although the electron diffraction pattern of the as-deposited films showed only a broad ring. (a) The BF image showing ellipse-shaped Ni3B crystals formed as a result of heat treatment. (b) The DF image taken from one of the Ni3B diffraction spots. (c) The DF image obtained from the Ni diffraction spot.

187

188

Nano-Plating

Figure 4.14. Crystallization processes of electroless Ni–B amorphous films [(a)–(c)]. (d) An enlarged image from the part of (c) and (d0 ) its DF image of the Ni crystals.


189

Figure 4.15. A schematic diagram illustrating structural changes occurring in electrolytic Ni–B alloy films by heat treatment from unstable to stable forms. The as-deposit structure of the Ni–B alloy films was an amorphous or a crystalline phase.

190

Nano-Plating

Figure 4.16. High-resolution TEM lattice images of electrolytic Ni–P alloy films containing various P concentrations.


191

throughout the film and thus this film was also crystalline (Figure 4.16(b)). In the case of 24.3 at.% P film, a lattice image was not present, indicating that the film was amorphous (Figure 4.16(c)). The 21.5 at.% P film contained several nm crystallites highlighted by lattice images. These crystallites, however, did not grow after annealing. It is possible that the crystallites dissolved into the matrix after annealing because their size was below the so-called critical nuclei size. In general, the energy of a small crystallite consists of volume and surface energies. If the volume energy becomes larger than the surface energy, the crystallite will grow by heating. Conversely, if the surface energy is larger than the volume energy, the crystallite will disappear by dissolution into the matrix. The critical size can be determined by equating the surface energy to the volume energy (Japan Inst. Metals, 1982). A crystallite, with a size smaller than the critical size, is called an embryo (Japan Inst. Metals, 1982). If these embryos are annealed at a temperature below the crystallization temperature of the associated intermetallic compound, they will dissolve into the matrix. Consequently, the structure of the amorphous film becomes more uniform. This embryo dissolution process is probably connected to a phenomenon known as the structural relaxation of an amorphous film. The presence of embryos was also observed (Wang et al., 2000; Ito et al., 2001a,b) in other electrolytic amorphous films. From this observation, we can conclude that there is atom segregation in plated films and the extent of the segregation corresponds to the embryo size. The embryo size is not constant but varies depending on the type of alloys (for example, Ni– P, Ni –B, Ni– W, Fe –W, etc.). The structure shown in Figure 4.15(c), which corresponds to both Figures 4.10(c) and 4.13, is not uniformly amorphous but contains fine Ni crystallites (Figure 4.15(1)). These Ni crystallites grew by low-temperature annealing (Figure 4.15(2)). At the crystallization temperature of the intermetallic compound (Ni3B), Ni3B crystals suddenly nucleated (Figure 4.15(3)). The Ni3B crystal then grew radially, while its crystal front migrated outward and passed around pre-existing Ni crystallites. In Figure 4.15(d), the film is made by an assembly of fine Ni crystallites (Figure 4.15(1)). These Ni crystallites grew continuously with increasing temperature (Figure 4.15(2)) and excess boron atoms have segregated at the Ni grain boundaries. A further increase in temperature initiated the formation of Ni3B (Figure 4.15(3)), which finally led to the stable two-phase alloy structure of Ni and Ni3B (Figure 4.15(4)). We have described how the microstructure of plated alloy films having various compositions and structures, changes from an unstable form to a stable form with increasing temperature. Crystalline films can be transformed to an amorphous phase by continuously reducing the grain size, which is in turn varied by changing the composition (Wang et al., 2000; Ito et al., 2001a,b). Here, it is reasonable to assume that general grain boundaries are

192

Nano-Plating

made up of a kind of amorphous phase. Then, an increase in the grain boundary area (or a decrease in the grain size) corresponds to an increase in the amorphous phase, i.e. the volume of an amorphous phase increases with decreasing grain size. Accordingly, at the composition immediately before the whole film becomes completely amorphous, the structure is an assembly of fine crystallites and embryos distributed uniformly in an amorphous matrix. When these crystallites and embryos are gone, the whole film becomes amorphous. Electrodeposited amorphous Ni – B films readily change to their stable phase by heat treatment. Some films go through two or three stages of meta-stable intermediate phase formation before reaching the final phase (Masumoto and Maddin, 1975; Masumoto et al., 1977; Masumoto, 1982). Thermal stability of an amorphous phase and its crystallization temperature depend on the heating rate and holding temperature. For this reason, the most reliable method for determining these properties is to construct a time –temperature – transformation (TTT) curve (Watanabe and Scott, 1980; Oriya et al., 1987).

REFERENCES Augis, J.A. & Bennett, J.B. (1978) J. Electrochem. Soc., 125, 330. Bennett, J.B. & Tompkins, H.G. (1976) J. Electrochem. Soc., 123, 999. Clarke, M. & Dutta, P.K. (1971) J. Phys. D: Appl. Phys., 4, 1652. Dill, J., Charlier, J. & Winand, R. (1998) J. Mater. Sci., 33, 2771. Dutta, P.K. & Clarke, M. (1968) Trans. Inst. Metal Finish., 46, 20. Enomoto, H., Fujiwara, Y., Isaki, M. & Ono, H. (1982) J. Metal Finish. Soc. Jpn, 33, 369. Ito, K., Wang, F. & Watanabe, T. (2001a) J. Jpn Inst. Metals, 65, 1025. Ito, K., Wang, F. & Watanabe, T. (2001b) J. Jpn Inst. Metals, 65, 495. Japan Inst. Metals,, Editor (1982) Metal Handbook, Maruzen Pub. Co., p. 332, 374, 1445. Lo, C.C. (1980) J. Appl. Phys., 51, 2007. Masui, K., Masuda, M., Maruno, S. & Kawaguchi, T. (1985) J. Metal Finish. Soc. Jpn, 36, 50. Masumoto, T., Editor (1982) Fundamentals of Amorphous Metals, Ohmu Pub. Co., p. 91. Masumoto, T. & Maddin, R. (1975) Mater. Sci. Engng, 19, 1. Masumoto, T., Inoue, A. & Kimura, H. (1977) J. Jpn Inst. Metals, 41, 730. Onoda, M., Tsuchiya, T., Ogawa, K. & Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 388. Onoda, M., Tsuchiya, T., Shimizu, K. & Watanabe, T. (1992) J. Surf. Finish. Soc. Jpn, 43, 138. Oriya, H., Hasegawa, N., Misaki, Y. & Tanabe, Y. (1987) J. Metal Finish. Soc. Jpn, 32, 637. Schubert, R. (1973) J. Electrochem. Soc., 125, 1215. Suda, H., Watanabe, T., Misaki, Y. & Tanabe, Y. (1981) J. Jpn Inst. Metals, 45, 117. Tamura, T. & Hosokawa, K. (1977) J. Metal Finish. Soc. Jpn, 28, 564. Tanabe, Y. & Shimizu, Y. (1975) J. Metal Finish. Soc., 26, 406. Wang, F., Ito, K. & Watanabe, T. (2000) J. Jpn Inst. Metals, 64, 1133. Watanabe, T. (1989) J. Surf. Finish. Soc. Jpn, 40, 375. Watanabe, T. (1993) Current topics in amorphous metals, Physics & Technology, Elsevier Science Publishers B.V., p. 137. Watanabe, T. & Scott, M. (1980) J. Mater. Sci., 15, 1131.

Microstructural Changes in Plated Films During Heat Treatment Watanabe, T. & Tanabe, Y. (1975) J. Jpn Inst. Metals, 39, 831. Watanabe, T., Arai, K., Hirose, T. & Chikazawa, M. (1999a) J. Jpn Inst. Metals, 63, 496. Watanabe, T., Hirose, T., Arai, K. & Chikazawa, M. (1999b) J. Jpn Inst. Metals, 63, 489. Wilson, G.C. (1972) Trans. Inst. Metal Finish., 50, 109. Wynne, B.E., Edington, J.W. & Rothwell, G.P. (1971) Met. Trans., 3, 301.

193


Chapter 5

Control of Macrostructure in Plated Films and Fabrication of Three-Dimensional Microstructure 5.1. 5.2.

Introduction Macrostructure control in plated films 5.2.1 Columnar structure 5.2.2 Fine-crystal structure (nanocrystals) 5.2.3 Amorphous structure 5.2.4 Single crystals 5.2.5 Multi-layer (ML) film 5.2.6 Multi-layer (ML) containing alternating layers of crystalline and amorphous phases 5.2.7 Epitaxial multi-layer 5.2.8 Amorphous/crystalline graded structure 5.2.9 Composite coating 5.2.10 Graded composite coating 5.2.11 Other structures 5.3. Fabrication of three-dimensional microstructural body References

197 197 197 197 199 199 199 200 200 202 202 202 202 204 204


Chapter 5

Control of Macrostructure in Plated Films and Fabrication of Three-Dimensional Microstructure 5.1. INTRODUCTION

Film macrostructure described in this chapter can be defined as a macroscopic representation of the overall film structure. This description does not involve microscopic characteristics, such as the size or the orientation of grains. Using a structure control method for plated films (see Figures 1.1 and 1.3), we can fabricate any type of macrostructure, which in turn allows us to tailor new functional materials in plated films. The types of macrostructures are listed in Figure 5.1. In this chapter, we will describe how to fabricate these macrostructures, as well as three-dimensional microstructures (Watanabe and Fujita, 1995), in plated films.

5.2. MACROSTRUCTURE CONTROL IN PLATED FILMS

5.2.1 Columnar structure A columnar structure is a frequently observed growth form in plated films. Pure metal deposits often take this growth form, which initiates as fine grains followed by the nucleation of columnar grains. The columnar grains tend to grow along the fast growth direction of the material. When these oriented columns fill the entire film surface, it is said that the preferred orientation (or film texture) is developed. 5.2.2 Fine-crystal structure (nanocrystals) This structure consists of fine crystals and often shows up in high melting point (MP) metals or in alloy deposits. Low MP metals generally grow large grains, which can be refined by alloying. Even for high MP metals, alloying will make the grain size much smaller. Alloying with a small amount of non-metallic light elements, such as phosphorus, boron, carbon, sulfur, nitrogen, arsenic, etc. is highly effective in refining the grains. When the concentration of alloying element exceeds a certain value, the film may become amorphous. Alloying with metallic elements can also help in refining the grains. The degree of the refining effect depends on the type of metallic element. When the grain size becomes extremely fine and reaches the nanometer scale through alloying, we can call the grains nanocrystals (Czira´ki et al., 1994; Bakonyi et al., 1996a; Lu et al., 2000). Grain refining for pure metal films, however, is very difficult, especially when the purity 197

198 Nano-Plating

Figure 5.1. Various macrostructures observed in plated films.

Control of Macrostructure in Plated Films

199

is high. This is because the recrystallization temperature of pure metals decreases with increasing purity. Additionally, pure metals can sometimes recrystallize at ambient temperatures. It is of interest to mention that there have been some efforts to produce nanoscale electrodeposits in a powder form (Delplancke et al., 1995, 1997). 5.2.3 Amorphous structure Plated amorphous films were described in Chapter 2. As shown in Table 2.1, an amorphous phase can be obtained in alloy deposition systems of metal/metal, metal/non-metal, and metal/hydrogen. Semiconductor alloy films often become amorphous (Omi, 1989). To understand the formation of amorphous films, refer to Chapter 2. 5.2.4 Single crystals Single-crystal films can be prepared by electroplating methods, which employ film growth on single-crystal substrates by epitaxy. Metals are generally used as substrate materials, onto which pure metal or solid solution alloy films can be plated. The initial stages of electrodeposition proceed by epitaxial growth (lattice-matched growth). There is a critical thickness, at which epitaxy can break down. The critical thickness depends on the degree of a lattice mismatch existing between the film and substrate as well as plating conditions. For a more detailed discussion on epitaxy, refer to Chapter 1. 5.2.5 Multi-layer (ML) film (Watanabe, 1994a,b) The fabrication of ML films by electrodeposition can be achieved by alternately plating two or more different metals in a layer form. An artificial superlattice can be constructed if a ML film is formed with a mono-atomic layer. In practice, however, the production of such a superlattice is very difficult. The reason for this difficulty stems from the initial nucleation stages of electrodeposition. The initial stages often proceed via a Volmer – Weber type growth mechanism, whereby film growth is initiated through island formation. In this growth mode, the minimum thickness at which the whole substrate surface is covered by a plated film corresponds to the thickness at which island coalescence is completed. Thus, it is easy to visualize that the minimum layer thickness obtainable by an electrodeposition method is equivalent to the height of coalesced islands. As long as the film grows in this manner, the production of any artificial superlattices, involving monoatomic layers, is impossible. There are several reports (Bakonyi et al., 1996b; Jyoko et al., 1996a,b; Nabiyoun and Schwarzacher, 1996; Czira´ki et al., 1997) claiming the successful fabrication of ML films, which consist of a few nanometer-thick alternating layers. These films exhibited the giant magneto-resistance (GMR) effect. The reason why the GMR effect was observed in these ML films is not understood at the moment.

200

Nano-Plating

Two-bath and one-bath methods are available for producing ML films by electrodeposition. There is also an additional method that can produce ML films by alternately running two kinds of plating solutions on a cathodically biased substrate surface. In the two-bath method, the substrate is plated alternately in two kinds of plating baths, while inserting water-rinsing/drying steps between the two platings. An example is shown in Figure 5.2(a), which displays the cross-section of a ML of Cu and Fe– 80 at.% Ni (Permalloy). This film exhibited the GMR effect, although not very strong. In the one-bath method, two kinds of metal salts, having different reduction potentials, are dissolved into one plating bath. ML films are obtained by increasing/decreasing the overpotential (current density) periodically. Figure 5.2(b) shows a ML of Cu and Ni made by the one-bath method. In this method, it is impossible to make two layers consisting of two pure metals. Strictly speaking, the Cu – Ni ML is made of alternating layers of pure Cu and Ni-rich Cu – Ni alloy. 5.2.6 Multi-layer (ML) containing alternating layers of crystalline and amorphous phases Alternating MLs of crystalline and amorphous phases are easily constructed in any plating bath that produces an amorphous deposit. As shown in Figures 2.18 and 2.19, the phosphorus, boron, or metal content of the film can be varied by adjusting the overpotential (current density), i.e. a periodic change in the overpotential (current density) produces alternating layers of amorphous/crystalline phases. Figure 5.2(c) is the cross-sectional view of an amorphous/crystalline ML obtained from a Ni –B alloy-plating bath (Onoda et al., 1993, 1994). Each layer is 12.5 nm thick. Electron diffraction patterns from the amorphous and crystalline layers are also shown in Figure 5.2(a) and (b). The structure of each layer changes with the boron content, which in turn is controlled by the current density. Therefore, it is possible to change the grain size and ultimately the physical properties of ML films. 5.2.7 Epitaxial multi-layer Two dissimilar metals with the same crystal structure and a small lattice mismatch can be made to form a multi-layer film by epitaxy. Figure 5.2(b) is an example of epitaxially grown Cu/Ni multi-layer film. Both Cu and Ni are face-centered cubic (fcc) and their lattice mismatch is only 2.3%. As indicated by the uniformly distributed black –white bend contour, the whole Cu/Ni composite film appears uniform without any structural discontinuities and behaves like a single crystal. If two metals with different crystal structures are chosen (for example fcc and bcc metals), the resulting multi-layer films will have a structure similar to that of a plywood sheet.

Control of Macrostructure in Plated Films Figure 5.2. Multi-layered (ML) films fabricated by plating methods. (a) An Fe– Ni ML film grown on copper by the two-bath method. (b) A Cu –Ni ML film obtained by the one-bath method, and (c) a crystalline Ni–Co –B/amorphous Ni –Co –B ML film by the one-bath method.

201

202

Nano-Plating

The selection of two structurally different metals will therefore make it possible to produce high-strength materials, which could endure fracture by plastic deformation. 5.2.8 Amorphous/crystalline graded structure Plating baths capable of forming amorphous/crystalline multi-layer films can be used to produce films with structurally graded crystalline-to-amorphous or amorphous-to-crystalline layers. In this case, the current density is varied continuously. 5.2.9 Composite coating Composite films containing various types of powder are manufactured commercially and can be obtained by electrodeposition. The type of powder is not limited to metals, indeed a variety of materials are used including ceramic, polymer, and diamond powders. One application is found in plated chromium films. Plated chromium films often contain fine surface cracks due to their intrinsically high internal stresses. It has been shown that high wear-resistance films can be obtained by filling such cracks with ceramic powder during plating. Most recently, Ni films containing fine TiO2 or Ag particles were manufactured to produce virus-resistant films (Ishiguro et al., 1999; Zhao et al., 1999). Composite films of nickel and Teflon powder were also produced to obtain a water-repellant metal surface. 5.2.10 Graded composite coating A graded composite structure can be obtained by changing the concentration of composite powder continuously with time. One successful example is Ni film plated with zirconia powder. Here, a graded structure was obtained by deliberately varying the quantity of zirconia powder during plating in such a way that the internal region is more metallic, while the surface region is more ceramic. 5.2.11 Other structures In Figure 5.1(a) – ( j), we describe plating methods for producing 10 types of macrostructures. It is possible that there are other methods for making these macrostructures. Nevertheless, each macrostructure exhibits characteristic physical properties. As discussed in Chapter 4, it is important to remember that plated films, especially those containing amorphous or meta-stable phases, undergo structural changes with heat treatment. The amorphous phase will always transform to a hard intermetallic compound. The formation of such a hard compound is beneficial for obtaining high-strength materials. For example, heat treatment of amorphous Fe– W alloy films plated on the cutting edge of a drill successfully produced a highly wear-resistant ball mill by forming the intermetallic compound (Watanabe et al., 1997). Another example is the fabrication of a roller, on which an amorphous Cr –C alloy film was plated and then heat-treated to form the intermetallic


Figure 5.3. Fabrication of a three-dimensional microstructural body by plating methods. (a) A replica method, where the plated film is used as a replica for the underlying substrate surface structure. (b) A method using a masking or a photolithography technique. (c) A localized plating method using a laser/light illumination or a needle-shaped cathode.

203

204

Nano-Plating

compound (Morikawa et al., 1991). It is interesting to investigate the relationship between the structural change by heating and their properties.

5.3. FABRICATION OF THREE-DIMENSIONAL MICROSTRUCTURAL BODY (WATANABE ET AL., 1995)

A three-dimensional microstructure is defined as a microscopic solid body. There are many methods for fabricating three-dimensional microstructures. In Figure 5.3, we illustrate three methods. Figure 5.3(a) is a replica method, where a substrate surface containing irregularities is plated with a film. The plated film will replicate the underlying substrate surface irregularities. The original substrate can be metal, plastic, or ceramic material, but the final replicating material is a metal. Plating can be repeated on the same substrate, and if the plated film adheres properly to the substrate, it will not peel off easily. The CD manufacturing method uses the replicated metal film as a stamper; in this case, the film surface is oxidized deliberately to allow easier peeling. Figure 5.3(b) is a method that uses a photolithographic technique. Most recently, progress has been made to apply this advanced technique to fabricate electronic devices and micro-machines in the electronics industry. Multi-level photolithography is used to fabricate copper interconnects in ULSI circuits by the damascene technique and to manufacture an actuator in a static motor by inserting a sacrificial layer. The process for producing micro-machines by plating methods is called Lithographie Galvanoformung Abformung (LIGA). Figure 5.3(c) illustrates a method for preferentially depositing a metal on the laserilluminated area of a substrate only, which is submerged in a plating solution. If a semiconductor material is used as a substrate, even visible light will induce the plating. The plating, therefore, has to be performed in a dark room.

REFERENCES ´ ., Gero¨cs, I., Varaga-Josepovits, K., Arnild, B. & Bakonyi, I., To´th-Ka´da´r, E., Poga´ny, L., Czira´ki, A Wetzing, K. (1996a) Surf. Coating Technol., 78, 124. ´ ., Gero¨cs, I., Nabiyouni, G. & Bakonyi, I., To´th-Ka´da´r, E., Becsei, T., To´th, J., Tarno´czi, T., Czira´ki, A Schwarzacher, W. (1996b) J. Magn. Magn. Mater., 156, 347. ´ ., Fogarassy, B., Gero¨cs, I., To´th-Ka´da´r, E. & Bakonyi, I. (1994) J. Mater. Sci., 29, 4771. Czira´ki, A ´ ., Gero¨cs, I., Fogarassy, B., Arnild, B., Reibold, M., Wetzig, K., To´th-Ka´da´r, E. & Czira´ki, A Bakonyi, I. (1997) Z. Metallkd., 88, 781. Delplancke, J.L., Dibella, U., Reisse, J. & Winand, R. (1995) MRS Symp. Proc., 372, 75. Delplancke, J.L., Bouesnard, O., Reisse, J. & Winand, R. (1997) MRS Symp. Proc., 451, 383. Ishiguro, F., Nishikawa, T., Amano, R. & Koyanagi, T. (1999) Materia, 38, 64.


205

Jyoko, Y., Kashiwabara, S. & Hayashi, Y. (1996a) J. Surf. Finish. Soc. Jpn, 47, 1025. Jyoko, Y., Kashiwabara, S. & Hayashi, Y. (1996b) J. Magn. Magn. Mater., 156, 35. Lu, L., Sul, L. & Lu, K. (2000) Science, 287, 1463. Morikawa, T., Yokoi, M., Eguchi, S. & Fukumoto, Y. (1991) J. Surf. Finish. Soc. Jpn, 42, 100. Nabiyoun, G. & Schwarzacher, W. (1996) J. Magn. Magn. Mater., 156, 355. Omi, T. (1989) J. Surf. Finish. Soc. Jpn, 40, 368. Onoda, M., Shimizu, K., Tsuchiya, T. & Watanabe, T. (1993) J. Magn. Magn. Mater., 126, 595. Onoda, M., Shimizu, K., Tsuchiya, T. & Watanabe, T. (1994) J. Surf. Finish. Soc. Jpn, 45, 714. Watanabe, T. (1994a) J. Surf. Sci. Soc. Jpn, 1, 637. Watanabe, T. (1994b) J. Surf. Finish. Soc. Jpn, 45, 1244. Watanabe, T. & Fujita, S. (1995) J. Precision Engng Soc. Jpn, 61, 1381. Watanabe, T. & Watanabe, S. (1997) J. Surf. Finish. Soc. Jpn, 48, 549. Zhao, Z.-H., Kusakari, S., Sakagami, Y. & Osaka, T. (1999) J. Surf. Finish. Soc. Jpn, 50, 25.


Chapter 6

Characterization Methods for Plated Films 6.1.

Structure of metal substrates 6.1.1 Deformation and annealing texture 6.1.2 Surface deformation layer 6.1.3 Single-crystal substrate 6.1.4 Pre-treatment methods for polycrystalline substrates 6.1.4.1 Annealing treatment 6.1.4.2 Electropolishing 6.2. Structural determination of plated films 6.2.1 XRD method 6.2.2 Structural analysis by TEM 6.2.2.1 Observation of the initial stages of electrodeposition 6.2.2.2 Plan-view observation of thick plated films 6.2.2.3 Crystallographic matching relationship between plated film and substrate 6.2.2.4 Preparation of TEM cross-section samples 6.3. Grain-size measurements 6.3.1 Direct measurements by TEM observations 6.3.2 X-Ray diffraction method 6.4. Observations of surface morphology 6.4.1 Scanning electron microscope (SEM) 6.4.2 Replica method for TEM observations 6.4.3 Observations by AFM 6.4.4 Measurements by a surface profilometer 6.5. Measurement of preferred orientation 6.5.1 Definition of preferred orientation and its description 6.5.2 Measurement of texture in plated films References

209 209 209 212 212 212 213 215 215 216 218 221 223 223 235 235 238 240 240 246 248 248 248 250 250 252


Chapter 6

Characterization Methods for Plated Films The structure of plated films can be classified into seven types as shown in Figure 1.1. In this chapter, we will describe how to characterize these structures.

6.1. STRUCTURE OF METAL SUBSTRATES

6.1.1 Deformation and annealing texture The use of a well-defined substrate structure is essential for interpreting the structure and property of plated films. Thus, it is important to give an appropriate pre-treatment to the surface and structure of the substrates. Although ceramic and plastic materials are also used as substrates, we will describe only the pre-treatment of metal substrates. It is important to remember that ceramic and plastic substrates are exclusively used for depositing electroless films. For the pre-treatment of such substrates, refer to the description of electroless plating in Section 2.2.1. Rolled or forged metal sheets are generally used as substrates, although cast materials are sometimes employed. Commercially rolled sheets contain grains that are stretched along the rolling direction and therefore are not uniform in a crystallographic sense (see Figure 6.1). This structural anisotropy also appears in the mechanical properties, such as hardness and bending/tensile strengths. The structure of plated films grown on such substrates can be affected by the anisotropy of the substrate structure. The development of such a structural anisotropy in rolled metal sheets can be attributed to the mode of a plastic deformation by the rolling process, which proceeds by the slipping of dislocations on specific crystallographic planes along specific directions. This process will rearrange the grains according to the deformation mode induced by rolling and cause a so-called deformation texture. Plastically deformed metal sheets can undergo recrystallization and grain growth processes after annealing. They also develop their own unique texture called annealing texture. 6.1.2 Surface deformation layer In the previous section we described how rolled sheets are subjected to a severe plastic deformation through thickness reduction between two rollers. In particular, the surface region of the rolled sheets receives the highest plastic deformation via a frictional force from the rollers and sometimes becomes amorphous (see Figures 6.1 and 6.2). Such an 209

210

Nano-Plating

Figure 6.1. The grain structure and surface deformation layer of a rolled metal sheet.

amorphous surface layer is called a Laves layer. Commercially rolled sheets always contain a severely deformed surface layer, known as a surface deformation zone (SDZ). The SDZ layer also forms during other mechanical processes, such as metal cutting or grinding. The thickness of the SDZ depends on the type/structure of metals/alloys and the machining method. The SDZ can grow as thick as several hundred mm. Figure 6.3 shows surface morphology changes on a rolled pure copper sheet during chemical etching. Here, the SDZ layer dissolves after prolonged etching.

Figure 6.2. A surface deformation layer in rolled materials.

Characterization Methods for Plated Films Figure 6.3. A surface morphology change with etching time in a polycrystalline copper sheet. A surface deformation layer is first dissolved by the etching, followed by the appearance of the crystallographic facets.

211

212

Nano-Plating

6.1.3 Single-crystal substrate Plated films generally grow epitaxially on single-crystal substrates and form a low-energy interface with the substrates. To investigate the structural evolution of plated films, therefore, it is convenient to use single-crystal substrates. Polycrystalline substrates are not ideal for such a study because their deformation or annealing texture may influence the structure of plated films up to a certain film thickness. A single crystal of pure copper or aluminum can be easily grown using a uni-directional solidification method known as the Bridgeman technique. Details of the equipment and the growth method can be found in the following review article (Watanabe and Nakata, 1994). In short, single crystals of 20-mm-diameter copper and 5-mm-diameter aluminum were successfully grown by this technique. If larger crucibles and electric furnaces are used, then larger single crystals can be obtained. Firstly, copper charge is put into a graphite crucible with a pointed bottom. Then the crucible is inserted into a quartz tube and the tube evacuated using a rotary pump. A single crystal is obtained by a uni-directional solidification method, which requires the raising/lowering of an electric furnace around the tube three times. The length of the crystal is controlled by the length of the crucible; the longer the crucible, the longer the crystal. After determining the orientation using a Laue X-ray diffraction (XRD) method, a wafer of the single crystal is sliced. After the slicing, the surface is finally electropolished to obtain an SDZ-free wafer with a mirror-finish surface. 6.1.4 Pre-treatment methods for polycrystalline substrates 6.1.4.1. Annealing treatment. Commercial metal sheets are usually made by rolling and are usually fine-grained. The internal stress is high due to the work-hardening induced by rolling. In addition, they are highly defective and contain a high density of lattice defects, such as dislocations, stacking faults, vacancies, and interstitials. The degree of work-hardening depends on the amount of rolling treatment and the type of material. As previously discussed, these metal sheets are expected to contain the SDZ layer. To obtain plated films that are not affected by the structure of the substrate, it is important to homogenize the substrate structure by reducing its structural anisotropy and imperfections. The first step to this homogenization is to anneal the substrate. The annealing treatment is done by holding the substrate at temperatures above the recrystallization temperature and then cooling it slowly. It is important to note that the annealing treatment does not necessarily remove the texture or the SDZ layer completely. If the substrate is an alloy metal, it may possess characteristic structural features, such as foreign or mixed phases (Figure 1.3). To eliminate such phases, the annealing temperature, its duration, and cooling rate have to be carefully determined using the equilibrium phase diagram. In some cases, the presence of foreign phases cannot be avoided. (Figure 1.45).

Characterization Methods for Plated Films

213

6.1.4.2. Electropolishing. The structures of plated films are affected by surface irregularities, oxides, contaminants, and an SDZ layer present on substrates. To insure that the plating proceeds without the adverse effect of the substrate, we have to remove these surface features. The best method for this is to use an electropolishing technique. A chemical etching method is also available but leaves surface irregularities due to etching. Next, we will describe the principle of surface smoothing by an electroplating method. Figure 6.4(a) is a schematic diagram of a current – voltage curve for a metal anode. At low voltage, a steep slope appears and this region is called an active region, in which a metal dissolves. After passing through the peak, the curve becomes flat. This flat section is called an inactive region, where only the grain boundaries dissolve, thus leaving an irregular surface. With a further increase in voltage, another steep slope appears. In this region, oxygen evolution takes place. To obtain a bright mirror surface, we must set the voltage between the inactive and oxygen evolution regions as shown by a circle. There are two possible mechanisms that produce mirror surfaces free of irregularities by electropolishing. The first mechanism is based on current density, which is higher in the protruded portion than in the valley region. Figure 6.4(b) illustrates how the current is concentrated at the protruded portion of the surface during electropolishing. The protruded portion will be polished preferentially, thus making the surface smoother. The second mechanism involves the dissolved viscous products at the protrusions, which migrate down to the neighboring valley region during electropolishing. This viscous

Figure 6.4. A schematic diagram showing (a) an anodic polarization curve for a metal and (b) a mechanism for electropolishing a metal surface.

214

Nano-Plating

product tends to protect the valley region from dissolution, thus promoting the preferential dissolution of the protruded region. Nevertheless, to find a suitable voltage for obtaining a mirror-like surface, it is important to construct the current –voltage curve experimentally. From this curve, one can determine the polishing region. If the voltage moves into the oxygen-evolution potential, many sharp pits will appear on the surface, accompanied by the evolution of oxygen gas. This results in an irreproducible surface finish. The electropolishing solutions used are listed in Table 6.1. An electropolishing setup is shown in Figure 6.5. We found that a freshly made polishing solution does not function well. Instead, the used polishing solution performs better because the presence of dissolved metal ions in the solution is beneficial. Metal ions can be introduced into a fresh solution by electropolishing a dummy sample until the solution is saturated with the metal ion. For a cathode, we used a stainless steel sheet. It is recommended not to remove a metal film deposited on the cathode during polishing. Although the deposited metal film on the cathode leaves an irregular surface, it offers a beneficial effect on polishing by increasing the surface area. Furthermore, the solution cooling is desirable for obtaining good polishing. The solution can be cooled by circulating the solution through a coil chilled in an ice bath. One of the key steps in obtaining the mirror surface is to take a specimen out of the solution after a 10-s polishing and then to remove the surface residue by wiping with cotton wool pick. After repeating this process two or three times, we can acquire a mirror-like surface. It is important not to use water for this step. As described earlier, at least a 15-min polishing will be necessary to remove the deformed surface layer.

Table 6.1. Electropolishing solutions and their polishing conditions. Metal

Electropolishing solution

Polishing condition

Cu and Cu alloys

Phosphoric acid : water (2:1)

Fe, Ni, Al and their alloys. Stainless steel

60% Perchloric acid : 100% acetic acid (1 : 3 , 4) (Jacket solution)a 60% perchloric acid ¼ 105 ml; 100% acetic anhydride ¼ 386 ml; water ¼ 9 ml

1.5 V when the substrate polishing area is 1 , 5 cm2 15 A/dm; 15 8C The condition depends on the type of metals, but the lower temperature produces better results

a

Slowly add perchloric acid to acetic anhydride solution with a pipette. Since heat will be generated during the addition, a beaker containing the solution

should be chilled in an ice water bath. Make sure that the solution is agitated during the addition, so that the solution temperature does not rise more than 30 8C. The temperature initially rises but heating will not occur at the end. Finally add water. Be careful that the water addition will also generate heating.


215

Figure 6.5. An experimental setup for electropolishing substrate materials.

6.2. STRUCTURAL DETERMINATION OF PLATED FILMS

XRD and transmission-electron-microscope (TEM) methods can be used for a structural analysis of plated films. 6.2.1 XRD method There are a number of excellent books describing XRD methods for a structural analysis of general crystals (Cullity, 1978). In this chapter, we focus primarily on an analysis of plated films. The structure of plated films is the same as the structure of the bulk solids prepared by the standard metallurgical processes. In these processes, solids are prepared through first melting in a crucible, followed by slow cooling to solidification. Methodological aspects of

216

Nano-Plating

the X-ray technique for plated films are the same as the technique used for bulk solids. There are, however, several points of caution in analyzing plated films by an X-ray method. An XRD method assumes that the structure of plated films is uniform along the film thickness direction. Although plated films are generally thin and fine-grained, their structure still changes along the thickness direction (see Figures 1.51 –1.55). Some plated films have a layered structure that forms due to the effect of the substrate structure or due to a change in composition during plating. This type of layer structure can be easily recognized using a cross-section TEM, but cannot be determined by an X-ray method. In our experience, the structural analysis by an XRD method is not reproducible. This resulted in a scattering of the property data. For example, the structure shown in Figure 1.55 is determined to be crystalline when the film is thin, but amorphous when it is thick. Here, the thick film is incorrectly assumed to be completely amorphous by an X-ray analysis, although the film consists of two layers of crystalline and amorphous phases. Similarly, the film shown in Figure 1.54(A) was determined to be fine-grained for a thin deposit and large-grained for a thicker deposit. It is difficult to obtain an XRD pattern if the film thickness is less than 1 mm. The effective film thickness (the effective depth of X-ray penetration), X, for obtaining the maximum intensity can be given by the following equation: X ¼ KX sin u=2m

ð6:1Þ

where X is the effective film thickness (in cm) for obtaining 99% intensity, KX, the constant, and m, the absorption coefficient. For example, using an X-ray beam with a wavelength of 9 nm and an absorption coefficient of 200 cm21, we find the effective thickness, X, of a platinum film for the 99% intensity of the {111} peak appearing at 39.88 to be 34.9 mm. If the film is less than 34.9 mm thick, the diffraction intensity is not strong enough to be useful. Also, in the case of plated films, diffraction peaks appear, which are not listed in the JCPDS cards or have never been reported in the literature. As explained in Section 1.3, the structure of plated films is similar to the structure obtained by rapid quenching from high temperatures and thus can become amorphous or meta-stable phases, which are not listed in the equilibrium phase diagram. If such unknown peaks appear, it becomes difficult to conduct a structural analysis. More recently, however, there exist excellent software programs that can determine the lattice parameter of a crystal. 6.2.2 Structural analysis by TEM (Sotomura, 1989; Kanto Division of Japanese Society of Electron Microscopy, 1991; Japanese Society of Electron Microscopy Summer School, 1994,1995,1996a,b; Hong and Saka, 1997; Saka, 1997; Tanaka et al., 1997) A structural analysis of plated films by TEM can be performed using electron diffraction patterns. The area of interest, where an electron diffraction pattern is obtained, can be


217

chosen by using a selected area aperture, whose size varies from 80 nm to 300 mm. Clearly, the TEM can analyze a much smaller area than the X-ray method. A modern TEM equipped with a field emission gun can produce a focused beam less than 1 nm, which allows not only an electron diffraction analysis at a very small area, but also a qualitative/quantitative compositional analysis by X-ray and energy loss spectroscopy. Figures 1.28, 2.28(b) and (c) were obtained using these techniques. The purpose of using TEM for plated films is schematically illustrated in Figure 6.6. Here the TEM is used to observe four features in plated films: (1) the initial stages of

Figure 6.6. TEM specimen configurations suitable for studying various aspects of plated films. Specimens for studying: (1) the initial stages of electrodeposition, (2) the microstructure at different film depths, (3) crystallographic relationship between the plated film and the substrate, and (4) the overall cross-section.

218

Nano-Plating

electrodeposition; (2) structural changes occurring at intermediate stages; (3) crystallographic relationship at the film/substrate interface; and (4) cross-sectional structure. The TEM specimens have to be prepared according to each objective. An optimum TEM specimen thickness depends on the type of metals, but has to be below 100 nm, generally less than 50 nm. There are a number of TEM specimen preparation methods. For further information, read the author’s articles (Watanabe, 1982, 1989a,b, 1995, 1996) on TEM specimen preparation methods and observations for plated films. 6.2.2.1. Observation of the initial stages of electrodeposition. We describe a TEM observation method for the nucleation and growth stages of electrodeposition as seen in Figure 2.2. For this observation, we use the plated film itself as a self-supporting TEM sample after stripping off the substrate, which meant that the plated film had to be thin enough to be electron transparent. After a short period of electroplating, the deposit – substrate composite film is dipped into an etching solution in which the substrate preferentially dissolves. Prior to dipping, an approximately 2-mm square mesh is scarred on the deposit – substrate sheet with a knife, so that a 2-mm square film will float off the substrate (see Figure 6.7). After several rinses in water, the floated film is then placed on a copper grid. For handling these thin films in an etching solution, it is convenient to make a small skimmer with a platinum mesh. It is difficult to find a solution that dissolves only the substrate and not the deposit. The stripping solutions and conditions that we developed are shown in Table 6.2. It is clear that the type of stripping solutions depends on the type of deposit – substrate combinations. If Ni or Ni alloy films are grown on pure iron substrates, stripping is extremely difficult because their ionization potentials are very close to each other. We describe our stripping method as an example (Tanabe and Watanabe, 1971). Here, we used a saturated nickel acetate solution as a stripping solution. The voltage was applied between a scarred deposit –substrate composite sheet as an anode and a carbon sheet as a cathode in this solution (see Figure 6.8(a)). By varying the voltage, we found the potential

Figure 6.7. Before floating a plated film off the substrate, a 2-mm square mesh is scarred on the surface of the deposit–substrate composite sheet with a knife.

219

Characterization Methods for Plated Films Table 6.2. Stripping solutions for plated films and their stripping conditions. Substrate

Plate metal

Reagent and condition

Reference

Fe, Fe –C

(a) Au, Pt, Cr

10% Nitric acid solution

(Saka, 1997)

Cu, Cu alloy

(b) Ni, Ni–P, Ni–B, Ni– Sn, Ni– S, Cr, Cr–W, Cr– Mo, Pd, Ag, Pt, Au, Au–Pd, Au–Ni, Ni– Fe–Mo (c) Co, Co–B, Co –Ni, Co– P, Co –Ni–B, Co– Fe, Co –Fe–S (d) Ni– Fe

Chromic acid anhydride 250 g/l; Sulfuric acid 25 g/l

(Hong and Saka, 1997)

Potassium cyanide 10 g/l

(Watanabe, 1982)

Fe

Plastics

(e) Ni, Ni– B, Ni –P

PP resin (f) Ni–P, Ni–B, Co– P, Co– B, Cu ABS resin

Mechanically Stainless (g) Ni and Ni–Co steel system composite peel off film for example Ni– TiH2, Ni– Al2O3 Cr plated (g) Ni– Fe film

Persulfuric acid ammonium 10%; (Watanabe, 1989) Ammonium solution 20% or Sodium cyanide solution 20 g/l (Watanabe, 1989) Electrolytically peel off method; Nickel acetate saturated solution; pH 4 adjusted with glacial acetic acid; voltage: 25 V; cathode: graphite plate

(Watanabe, 1995)

Xylene Methylethylketone

(Watanabe, 1996)

The thicker film than few mm is easy to peel off mechanically. Substrate must be polished by emery paper (No. 500) before plating.

(Tanabe and Watanabe, 1971)

(Sakai, 1977)

region where only the iron substrate dissolves. The voltage was , þ 0.2 V. This stripping solution was developed using a Poulbaix diagram (potential –pH diagram). This principle should be applicable to developing a stripping solution to other metal systems. An example of a TEM specimen is shown in Figures 2.24 and 2.26(b). An in situ heating experiment was also conducted inside a TEM using Ni –B alloy films prepared by the above electrolytic stripping method. TEM micrographs in Figures 4.11 –4.14 show the crystallization processes. If a plated deposit is not a uniform film but consists of islands, the stripping method will disintegrate the deposit. For an island-type deposit, a thin carbon film is deposited on the top surface of the deposit – substrate composite sheet, which is then dipped into a stripping solution. By doing this, the islands are fixed on the carbon film after stripping off the substrate. Since carbon film is amorphous and electron-transparent, it will not interfere with the fine structure of a plated deposit (Figure 6.9).

220

Nano-Plating

Figure 6.8. (a) An experimental setup for stripping nickel films off iron substrates electrolytically and (b) a dissolution rate of iron and nickel plotted against a potential.


221

Figure 6.9. A method for making TEM specimens suitable for observing the initial stages of electrodeposition.

In cases where a stripping solution cannot be found, we can use a very thin substrate and examine the deposit –substrate composite film together without any stripping operation. This method is illustrated in Figure 6.10. A thin gold film is first deposited on both sides of a pure iron foil substrate, and the gold deposit on the back surface is scraped off with a knife. After masking the edge region with a lacquer, the gold – iron foil is dipped into nitric acid. The exposed iron on the backside will dissolve. This process will create a thin gold membrane supported with an iron frame. This gold membrane will be the substrate for a to-be-plated film. After plating on one side of the membrane, we cut out a 2-mm square from the membrane and place it on a TEM grid for TEM observations. For this type of specimen, we will be examining both the plated film and the gold membrane simultaneously. 6.2.2.2. Plan-view observation of thick plated films (Watanabe, 1996) 6.2.2.2.1. Electropolishing method. To prepare plan-view TEM samples for thick plated films, electropolishing, ion-milling, and focused ion-beam (FIB) methods are used. Among these three techniques, the electropolishing method provides samples with the least damage.

222

Nano-Plating

Figure 6.10. A method for making pure gold film substrates.

An electropolishing machine for preparing reliable TEM samples from a sheet or a foil is available commercially (see Figure 6.11). TEM micrographs in Figure 1.25 show nickel specimens prepared by this machine. In the electropolishing method, a plated film has to be stripped off the substrate using stripping solutions listed in Table 6.2, so that both surfaces can be polished simultaneously through the twin jet nozzles. If the substrate is thin, we first polish off the substrate while the surface of the plated film is covered with adhesives or tape, and then polish the plated film from both surfaces. It is convenient to use stainless steel or titanium sheets as the substrate because films plated on these substrates can be peeled off mechanically with ease. These substrates are useful in obtaining plated films

Figure 6.11. A schematic illustration of a twin-jet electropolishing setup (C: a cathode, E: a DC power supply, H: an electrolyte entrance, L: a light source, M: an optical microscope for inspection, N: a nozzle, R: a stirrer, S: a specimen, W: a switch, and V: a transparent tank).


223

free of the substrate effect because the substrate is generally covered with an amorphous passivation layer, which does not cause any epitaxy to the subsequently deposited films. Preparation of a plan-view TEM sample from a certain depth requires several steps. First, the sample has to be polished from one side down to the desired depth. The polished surface is thus protected by an adhesive layer, and the other surface is then electropolished until perforation. By doing this, we can obtain a plan-view sample located at the desired depth. At the very moment that a specimen is perforated, a photo sensor detects light and shuts the voltage. For this reason, it is important to use a transparent protective layer such as an adhesive for light transmission. Figure 1.25 is a plan-view TEM micrograph showing the microstructure of various films obtained by the two-side electropolishing technique. The microstructure represents roughly the middle of the total depth. 6.2.2.2.2. Ion-milling method. The principle of the ion-milling method is described in Figure 6.12. This technique is similar to the sputtering method. A specimen is placed in the cathodic side and thinned with ionized argon gas. The ion beam is focused on the 3-mm specimen area. Two guns are arranged to sputter-deposit both surfaces. The specimen temperature is thought to rise up to 80 8C during ion milling. The ion-milling method is simple but sometimes presents difficulty in obtaining smooth surfaces due to preferential etching, especially when a plated film contains foreign phases. Table 6.3 lists sputtering rates for various metals. Figure 6.13 is Cu – Sn alloy films prepared by an ion-milling method. 6.2.2.3. Crystallographic matching relationship between plated film and substrate. As long as the substrate material is a metal, a metal grows with some epitaxial relationship on the substrate regardless of plating methods, i.e. electrolytic or electroless (see Table 1.1). Conversely, if the substrate is made of plastic or ceramic materials, the plated films will not form strong chemical bonds and thus crystallographic matching is not expected. In order to study the epitaxial relationship by a TEM, the plated film has to be examined together with the substrate or cross-section sample. Unfortunately, the cross-section method cannot fully provide information on the epitaxial relationship. Thus, although this method is more tedious, the former method has to be used. The difficult step of the simultaneous observation method is to thin the substrate to electron transparency and then to deposit a film. The total thickness of both the film and the substrate has to be thin enough for TEM observations. Misfit dislocations are present at the deposit –substrate interface. If the substrate is removed, interfacial phenomena such as the presence of misfit dislocations will not be observed. 6.2.2.4. Preparation of TEM cross-section samples. The cross-section of a plated film represents the microstructure along the film’s thickness. From a cross-section, it is possible

224 Nano-Plating

Figure 6.12. A schematic diagram of an ion-milling machine.

225

Characterization Methods for Plated Films Table 6.3. Sputtering rates for various metals. Ar 200 eV

Ar 600 eV

Element

Atom/ion

Element

Atom/ion

Ag Cu Au Pd Cr Ni Pt Co Fe Ir Ru Mo Re Os Al Hf V W Zr Ta Nb Ti Be

1.58 1.10 1.07 1.00 0.67 0.66 0.63 0.57 0.53 0.43 0.41 0.40 0.37 0.36 0.35 0.35 0.31 0.29 0.28 0.28 0.25 0.22 0.18

Ag Au (500 eV) Pd Cu Pt Ni Co Cr Ru Fe Al Ir Os Mo Re Hf Be Zr V Nb Ta W Ti

3.40 2.43 2.39 2.30 1.56 1.52 1.36 1.30 1.30 1.26 1.24 1.17 0.95 0.93 0.91 0.83 0.80 0.75 0.70 0.65 0.62 0.62 0.58

to trace back the growth history to see how the film has grown. Some cross-section samples can be seen in Figures 1.52 – 1.54. There are three methods for preparing cross-section TEM samples, i.e. (a) ultra-microtoming, (b) ion-milling, and (c) FIB milling methods. Advantages/disadvantages of each technique will be discussed in the following section. 6.2.2.4.1. Ultra-microtoming (Akakura 1977; Asakura 1999; Sakai 1977; Sakai 1989; Wood 1990) . For preparing cross-section TEM samples, an ultra-microtoming technique uses a diamond knife, which slices thin cross-sections of plated films grown on a substrate. This procedure is illustrated in Figure 6.14. The substrate used for plating metals should be a soft material, such as pure copper. The material should also be thin. A sheet containing plated film –substrate composites is cut into a 2 £ 15 mm2 strip using scissors. The strip is then placed vertically and embedded into a commercial plastic cell with a resin. For the resin, we used an epoxy or a lightsensitive resin. The property of these resins critically determines the degree of difficulty in slicing, therefore, it is important to consult an experienced engineer for the selection of

226 Nano-Plating

Figure 6.13. TEM micrograph and its diffraction patterns for an electrolytic Cu –3.4 at.% Sn alloy film prepared by an ion-milling method.


227

Figure 6.14. A method for preparing cross-section TEM specimens for plated films using a microtome method.

appropriate resins. After the resin is cured, we remove a resin-fixed specimen from the plastic cell and start trimming the edges carefully until a fresh cross-section appears. We continue to trim the exposed cross-section with a glass knife attached to the microtome. A thin section is then cut with a diamond knife. Sliced specimens are finally placed on a copper grid (see Figure 6.15(a)). The lateral dimension of these microtomed samples is 50 £ 200 mm2 : Thus, the specimen has to be trimmed with a glass knife down to this dimension prior to slicing with a diamond knife. The cross-section of a 2-mm thick plated film is visible at the lower left part of Figure 6.15(b). Since a microtomed metal specimen undergoes a severe plastic deformation during cutting and thus contains undesirable artifacts, it is often thought to be unreliable for TEM studies. Careful preparation of such samples, however, can provide interesting microstructural information.

228 Nano-Plating

Figure 6.15. Microtomed cross-section TEM specimens placed on a TEM grid.


229

Figure 2.25 is an example of a good sample. The film is bent as indicated by a straininduced bend contour in the copper substrate. Figure 6.16 is a multi-layer specimen containing alternating layers of Ni –Co – Cu alloy and pure copper. This multi-layer film was fabricated to test the presence of a giant magneto-resistance effect. The micrograph clearly reveals dendritic growth along the thickness direction (Figure 6.16(a)). At higher magnification, alternating black –white lines from the multilayer are seen to follow along the profile of dendrite arms (Figure 6.16(b)). Specimens shown in Figures 1.28, 1.52, 1.53, 1.54, 2.20, and 2.28 were also prepared by microtoming. Since the edge angle of a commercial diamond knife is 458 and the escape angle is 78, the cutting will be done at the total angle of 528. Due to this large angle, a specimen tends to be pushed up during slicing and the side face of the specimen is rubbed against the diamond edge (Figure 6.17). If the edge is new and sharp, it is possible to cut a specimen with less plastic deformations. As the edge becomes dull with usage, however, various artifacts will appear on the cut specimen surface as seen in Figure 6.18. Figure 6.18(a) is a TEM micrograph showing the cross-section of a plated film on a substrate obtained by microtoming. The specimen is wrinkled along the interface direction, which is created during cutting with a diamond edge. Structural information along the film thickness direction, however, appears to be preserved without any artifacts. Microtomed TEM specimens were also examined using a scanning electron microscope (SEM). Figure 6.18(b) and (c) are SEM micrographs showing one surface touched by the cutting edge and the opposite untouched surface, respectively. Both surfaces are seen to be wavy along the cutting direction. Figure 6.18(d) is the magnified picture of (b), indicating that the specimen is apparently rifted due to a large plastic deformation. In addition, the deposit –substrate interface is no longer a straight line, but a curved line instead. When the grain size of a plated film is large, each grain undergoes different modes of plastic deformation according to their available slip system for deformation. Such an anisotropic deformation also contributes to the formation of an uneven specimen. As the knife-edge becomes dull, the cutting efficiency deteriorates. The cutting also gets worse with increasing slice thickness and increasing cutting speed. It is important to note that there is a work-hardened zone created by a glass knife on a specimen surface prior to the final diamond slicing (see Figures 6.1(b) and 6.2). The edge of a glass knife is more brittle and becomes dull faster than a diamond knife. For this reason, the workhardened zone generated by a glass knife will be thicker than that of a diamond knife (c.f. Figure 6.19). To lower the thickness of the work-hardened zone, a fresh glass knife should be used and a thinner slice should be cut. After the glass trimming, several slices of specimens are cut with a diamond knife and are placed in water. Next, the sliced specimens are scooped on a copper grid. If the grid surface is oxidized or contaminated, it tends to repel water, making the scooping job difficult. In this case, it may be helpful to

230

Nano-Plating

Figure 6.16. A microtomed cross-section TEM specimen of a Ni–Co –Cu multi-layered film grown on copper (Courtesy of Prof. Seizo Kainuma).

give the grid a short etching in a dilute nitric acid solution. By doing this, the copper grid becomes wettable and the specimen scooping becomes easier. There are four possible specimen-cutting edge configurations as seen in Figure 6.20. Figure 6.20(c) is the best direction. In configuration (a), a plated film tends to separate from the substrate during cutting. In (b), a plated film will be deformed along the thickness direction by a compressive force exerted with the knife-edge. In (c), a plated


231

Figure 6.17. A schematic view illustrating how a TEM specimen is made with a diamond edge.

film is pushed down during cutting and thus is not easily separated. In (d), a plated film is likely to separate from the substrate. A knife-edge scar called a knife mark can be introduced on the surface of a specimen (see Figure 1.52 –1.54). If this scratch mark is parallel to the interface, it is difficult to distinguish it from the layer structure formed along the film thickness direction. It is impossible to make a cross-section specimen for a plated film alone because it is brittle and may disintegrate during cutting. 6.2.2.4.2. Ion milling. The basic principles for obtaining TEM samples by an ion-milling (or ion-thinning) method are explained in Section 6.2.2.2.2. One of the specimen preparation methods for making TEM cross-sections is described in Figure 6.21. Several plated specimens are sandwiched together with glue and then inserted into the thin (0.5-mm wide) slot of a 2-mm diameter brass rod. The brass rod is then inserted into a metal tube with a 3 mm outer diameter and 2 mm internal diameter. Next, epoxy is poured into the tube. Care should be taken not to introduce air bubbles in the epoxy. Also, it is important to control the amount of epoxy. If the resin is too thin, adhesion to the specimen is weak, whereas if it is too thick, the resin will be milled preferentially and the specimen may fall out of the tube. Once the epoxy is cured, a thin (, 0.25 mm) disc is sliced mechanically from the tube, followed by mechanical thinning to about 100 mm using an emery paper or buffing. The central area of the sample is thinned using a dimpler. If both surfaces can be dimpled as shown in Figure 6.21(c), it is ideal, but this may not be possible in practice. If not possible, then at least one side should be dimpled. The dimpled sample is polished in an ion-milling machine from both sides.

232

Nano-Plating

Figure 6.18. TEM/SEM examinations of specimens prepared with a microtome. (a) A TEM image, (b) an SEM image of the free surface, (c) an SEM image of the cut face, (d) an SEM image of the free surface, and (e) an SEM image of the cut face.


233

Figure 6.19. A schematic diagram showing how a layer of work-hardened zone forms on a sample during cutting with (a) a glass knife and (b) a diamond knife.

A profile of the specimen during the ion-milling process is shown in Figure 6.21(d). Such an uneven profile originates from the difference in milling rates for a plated film, substrate, and resin. One way to reduce this preferential milling is to rotate the specimen stage in such a way that ion milling is performed only when the direction of an ion beam becomes perpendicular to the long direction of the specimen. One practical method for this is to shut off, or shield, the ion beam when it becomes parallel to the long direction of the specimen. By doing this, the ion-milling time can become lengthy, ranging from a few hours to several days. A commercial ion-miller is generally equipped with a terminator, which stops the machine by detecting the initiation of a small hole in the specimen.

Figure 6.20. Relationship between specimen and cutting direction.

234 Nano-Plating

Figure 6.21. Steps for preparing a cross-section TEM specimen of plated films by an ion-milling method.


235

However, the location where a hole initiates cannot be controlled. Milling rates for various metals are listed in Table 6.3. Figures 1.32 and 6.22 are examples of ion-milled samples. The micrograph in Figure 6.22 is the cross-section of an electrolytic Ni– 31.6 at.% P alloy film grown on a copper substrate. The Ni – 31.6 at.% P/Cu interface is clearly visible. From this micrograph, we note that there is no sign of plastic deformation in either the film or the substrate. The plated film actually contained fine-grained and amorphous regions. Selected area diffraction patterns from three regions (circled with white) are shown in Figure 6.22(a)– (c). An XRD pattern from this specimen is also shown on the left-hand side of the figure and indicates the presence of various intermetallic compounds related to Ni– P binary alloys. Figure 6.23 is another example of a specimen prepared by an ion-milling method. The micrograph shows a Ni/Cu multi-layer fabricated by varying the potential periodically in a single bath. Again, no plastic deformation is visible in this specimen and the Ni/Cu multilayer is clearly visible. Furthermore, a characteristic bend contour is continuous across the Ni –Cu interfaces, indicating that each layer has grown epitaxially to become a single crystal. 6.2.2.4.3. Focused ion-beam (FIB) method (Nikaw et al., 1989; Adachi, 1991; Phaneuf et al., 1997). An FIB method uses a focused (, 0.1 mm in diameter) gallium beam to prepare TEM specimens. As long as the specimen is a solid, the FIB can cut it regardless of its stiffness. This is ideally suited for preparing the cross-section TEM specimen of plated films. One advantage is that the sample can be imaged simultaneously using secondary electrons emitted while being cut. This allows us to monitor the location and polishing condition simultaneously. These images provide high-contrast images, which cannot be obtained by a SEM.

6.3. GRAIN-SIZE MEASUREMENTS

6.3.1 Direct measurements by TEM observations One method for measuring grain sizes in thin films is to use TEM for direct observations. Two types of specimens, i.e. plan-view and cross-section samples have to be prepared for this measurement. An example of a plan-view sample prepared by an ion-milling method is shown in Figure 6.13. It is important to remember that the grain varies in shape, e.g. columnar or elliptical and its size also changes with the film depth. Grain-size measurements for columnar-, elliptical-, or needle-shaped grains were described in Section 6.2.2.4.2. The cross-sectional view is best for obtaining information on the grain shape. Well-prepared TEM

236 Nano-Plating

Figure 6.22. A cross-section TEM specimen of an electrolytic Ni –31.6 at.% P alloy film plated on a copper substrate prepared by an ion-milling method.


237

Figure 6.23. A cross-section TEM specimen of a Cu/Ni multi– layer film prepared by an ion-milling method (Courtesy of Dr. Go¨ran Holmbom).

cross-section specimens can provide both dimensional and crystallographic information along the lateral and vertical directions of elongated grains (Figure 1.31). Care should be taken for a microtomed specimen because the grains may be highly distorted as a result of cutting with a diamond knife (see Figure 6.18(b) and (c)). Finegrained films are resistant to a plastic deformation and are thus not easily damaged by microtoming. Therefore, it is possible to obtain reliable grain-size measurements from these films (see Figure 1.26). Figure 2.20 is another successful example. If the grains have a spherical shape and are distributed uniformly along the film thickness direction, their size can be measured from the film’s normal direction by preparing plan-view samples. Such samples can be made using electropolishing or ionmilling methods. If the grain size is less than the film thickness, the overlapped images of grains can be seen. Sometimes this image overlapping makes the size measurement more difficult. In this

238

Nano-Plating

case, it is convenient to use dark-field (DF) images instead of bright-field images because DF images have less image overlapping. This example is seen in Figures 4.13 and 4.14 (d0 ). Conversely, if the film thickness is less than the grain diameter, it is not easy to determine the size precisely. In the case of nearly amorphous and fine-grained films, a lattice imaging technique can be applied (see in Figures 6.24 and 2.28(c)). The absence of lattice images in a film does not necessarily indicate the absence of crystalline phases because crystals may simply happen to be oriented away from the strong Bragg condition. 6.3.2 X-Ray diffraction method (Scherrer, 1920; Warren, 1934; Phaneuf et al., 1997) One of the simplest ways to measure the grain size is by an XRD technique. Assume that the average shape of the grains is spherical and that the grains are distributed uniformly. The grain size can be determined by the full width at half maximum (FWHM) of a diffraction peak using Scherrer’s equation: Dhkl ¼ K l=B cos u

ð6:2Þ

where Dhkl is the grain size in the direction normal to the reflecting plane {hkl}; l, the X-ray wavelength; B1, the FWHM (in radian) due to the grain size; u, the Bragg reflection angle; and K, the Scherrer’s constant (2 0.7 , 1.7). If the grain is a cube and their size is uniform, we obtain D ¼ Dhkl and K ¼ 0:89: The desirable intensity value at the half-maximum should be beyond 10,000 counts after a subtraction of background. B can be determined using the following equation: B2 ¼ B2M 2 B2S

ð6:3Þ

where BS is the FWHM of a standard sample having a grain size more than 25 mm and BM, the measured FWHM of a material. The applicable grain-size range of this equation is 1 , 100 nm: As already discussed, grains of plated films are often elongated along the film thickness direction, forming columnar, elliptical, or needle-shaped grains. For such grains, it is necessary to measure the grain size from both the long- and short-axis directions. The choice of diffraction peaks depends on the purpose of the study. We used strong diffraction peaks because we can evaluate the grain size easily with such peaks. However, it is questionable whether such an arbitrary choice is valid for obtaining the correct grain size. If the film has a texture, diffraction peaks originating from the planes normal to the preferred orientation are used. Nevertheless, it appears more logical to choose diffraction peaks after determining the grain shape and its distribution by some other means. The shape of the diffraction peaks also changes with a lattice distortion in the crystal (Unga´r et al., 1999). In addition to the grain size, a layer spacing in multi-layer films can 1 Generally the FWHM of a diffraction peak is used but the area under the peak divided by the peak intensity can also be used (von Laue, 1926).

Characterization Methods for Plated Films Figure 6.24. High-resolution TEM lattice images of electrolytic Ni –P films. In this example, the lattice images were used to measure crystal sizes. (a) Ni – 20.0 at.% P and (b) Ni– 24.7 at.% P alloy films.

239

240

Nano-Plating

also be determined from the positions of satellite peaks appearing around the main diffraction peaks (Segmuller and Blakeslee, 1973; Michaelsen, 1995) (see Figure 5.2).

6.4. OBSERVATIONS OF SURFACE MORPHOLOGY

The surface morphology of plated films can be studied using an optical microscope, an SEM, or an atomic force microscope (AFM). In addition to the microscope techniques, other surface characterization methods, such as a profilometer, provide a quantitative value of surface roughness. 6.4.1 Scanning electron microscope (SEM) (Watanabe, 1995) Imaging of surface morphology is easily performed using a SEM by inserting a specimen, which does not require lengthy specimen preparation, into a vacuum chamber. Like metals, SEM specimens have to be electrically conductive. Ceramic or plastic materials can also be examined if the surface is coated with a thin conducting film, such as gold or gold – palladium. A modern SEM operated at low accelerating voltages is capable of imaging organic or inorganic materials without a conductive coating, although its resolution is somewhat lower. Furthermore, an environmental SEM is available to examine wet materials or liquid droplets at ambient atmosphere. Nevertheless, metal-coated specimens still provide the best images in terms of sharpness and resolution. Metal specimens can become susceptible to electron-beam damage if their surface is covered with a passivation film. Such specimens cannot be imaged at high resolution. If so, they usually yield images of poor quality. Materials giving low-contrast images should always be coated with a gold or a gold – palladium film by sputtering. If the sputtering time is too short, the coated metal will not be in a continuous film but in an island form. For example, gold deposits in Figure 2.2(a) –(c) are a discontinuous film and thus are not suitable as the coating layer. In our experience, 5-min sputtering appeared to be necessary to obtain a continuous film of 15 nm in thickness. Oneminute sputtering yielded a discontinuous islanded film with the average grain size of 5 nm. An SEM technique can be applied to study various properties of plated films. One of our applications is shown below as an example. Figure 6.25(a)– (d) illustrates the experimental steps in preparing a SEM specimen for studying the fracture behavior of plated films. In this application, a thin strip is cut with scissors from a composite sheet containing a plated film and the substrate. A notch is cut at the middle section of the strip and then this strip is pulled apart from each end so that a fracture occurs at the bottom of the notch. By doing this, a fracture surface containing both the film and the substrate can be obtained for SEM observations. After the fractured specimen is bent to 458 at the middle section, it is glued on a SEM holder through conducting tape (see Figure 6.26(a)).


Figure 6.25. A method for preparing a fractured sample of plated films for SEM observations.

241

242

Nano-Plating

Figure 6.26. Convenient specimen configurations for observing both the surface morphology and the fracture surface of plated films by SEM (I).

For hard-to-fracture specimens, cracks can be introduced by applying 1808 bending to the specimens. In these types of specimens, cracks are oriented at various angles and thus can be examined simultaneously at different viewing angles without tilting the specimen (see Figure 6.27(a)). For observation of the fracture characteristics of plated films as a function of plating conditions, multiple specimens can be mounted on a single SEM stage (see Figures 6.26(b) and 6.27(b)). This type of specimen configuration is also beneficial in quickly surveying an experimental trend as to how various plating conditions affect the fracture behavior. After this quick survey, one can go back to a specific specimen for more detailed examinations. From both quick and detailed observations, we can formulate a consistent picture of fracture characteristics in plated films. For conducting comparative studies, it is important to take micrographs at the same magnifications and at the same specimen tilt angles. The degree of magnification depends on the feature size of interest. The proper selection of

Figure 6.27. Convenient specimen configurations for observing both the surface morphology and the fracture surface of plated films by SEM.


243

location, magnification, and specimen tilt angles, of course, requires a sound knowledge of the overall trend. Specimen configurations shown in Figures 6.26(a) and 6.27(a) allow us observe both the surface morphology and the fracture surface. The morphology of the fracture surface provides important material information on the structure and properties of plated films. As illustrated in Figure 1.13, the surface morphology of plated films changes with film thickness. The fracture surface can provide information on the film thickness at which the surface roughness/leveling changed or a layer structure formed. At the same time, the surface irregularities can be examined by tilting a specimen. Specimen tilting is needed because the shape and height of surface irregularities are not easily observed from the film normal direction, but are recognized in a tilted specimen as seen in Figure 6.28. Another example demonstrating how observation angles affect the images is shown in Figure 6.29. In this figure, the observation angle was changed by tilting the specimen. The surface image taken from the normal direction appears to be completely different from the image observed at an angle tilted away from the film normal. Note that the location is identical as the corresponding sites are numbered as (1) – (5). The surface image viewed from the normal direction contains rounded features (Figure 6.29(a)), whereas the surface of a tilted specimen shows pointed crystals (Figure 6.29(a)). The latter micrograph undoubtedly represents the correct surface morphology while the former does not. The reason why such markedly different images were obtained in the SEM can be explained using the origin of SEM contrast. In the SEM, an incident electron beam (primary electrons) excites free electrons inside a specimen and the excited electrons (called secondary electrons) are subsequently emitted from the specimen surface (see Figure 6.30). The quantity of these secondary electrons, which are collected by a detector, changes with the location of the surface and contribute to variations in the image contrast. A pointed site produces more secondary electrons and thus appears white, whereas a flat region emitting fewer, or zero electrons shows up black. As a result, surface irregularities can be seen as black – white images. In the normal accelerating voltage of 20 , 30 kV, a large quantity of secondary electrons are emitted from the tip of pointed crystals. From the normal direction, therefore, these electrons dominate the image contrast and tend to wash out details of surface features lying below the tip region. Consequently, these pointed crystals are seen as rounded white features without morphological details of the surrounding area. If the specimen is tilted, the side face of these pointed crystals become sites for enhanced secondary electron emission. This orientation, therefore, provides a better representation of surface morphologies. One of the unavoidable mistakes we often encounter in SEM imaging is an illusion that protrusions and valley regions are switched if the micrograph is placed upside down. This example is shown in Figure 6.31. In Figure 6.31(a), the surface appears to consist of pointed crystals, whereas in (b), it shows more rounded crystals. The two micrographs are actually identical except that (b) is

244 Nano-Plating

Figure 6.28. An SEM image change for the surface of plated nickel films due to different observation directions.


245

Figure 6.29. An SEM image change for the surface of plated nickel films due to different observation directions. SEM images (a) taken from the normal direction and (b) taken at an angle tilted away from the film normal. The corresponding locations are indicated with numbers from 1 to 5.

Figure 6.30. An SEM contrast formation mechanism.

246

Nano-Plating

Figure 6.31. An optical illusion arising from placing SEM micrographs upside down.

placed upside down. The problem in dealing with this illusion is that the correct orientation of the micrograph is known only to the person who took the micrograph. Therefore, one should be aware of this type of confusion in SEM micrographs and be sure that the micrograph is placed in a correct orientation. 6.4.2 Replica method for TEM observations One of the weaknesses in a conventional SEM is its inability in taking high-contrast micrographs from smooth surfaces. Imaging of fine-scale surface irregularities is limited by the resolution of the SEM. This difficulty is, of course, connected to the fact that the quantity of secondary electrons is the same everywhere on a smooth surface, thus the resulting images have no contrast variations. An alternative method is to prepare a surface replica and to image it by a TEM. An example of surface replica images is given in Figures 2.6(1) and (2). It should be pointed out that these surface features did not show up clearly by SEM imaging. A method for making replica specimens is illustrated in Figure 6.32. In this technique (a two-stage replica method), a drop of acetone is first placed on the surface of plated films


Figure 6.32. TEM specimen preparation steps using a replica technique.

247

248

Nano-Plating

(Figure 6.32(a)), followed by the placement of an acetylcerulose film (Figure 6.32(b)). After the film is dried, it is peeled off the surface (Figure 6.32(c)). This is the first-stage replica. Next, this replica is taped down with its replicated side up on a glass slide and given shadowing with chromium at an incidence angle of 458 (Figure 6.32(d)). A thin carbon film is then deposited on the shadowed surface at the normal incidence to enforce the chromium film (Figure 6.32(e)). The specimen is then taken out of the vacuum chamber. A square mesh pattern with 2-mm spacing is scarred on the surface with a knife. The replica is then soaked in acetone. After the acetylcerulose film is dissolved in acetone, the piece of the replica having the 2-mm square dimension will be released. The TEM replica specimen is then scooped with a platinum mesh, (Figure 6.32(f)). Sites having thick chromium films give a dark contrast, whereas those with thin chromium films give a white contrast (Figure 6.32(g)). Although steps for making the two-stage replica are lengthy, it is not too difficult. It is important to remember that there is an optimum thickness for both chromium and carbon films. 6.4.3 Observations by AFM AFM can image the surface morphology, but can also provide a quantitative value of the surface roughness. Figure 6.33 compares the effect of three types of brightener on the surface morphology of electrolytic nickel films. Both SEM and AFM micrographs are pictured next to each other for comparison. Root mean square (RMS) values indicating the average roughness along the film normal are also given. Although the RMS value is the same, the surface morphology could be different. The reason for this difference is that the RMS value measures only the average surface roughness and not the wavelength of the RMS over the film surface. 6.4.4 Measurements by a surface profilometer (Mori and Yokoi, 1998) Commercial surface profilometers are available in both contact and non-contact modes. Surface roughness can be measured from fine irregularities excluding a geometrical deviation or a long-wavelength undulation. More details on this technique can be found in publications by the Japan Industry Standard (JIS).

6.5. MEASUREMENT OF PREFERRED ORIENTATION

As shown in Figure 1.31(a), preferred orientation (or texture) is a phenomenon in which a number of fine crystals in a plated film grow in a similar crystallographic direction normal to the film plane. If the film has a preferred orientation, say hkl, the probability of the hkl plane appearing on the surface of a plated film is high. Conversely, if the film does not have the preferred orientation, the probability of any oriented crystals covering the surface is the same for all orientations.


249

Figure 6.33. SEM and AFM images showing the effect of brighteners on the surface morphology of electrolytic nickel films. (a) Brightener: No, RMS: 334, Appearance: Black; (b) Brightener: 1,5-naphthalenedisulfonic acid, disodium salt, RMS: 336, Appearance: Black; (c) Brightener: 1,3,6-naphthalenedisulfonic acid, trisodium salt, RMS: 257, Appearance: Bright but cloudy; (d) Brightener: 1-naphthalenedisulfonic acid, sodium salt, RMS: 114, Appearance: Bright.

250

Nano-Plating

An example of the preferred orientation is shown in Figure 1.32. Here, the cross-section TEM micrograph of electrolytic Ni –7.6 at.% P alloy film prepared by an ion-milling method and its electron/XRD patterns are shown. From the TEM micrograph, extremely fine needle-shaped crystals are seen to grow along the film thickness direction. The {111} spot/peak in electron and XRD patterns is the strongest and thus the film has a 111 preferred orientation. The preferred orientation varies with the type of crystal systems and metals, and sometimes changes significantly with plating conditions. Here, we will present a description of preferred orientation and its measurement technique. 6.5.1 Definition of preferred orientation and its description It appears that no unified description is available for defining a preferred orientation in plated films. We will show how to define the preferred orientation and how to express it. If the preferred orientation of plated films is to be used in a practical sense, it suggests two different meanings. The first deals with the question as to which crystal axis is vertical to the substrate plane. The second addresses the question as to which crystallographic plane is statistically most abundantly present on the surface of plated films. Both meanings are identical in a crystallographic sense, but they have to be changed according to the application. For example, for describing perpendicularly magnetized films, the direction of magnetization is more relevant, so the use of the indices of the line direction, [hkl ] or , hkl . , is more appropriate. Alternatively, the indices of the plane, (hkl) or {hkl}, are used in describing film surface (plane) properties, such as corrosion-resistance characteristics or catalytic property, adhesion property between paint and substrate. In this book, we simply use hkl by omitting the parenthesis in order to imply the two meanings, as used in the JCPDS files. As explained in the Section 1.4.4, the preferential growth plane and direction of a crystal are the same but do not necessarily coincide. Therefore, it is important to use this terminology carefully. 6.5.2 Measurement of texture in plated films A pole figure is generally used to display the degree of the in-plane texture in a foil or a sheet. The pole figure is not suited to plated films exhibiting a preferred orientation in the film thickness direction. In this case, the pole figure has to be taken in the cross-section. Such a measurement is not only time-consuming, but also impractical. Most recently, a new technique (Hebesberger et al., 2001; Suzuki, 2001) was developed to determine the orientation of individual grains using backscattered electron diffraction patterns taken inside the SEM. In this technique, backscattered electron


251

diffraction patterns are taken by tilting a specimen to 708. These patterns are stored and analyzed in a computer. The orientation of each grain is determined and subsequently each grain is color-coded according to its orientation. This automated system allows one to obtain a large amount of orientational information for a number of grains. This method could open a new avenue of investigation for analyzing crystal orientations on a nanoscale level. Here we describe a simple method of determining orientations using an XRD pattern. First, we take an XRD pattern from a plated film and compare it with the standard intensity of a randomly oriented powder specimen listed in the JCPDS file (Wilson and Rogers, 1964). With this method, the texture index, N{111}, of the {111} plane in an fcc crystal can be computed as: N{111} ¼

IF{111} IFR{111}

ð6:4Þ

Here, IF{111} is the relative intensity of an XRD intensity from the {111} plane and can be computed from the following equation: IF{111} ¼

I{111} I{111} þ I{200} þ I{220} þ I{311}

ð6:5Þ

where I{hkl} is the intensity of an XRD from the hkl plane. IFR{111} is determined by Eq. (6.5), with the exception that the I{hkl} intensity values are taken from the corresponding randomly oriented powder sample, such as those used for making the JCPDS file. The texture index for other low-index planes, such as {200}, {220}, and {311}, can be determined in the same manner. Although in a strict sense, all diffraction peaks are needed, only four or five low-index planes are used for determining the texture index, N{hkl}. If the texture happens to be a high-index plane, a low-index plane closest to the highindex plane will be assigned. If a plated film does not have a texture, the texture index is 1 for all planes. If the film has a texture, the texture index of some planes becomes greater than 1 and a plane having the largest texture index will be taken as the preferred orientation. From computed X-ray intensities in fcc crystals, we obtained the following texture index for the {111}, {110}, {100}, and {311} planes. The maximum texture index value is 1.83 for the {111} plane, 3.978 for the {110} plane, 9.150 for the {100} plane, and 10.764 for the {311} plane. These computed values are for pure copper metal, but are expected to vary with different types of metals. Therefore, small texture index values do not necessarily imply the weak texture. For this reason, the degree of texture is sometimes expressed by the percent change from the

252

Nano-Plating

corresponding maximum texture index. It should be noted that the preferred orientation of 100 is the same as that of 200. Similarly, the preferred orientation of 110 and 111 is the same as that of 220 and 222, respectively.

REFERENCES Adachi (1991) Semicon/Kaisai-Kyoto Technology Seminar 19, June. Akakura, H. (1977) Denshi Kenbikyo, 12, 21. Asakura, K. (1999) J. Surf. Finish. Soc. Jpn, 50, 722. Cullity, B.D. (1978) Elements of X-Ray Diffraction, Addison-Wesley Publishing Company, Inc. Hebesberger, T., Schafler, E., Zehetbauer, M., Pippa, R., Unger, T. & Bernstoff, S. (2001) Z. Metall., 92, 410. Hong, M.N. & Saka, H. (1997) Script. Met., 36, 1423. Japanese Society of Electron Microscopy Summer School (1994) TEM Fundamentals & Applications—Nano-scale Local Analysis, Gakusai Kikaku Co. Japanese Society of Electron Microscopy Summer School (1995) TEM Fundamentals & Applications—Advanced Specimen Preparation Techniques, Gakusai Kikaku Co. Japanese Society of Electron Microscopy Summer School (1996a) TEM Fundamentals & Applications—Dynamic Analysis Techniques for Nano-structures, Gakusai Kikaku Co. Japanese Society of Electron Microscopy Summer School (1996b) TEM Fundamentals & Applications—Road to Nano-world, Gakusai Kikaku Co. Kanto Division of Japanese Society of Electron Microscopy, Editor (1991) Electron Microscopy Techniques for Evaluation of Advanced Materials, Asakura Book Co. Michaelsen, C. (1995) Phil. Mag., A72, 813. Mori, I. & Yokoi, M. (1998) J. Surf. Finish. Soc. Jpn, 49, 963. Nikaw, N. et al. (1989) International Reliability Physics Symposium, 27th Ann. Proc., 1. Phaneuf, M.W., Rowlands, N., Carpenter, G.J.C. & Sundaram, G. (1997) Specimen Preparation for Transmission Electron Microscopy of Materials IV. Symp., San Francisco, CA, April, p. 39. Saka, H. (1997) Crystal Electron Microscopy for Materials Scientist, Uchida Co. Sakai, T. (1977) Denshi Kenbikyo, 12, 137. Sakai, T. (1989) Denshi Kenbikyo, 24, 121. Scherrer, P. Bestimmung der Grosse und der Inneren Struktur von Kolloidteilchen Mittels Ro¨ntgenstrahlen, Nach. Gesell. Wiss. Go¨ttingen, Zitzungsber., July 26, 1918, in Zsigmondy, R., Kolloidchemie (Otto Spamer, Leipzig, 1920), 3rd Edition. Segmuller, A. & Blakeslee, A.E. (1973) J. Appl. Cryst., 6, 19. Sotomura, A. (1989) Electron Microscopy Techniques, Maruzen Publ. Co. Suzuki, S. (2001) Mater. Jpn, 7, 612. Tanabe, Y. & Watanabe, T. (1971) Electron J. Microsc., 20, 131. Tanaka, M., Terauchi, M. & Tsuda, K. (1997) Easy Electron Diffraction & Elementary Crystallography, Kyouritsu Publ. Co. ´ . & Borbe´ly, A. (1999) J. Appl. Cryst., 32, 992. Unga´r, T., Dagomir, I., Re´re´sz, A von Laue, M. (1926) Z. Krist., 64, 115. Warren, B.E. (1934) J. Chem. Phys., 2, 551. Watanabe, T. (1982) J. Metal Finish. Soc. Jpn, 33, 318. Watanabe, T. (1989a) J. Surf. Finish. Soc. Jpn, 40, 280.

Characterization Methods for Plated Films Watanabe, T. (1989b) J. Surf. Finish. Soc. Jpn, 40, 1355. Watanabe, T. (1995) Denshi Kenbikyo, 29, 158. Watanabe, T. (1996) Kagaku to Kogyo (Sci. Ind.), 70, 71. Watanabe, T. & Nakata, T. (1994) J. Surf. Finish. Soc. Jpn, 45, 1100. Wilson, K.S. & Rogers, J.A. (1964) Tech. Proc. Am. Electroplaters Soc., 51, 92. Wood, G.C (1990) Trans. Inst. Metal Finish., 7, February.

253


Chapter 7

Database for the Microstructure of Plated Films 7.1. 7.2.

7.3. 7.4.

Introduction Plating methods 7.2.1 Plating conditions 7.2.2 Substrate materials 7.2.2.1 Polycrystalline copper sheets 7.2.2.2 Single-crystal copper ({100}, {110}, and {111} planes) 7.2.2.3 Stainless steel sheets (SUS-304) 7.2.2.4 Electrolytic Ni-25 at.% P alloy films 7.2.2.5 Splat-quenched amorphous alloy foils (Fe –Si –B or Fe– Co –Si –B alloys manufactured by Nippon Amorphous Metals Co., Ltd.) Microstructure observations and measurements methods for plated films Database for the Microstructure of Plated Films 7.4.1 Pure Metals 7.4.1.1 Electrolytic Ag 7.4.1.2 Electrolytic Au 7.4.1.3 Electrolytic Cd 7.4.1.4 Electrolytic Co 7.4.1.5 Electrolytic Cr 7.4.1.6 Electrolytic Cu 7.4.1.7 Electrolytic Fe 7.4.1.8 Electrolytic Ni 7.4.1.9 Electrolytic Sn 7.4.1.10 Electrolytic Zn 7.4.2 Pure Alloys 7.4.2.1 Electrolytic Ag – Cd 7.4.2.2 Electrolytic Ag – Co 7.4.2.3 Electrolytic Ag – Cu 7.4.2.4 Electrolytic Ag – Sn 7.4.2.5 Electrolytic Ag – Zn 7.4.2.6 Electrolytic Al– Mn 7.4.2.7 Electrolytic Au – Cu 7.4.2.8 Electrolytic Au – Ni 7.4.2.9 Electrolytic Au – Pd 7.4.2.10 Electrolytic Au – Sn

257 257 257 258 258 258 258 259 259 259 260 260 260 274 296 308 318 323 333 344 358 369 384 384 387 398 407 416 417 423 427 437 446

256

Nano-Plating

7.4.3

7.4.4

7.4.2.11 Electrolytic Cd – Sn 7.4.2.12 Electrolytic Cd – Zn 7.4.2.13 Electrolytic Co – Cu 7.4.2.14 Electrolytic Co – Fe 7.4.2.15 Electrolytic Co – Mo 7.4.2.16 Electrolytic Co – Ni 7.4.2.17 Electrolytic Co – Sn 7.4.2.18 Electrolytic Co – W 7.4.2.19 Electrolytic Cr –H 7.4.2.20 Electrolytic Cu – Ni 7.4.2.21 Electrolytic Cu – Pb 7.4.2.22 Electrolytic Cu – Sb 7.4.2.23 Electrolytic Cu – Sn 7.4.2.24 Electrolytic Cu – Zn 7.4.2.25 Electrolytic Fe –Mo 7.4.2.26 Electrolytic Fe –Ni 7.4.2.27 Electrolytic Fe –W 7.4.2.28 Electrolytic Fe –Zn 7.4.2.29 Electrolytic In – Sn 7.4.2.30 Electrolytic Ni –B 7.4.2.31 Electrolytic Ni –Mo 7.4.2.32 Electrolytic Ni –P 7.4.2.33 Electrolytic Ni –S 7.4.2.34 Electrolytic Ni –Sn 7.4.2.35 Electrolytic Ni –W 7.4.2.36 Electrolytic Ni –Zn 7.4.2.37 Electrolytic Sn –Zn Electroless Plating 7.4.3.1 Electroless Ni– B 7.4.3.2 Electroless Ni– P Displacement Plating 7.4.4.1 Displacement Ag 7.4.4.2 Displacement Au 7.4.4.3 Displacement Cd 7.4.4.4 Displacement Cu 7.4.4.5 Displacement Zn

455 457 459 466 476 482 495 501 510 518 525 532 533 544 546 555 566 577 589 599 605 611 624 630 642 652 655 657 657 666 674 674 680 686 690 692

Chapter 7

Database for the Microstructure of Plated Films 7.1. INTRODUCTION

As illustrated in Figure 1.2 of Chapter 1, the microstructure of plated films can be classified into seven types. In this chapter, we will systematically present: (1) the crystalline structure (phase structure); (2) the grain size (crystal size); (3) the surface morphology; and (4) the texture (preferred orientation) of various pure metal and alloy electrodeposits. As discussed before, the basic idea for explaining the microstructure of plated films is that The microstructure and crystal size of plated films depend primarily on its film composition but not on plating conditions. Although the plating solution temperature and the overpotential also affect the microstructure and grain size, the effect is considered to be secondary. In other words, if the composition is the same, then the microstructure and grain size of plated films should be same regardless of the plating conditions. The surface morphology and the texture, however, vary markedly with plating conditions. For pure metal electrodeposits, we will first discuss how the grain size, the surface morphology, and the texture change with plating conditions. For alloy deposits (binary alloys only), we will describe how the microstructure and the grain size vary with composition by referring to the corresponding equilibrium phase diagram. We will also explain their surface morphology and texture using experimental micrographs and graphs. We hope that this chapter will serve as a database for the microstructure of pure metal and alloy electrodeposits. Only the minimum amount of experimental results will be covered, and thus for details, readers should refer to the original papers listed in cited references. 7.2. PLATING METHODS

We first explain the experimental methods used in obtaining the data contained within this database. In order that we can compare experimental results among different metal/alloy electrodeposits, the plating methods have been standardized unless otherwise stated. 7.2.1 Plating conditions The quantity of plating solution used was 0.5 l and the solution was purified through an activated charcoal filter and an ion exchanger. Distilled water was used to make up the 257

258

Nano-Plating

solution by adding commercial super-grade reagents as metal salts. The pH was adjusted using a buffer containing the same anion as metal salts. The plating was carried out, if possible, without additives. For example, a brightener or a stress reducer was not added even if surface roughening, dendrites, or cracks were developed on the film surface. The main reason for using an additive-free solution is that additional atomic elements might go into the deposit from additives in the solution and affect the final microstructure. If an additive was used, we made sure that no additional elements were included in the deposits. The plating was conducted at a constant current density while the plating time was determined using a coulomb meter. The constant-voltage plating was performed using a saturated calorimeter electrode. The plating tank was made of Pyrex glass. † Anode: a platinum sheet ð20 £ 20 mm2 Þ: For silver plating, a silver sheet was used as an anode. † Inter-electrode distance: , 50 mm. † Agitation: no agitation (still bath) or agitation with a magnetic stirrer. † Plating temperature: controlled within a ^1 8C accuracy of the set temperature. 7.2.2 Substrate materials A flat area of 20 £ 20 mm2 on one side of the sheet/foil was left exposed for subsequent plating by masking the remaining area. By doing this, the current density of the exposed area is always kept constant. 7.2.2.1. Polycrystalline copper sheets. For the removal of internal stress and to reduce any lattice imperfections left in the substrate, annealing for 1 h at 600 8C was performed. In addition, the substrate was electropolished in a solution containing 2-parts water and 1-part phosphoric acid for 15 min, firstly to remove a deformation zone induced by rolling and secondly to obtain a mirror-finish surface. 7.2.2.2. Single-crystal copper ({100}, {110}, and {111} planes). A single-crystal copper rod 23 mm in diameter and 300 mm in length was grown in a Bridgeman crystal growth furnace. After the orientation was determined from a Laue photograph, a thin slice of variously oriented (100, 110, and 111) copper was cut with a diamond-cutting wheel. To remove the deformation zone left on the surface after cutting, the sliced surface was further electropolished until a mirror finish was obtained. 7.2.2.3. Stainless steel sheets (SUS-304). The surface of stainless steel is generally covered with a passive amorphous film. On this substrate, it is possible to plate films that are not affected by the crystal structure of the substrate. Another advantage is that the plated films do not stick well on these substrates and thus can be easily peeled off. For

Database for the Microstructure of Plated Films

259

surface treatment prior to plating, the substrate surface was first degreased in acetone, then washed in distilled water, followed by soaking in a dilute sulfuric acid (10%) solution, and finally washed in distilled water. Plating was conducted immediately after the final wash while the substrate was still wet. 7.2.2.4. Electrolytic Ni –25 at.% P alloy films. Electrolytic Ni –25 at.% P alloy films are amorphous. Similar to the case of the stainless steel substrates, films plated on these substrates take an intrinsic structural form that is not affected by the crystal structure of the substrate. Therefore, the use of amorphous substrates is convenient in studying the crystal size or texture of crystalline films grown under various plating conditions. The same surface treatment used on the stainless steel sheets was used for preparing this substrate surface. 7.2.2.5. Splat-quenched amorphous alloy foils (Fe –Si –B or Fe –Co – Si – B alloys manufactured by Nippon Amorphous Metals Co., Ltd.). This substrate was used for the same purpose as the stainless steel sheets or electrolytic Ni – 25 at.% P alloy films. Again, the surface treatment follows the same procedure as the stainless steel sheets and electrolytic Ni – 25 at.% P films.

7.3. MICROSTRUCTURE OBSERVATIONS AND MEASUREMENTS METHODS FOR PLATED FILMS

Microstructural observations and measurement methods for plated films follow the procedure described in Chapter 6. It was shown that either the orientation (, hkl . ) or the plane ({hkl}) can be used to describe the texture (preferred orientation) of the film. For convenience, however, Miller indices without any parentheses, i.e. hkl, were used for expressing the film texture.

7.4 Database for the Microstructure of Plated Films 7.4.1. Pure Metals 7.4.1.1. Electrolytic Ag (Watanabe et al., 2002) 7.4.1.1.1. SILVER METAL

Silver metal is face-centered cubic (fcc) and the lattice constant, a, is 0.4086 nm. The melting point is 960.80 8C and the resistivity is 1:59 £ 1026 V=cm:

7.4.1.1.2. PLATING METHOD

A silver iodide (AgI) bath was used to plate Ag films. The bath was prepared by adding potassium iodide (KI), which can form a silver complex because silver iodide alone does not dissolve easily in water. The composition of the bath was 0.05 mol/l AgI and 2.0 mol/l KI. A silver sheet was used as an anode and plated amorphous Ni– P films were substrate materials. The bath was agitated using a magnetic stirrer during plating. The film thickness was determined from the measured film weight and estimated film density. The current efficiency was almost 100% under all plating conditions used. The effect of current density, bath temperature, film thickness, and pH on the film microstructure was investigated in this study.

7.4.1.1.3. FILM MICROSTRUCTURE

7.4.1.1.3.1 Crystal size and surface morphology Figure 7.4.1.1.1(A) shows a change in surface morphology with increasing film thickness in electrolytic Ag films grown at 150 A/m2 from the 60 8C bath (pH ¼ 2:0; adjusted using sulfuric acid). When the film was thin (, 0.25 mm), granular-shaped Ag crystals (, 0.5 mm in diameter) formed randomly over the amorphous substrate and grew larger with time (or with increasing thickness). The density of these fine crystalline Ag particles was initially 120 particles/100 mm2, but decreased to 2 particles/100 mm2 when the film reached 15 mm in thickness. This growth morphology is illustrated schematically in Figure 7.4.1.1.1(B). The surface morphology change with the current density is compared for the 10 mmthick films (see Figure 7.4.1.1.2). The grain size is large at low current densities but becomes smaller at higher current densities. This trend can be explained as follows. It was 260

Electrolytic Ag Figure 7.4.1.1.1. A change in the surface morphology of electrolytic silver films with increasing film thickness. The current density ¼ 150 A=m2 ; the solution temperature ¼ 60 8C; pH ¼ 2 (adjusted with H2SO4), and the substrate ¼ amorphous:

261

262 Nano-Plating Figure 7.4.1.1.2. A change in the surface morphology of electrolytic silver films with increasing current density. The film thickness ¼ 10 mm; the solution temperature ¼ 60 8C; pH ¼ 2 (adjusted with H2SO4), and the substrate ¼ amorphous:

Electrolytic Ag

263

discussed previously that metal atoms deposit through their discharge processes, which do not occur uniformly over the cathode surface but take place preferentially at protrusions. The density of such protrusions is low at low current densities and is high at high current densities. For low melting point metals, the size of the protrusions corresponds roughly to the grain size. The grain size, therefore, is large at low current densities and small at high current densities. Figure 7.4.1.1.3 shows the change in surface morphology with increasing bath temperature in 10 mm-thick electrolytic Ag films grown at 150 A/m2 from the Ag bath, whose pH was 2.0. It is clear that the grain size is small at low bath temperature but increases with increasing bath temperature. As described above, the surface smoothness depends primarily on the distribution of preferential metal ion discharge sites such as protrusions. In the present case, however, the grain size increase is due to an increase in the diffusivity of metal ions in the bath or due to an increase in the surface diffusivity of deposited adatoms. Figure 7.4.1.1.4(a) demonstrates how surface morphology changes occur with a change in the bath pH for 10 mm-thick electrolytic Ag films grown at 150 A/m2 in a 60 8C bath. The bath pH less than 6 was adjusted with H2SO4, whereas any pH larger than 9 was changed with KOH. The surface morphology does not appear to be effected by a change in the bath pH. In Figure 7.4.1.1.5, we show the surface morphology change with the type of buffer (the type of anion) used. The bath pH was adjusted to 2.0 using H2SO4, HNO3, and HCl. The film thickness was 10 mm. Regardless of the type of anions used, the grain size is large at low current densities and decreases slightly at high current densities. This grain size behavior with current density is consistent with the explanation given above. The effect of the type of anions used on the surface morphology is considered small here. Figure 7.4.1.1.6 is a surface morphology map for 10 mm-thick electrolytic Ag films, whose micrographs are arranged in terms of current density, bath temperature and texture. The grain size is large for the films plated at low current densities and in the hightemperature bath, whereas it is small at high current densities and in the low-temperature bath. This trend is again consistent with the above result. 7.4.1.1.3.2 Preferred orientation (texture) Texture index changes in electrolytic Ag films are plotted as a function of film thickness and current density. Although amorphous Ni –P films were used as the substrates, the Ag films initially exhibit a weak 111 texture, but the texture changed to 110 with increasing film thickness. The critical thickness, at which the texture changes from 111 to 110, is relatively thick at low current densities. For example, 5 mm at 50 A/m2, but is thin at high current densities (less than 1 mm at 250 A/m2). At 300 A/m2, the 110 texture starts from the beginning. The 110 texture is weak at low current densities but becomes stronger with

264 Nano-Plating Fig. 7.4.1.1.3. A change in the surface morphology of electrolytic silver films with increasing solution temperature. The current density ¼ 150 A=m2 ; pH ¼ 2; the film thickness ¼ 10 mm; and the substrate ¼ amorphous:

Electrolytic Ag

265

Figure 7.4.1.1.4. A change in (a) the surface morphology and (b) the texture of electrolytic silver films as a function of pH. The film thickness ¼ 10 mm; the current density ¼ 150 A=m2 ; the solution temperature ¼ 60 8C; and the substrate ¼ amorphous:

266 Nano-Plating Figure 7.4.1.1.5. A change in the surface morphology of electrolytic silver films as a function of current density and the type of a buffer used. (a) H2SO4, (b) HNO3, and (c) HCl. The film thickness ¼ 10 mm; the solution temperature ¼ 60 8C; and the substrate ¼ amorphous:

Electrolytic Ag

267

Figure 7.4.1.1.6. A change in the texture of electrolytic silver films as a function of current density and solution temperature. The film thickness ¼ 10 mm and the substrate ¼ amorphous:

increasing current densities. At the same time, the 111-textured grains start appearing at current densities above 250 A/m2 (Figure 7.4.1.1.7). A contour map of the texture index for electrolytic silver films is drawn in Figure 7.4.1.1.8 as a function of film thickness and current density. From this map, it is clear that the texture is not strong at low current densities. At 150 A/m2, however, the 110 texture emerges and becomes stronger with increasing film thickness. With a further increase in current density and film thickness, the texture changed to 111 for some films. This thickness range is enclosed with a dotted rectangular box in Figure 7.4.1.1.8. For example, the 10 mm-thick film started taking the 111 texture at current densities above 250 A/m2, which corresponds to the appearance of fine grains as seen in Figure 7.4.1.1.2. Based on this observation, we can conclude that the large grains obtained

268 Nano-Plating Figure 7.4.1.1.7. A change in the texture of electrolytic silver films as a function of film thickness and current density. The solution temperature ¼ 60 8C; pH ¼ 2 (H2SO4 as a buffer), and the substrate ¼ amorphous:

Electrolytic Ag

269

Figure 7.4.1.1.8. A contour map of the texture index for electrolytic silver films as a function of film thickness and current density.

Figure 7.4.1.1.9. Relationship between the texture index and solution temperature for electrolytic silver films. The current density ¼ 150 A=m2 ; pH ¼ 2 (H2SO4 as a buffer), and the substrate ¼ amorphous:

270

Nano-Plating

Figure 7.4.1.1.10. A change in the texture of electrolytic silver films as a function of current density and the type of a buffer used. The pH was adjusted to 2 using three types of adjuster: (a) H2SO4, (b) HNO3, and (c) HCl. The film thickness ¼ 10 mm; the solution temperature ¼ 60 8C; and the substrate ¼ amorphous:

Electrolytic Ag Figure 7.4.1.1.11. A change in the surface morphology of electrolytic silver films grown on (a) amorphous substrates and (b) copper foils as a function of film thickness. The current density ¼ 15 mA=cm2 ; the solution temperature ¼ 60 8C; and pH ¼ 2 (H2SO4 as an adjuster).

271

272 Nano-Plating Figure 7.4.1.1.12. A change in the surface morphology of electrolytic silver films grown on (a) amorphous substrates and (b) copper foils as a function of current density. The film thickness ¼ 10 mm; the solution temperature ¼ 60 8C; and pH ¼ 2 (H2SO4 as an adjuster).

Electrolytic Ag

273

at low current densities are 110-oriented, whereas the small grains at high current densities are 111-oriented. The texture index of electrolytic Ag films is plotted against the bath temperature in Figure 7.4.1.1.9. The films exhibited a weak 111 texture for low bath temperatures (25 8C) and displayed the 110 texture at 30 8C. The 110 texture index linearly increased up to 55 8C, but became discontinuous at 60 8C. Beyond 60 8C, however, the 110 texture index did not change. The surface morphology did not show any changes at 60 8C, as seen in Figure 7.4.1.1.3. It is important to remember that the film retained the 110 texture regardless of the presence of the discontinuity. Figure 7.4.1.1.4(B) shows the texture index change with bath pH. According to this graph, the 111 texture index appears to increase slightly with increasing pH (or increasing hydrogen ion concentration). In Figure 7.4.1.1.10, three types of buffers, i.e. (a) H2SO2, (b) HNO3, and (c) HCl were used to adjust the pH to 2 and the texture index change was investigated as a function of current density using these three baths. The texture index changes with current density and exhibits a peak at some current density for all cases. The current density, at which the texture index displays a peak, depends on the type of the buffer used; , 200 A/m2 for (a) H2SO2, , 210 A/m2 for (a) HNO3, and , 250 A/m2 for (a) HCl. Based on the above results, it can be concluded that the texture index change is not due to a change in bath pH (or hydrogen ion concentration), but is due to the type of anions dissociated from a buffer. Consequently, it is reasonable to assume that the amount of a buffer used will affect the texture index. Various surface morphologies of electrolytic Ag films are displayed according to current density, bath temperature, and texture in Figure 7.4.1.1.6. The 111 texture is prevalent at high current densities in the low-temperature baths, whereas the 110 texture is dominant at low current densities in the high-temperature baths. In addition, we note that for the 110-textured films, the grain size was large, whereas for the 111-textured films, not only did the grain size become small, but dendritic crystals also appeared. A surface morphology change with increasing film thickness in electrolytic Ag films is shown in Figure 7.4.1.1.11. Although two markedly different substrates, i.e. (a) an amorphous substrate and (b) copper foil, were used in this study, the surface morphology indicated no significant difference. A surface morphology change with a change in the current density is shown in Figure 7.4.1.1.12. Again, two different substrates ((a) an amorphous substrate and (b) copper foil) were used to test the effect of substrates. Very little substrate effect on the surface morphology was seen.

REFERENCE Watanabe, T. & Osada, M. (2002) J. Jpn Inst. Metals, 66, 614.

7.4.1.2. Electrolytic Au 7.4.1.2.1. GOLD METAL

The structure of gold metal is fcc and the lattice constant, a, is 0.4079 nm. The melting point is 1063.0 8C and the resistivity is 2:35 £ 1026 V=cm:


Chloride-based Au baths were prepared in order to study the early nucleation and growth stages of gold electrodeposition (see Table 7.4.1.2.1) (Tanabe and Kamasaki, 1970, 1971, 1972; Kamasaki and Tanabe, 1973, 1974a –g, 1975; Kamasaki, 1974). The bath pH was adjusted with HCl, K2CO3, and KCN. Rolled polycrystalline pure Fe and single-crystal {001} Fe sheets were electropolished to a mirror finish, and used as the substrates. For surface morphology and film texture studies, cyanide and sulfite Au plating baths were used (see Table 7.4.1.2.2) (Inoue et al., 2002). In this study, the substrate was prepared by electropolishing a polycrystalline pure Cu sheet, on which a 2 mm-thick amorphous Ni-25 at.% P film was grown by an electroless method. If the electroless Ni– P film is immersed in the Au plating solutions, it will dissolve in the solution, followed by deposition of Au by a displacement mechanism. To avoid such deposition, a strike Au film was grown on the substrate by an electrolytic plating method. The Au-covered substrates were then used to plate Au films in the two baths under various plating conditions (see Table 7.4.1.2.2). The strike Au bath is basically a cyanide bath with a low concentration of Au salt. The strike Au plating is generally performed at high current densities for a short time. The structure of strike Au films will be described in Figure 7.4.1.2.11.

7.4.1.2.3. INITIAL STAGES OF GOLD ELECTRODEPOSITION

Figure 7.4.1.2.1 plots a change in the average film thickness of electrolytic Au films with time. These Au films were deposited from four types of baths ((A), (B), (G), and (H)), which were prepared by varying the concentration of HAuCl4 and by adjusting the bath pH with HCl (see Table 7.4.1.2.1). The film thickness was measured using an optical interferometer. The slope of the plots in Figure 7.4.1.2.1 gives the deposition rate, which depends on the type of bath used. If the Au concentration in the bath is high, the deposition rate is high. If the Au concentration is the same, the deposition rate is higher for the bath containing more HCl. 274

275

Electrolytic Au Table 7.4.1.2.1. Bath compositions used for plating gold films. Bath name Composition

A B C D E F G H I

HAuCl4 (wt%)

K2CO3 (wt%)

K2CO3 (wt%)

HCl KCN (Vol) (wt%)

0.15 0.15 0.15 0.15 0.15 0.15 0.30 0.30 0.01

– – – – Added Added – – –

– – – – – – – – 0.06

– 0.8 2.0 20.0 – – – – 0.8

– – – – – – – – 0.001

pH

2.2 1.1 0.7 p 0.3 7.0 10.0 1.9 1.0 0.8

Temperature Voltage or current density

AT AT AT AT AT AT AT AT 65 8C

3V 3V 3V 3V 3V 3V 3V 3V 0.5 A/dm2

Vol ¼ ml/100 ml HAuCl4; AT ¼ ambient temperature.

Figure 7.4.1.2.2 displays plan-view TEM micrographs showing the initial nucleation and growth stages of electrolytic Au deposited on polycrystalline Fe substrates. Both the Au deposit and the substrate can be seen simultaneously in these micrographs. Micrographs in Figure 7.4.1.2.2(a) – (c) were taken from the three different areas of the same Au/substrate composite specimen, which was plated for 0.5 s. It is clear that the distribution of the Au nuclei ranging from 5 to 50 nm in size is random and changes markedly over the surface. If the Au was nucleated at the steps or kink sites of the substrate surface, the nuclei distribution is expected to be more orderly. The electron diffraction pattern from Figure 7.4.1.2.2(a) exhibited rings, further indicating that the nuclei were indeed distributed randomly. Here we can imagine that the Au has nucleated randomly without seeking stable sites such as kinks and steps on the substrate. The electron diffraction patterns from Figure 7.4.1.2.2(b) and (c) contained both rings and single-crystal spots. In addition, the density of the Au nuclei in Figure 7.4.1.2.2(a) is much lower than those in Figure 7.4.1.2.2(b) and (c). We found that the difference in the Au nuclei distribution arose from the different grain orientation of the substrate. It is important to remember that the substrate was polycrystalline. These results can be interpreted in the following manner. When the nuclei are small, they will not lattice-match to the substrate because of their high surface mobility. As the nuclei grow larger, however, they lose the Table 7.4.1.2.2. Two plating baths used to study the surface morphology and texture of gold films.

Amount pH

(a) Cyanide bath; KAu(CN)2

(b) Sulfite bath; Na3Au(SO3)2

10 g/l (0.05 mol/l) 6.2

10 g/l (0.05 mol/l) 8.2

276

Nano-Plating

Figure 7.4.1.2.1. Relationship between plating time and average film thickness for gold electrodeposits grown on polycrystalline iron substrates.

mobility and tend to adjust their positions to the state of a lower energy configuration by matching with the lattice of the substrate. Figure 7.4.1.2.2(d) is a continuation of Figure 7.4.1.2.2(a) – (c). In Figure 7.4.1.2.2(d), we show a thicker Au film, which was electroplated for 4 s. Here the Au film is seen to cover the whole substrate surface. In the electron diffraction pattern, both diffraction rings and the , 100 . zone axis single-crystal diffraction spots appeared. It is clear from this result that with increasing film thickness, Au grains with new orientations different from that of the lattice-matched grains were nucleated. Gold deposits after 0.5-s plating from the A, C, and D baths are seen in Figure 7.4.1.2.3(a) – (c), respectively. The substrates used were polycrystalline Fe sheets and the bath pH, which was adjusted to 2.2, 0.7, and p 0.3 with HCl for the A, C, and D baths, respectively. Similar to the case of the Au baths used in Figure 7.4.1.2.1, the deposition rate was high for the low-pH bath. Again, the difference seen in the Au nuclei size/distribution and their electron diffractions was found to be a result of the difference in the orientation of the substrate grains. Figure 7.4.1.2.4 displays plan-view TEM micrographs showing a time sequence of the initial nucleation and growth stages of electrolytic Au deposited on {100} single-crystal Fe substrates from the E bath (a gold-chloride-type bath). The bath pH was adjusted to 7 with the addition of K 2 CO 3. The deposition time was 0.5, 4, 8, and 20 s for Figure 7.4.1.2.4(a) –(d), respectively. The maximum size of the initial Au nuclei seen in

Electrolytic Au

277

Figure 7.4.1.2.2. Plan-view TEM micrographs showing the film morphology of gold films plated from the A bath on polycrystalline iron substrates. (a) 0.5, (b) 0.5, (c) 0.5, and (d) 4 s plating time.

278

Nano-Plating

Figure 7.4.1.2.3. Plan-view TEM micrographs showing the difference in the deposition rate of gold electrodeposits. The plating was conducted on polycrystalline iron substrates for 0.5 s in (a) the A bath ðpH ¼ 2:2Þ; (b) the C bath ðpH ¼ 0:7Þ; and (c) the D bath ðpH p 0:3Þ:

Electrolytic Au

279

Figure 7.4.1.2.4. A change in the film morphology of gold electrodeposits grown on the {100} surface of iron substrates from the E bath, whose pH was adjusted to 7 using K2CO3 as a buffer. The plating time was (a) 0.5, (b) 4, (c) 8, and (d) 20 s.

280 Nano-Plating Figure 7.4.1.2.5. A change in the film morphology of gold electrodeposits grown on the {100} surface of iron substrates from the E bath, whose pH was adjusted to 10 using K2CO3 as a buffer. The plating time was (a) 0.5, (b) 1, (c) 2, and (d) 3 s.

Electrolytic Au

281

Figure 7.4.1.2.6. A plan-view TEM micrograph and its electron diffraction pattern showing the structure of a gold film grown for 70 s from the I bath, whose pH was adjusted to 0.8 using KCN as a buffer.

Figure 7.4.1.2.4(a) was 20 nm, although these nuclei grew larger with time. After the 4-s plating, the electron diffraction pattern consisted of rings and thus the Au film was not lattice-matched to the substrate (see Figure 7.4.1.2.4(b)). After the 8-s plating (cf. Figure 7.4.1.2.4(c)), the , 100 . zone axis diffraction spots appeared in addition to the rings, indicating that the film had started lattice-matching to the substrate. After the 20-s plating, the Au film covered most of the substrate surface (see Figure 7.4.1.2.4(d)). The electron diffraction pattern indicated that the Au film consisted of lattice-matched as well as non-matched grains. Additional time-sequence TEM micrographs are shown in Figure 7.4.1.2.5, where the Au was plated in the F bath. Similar to the E bath, the bath pH was adjusted to 10 with K2CO3. The morphological change is almost the same as the E bath, except for the shorter plating time, i.e. (a) 0.5, (b) 1, (c) 2, and (d) 3 s. The deposition rate from the F bath is apparently much larger. As seen in Figure 7.4.1.2.4(b), even after the 1-s plating, the electron diffraction pattern is already exhibiting the , 100 . zone axis spots, indicating the lattice-matching of the Au to the substrate. The degree of the Au lattice-matching increased with increasing plating time because the intensity of the diffraction spots became stronger, while the rings became weaker (see Figure 7.4.1.2.4(c) and (d)). A bright-field (BF) TEM micrograph of a thick Au film grown for 70 s from the I bath is shown in Figure 7.4.1.2.6. The bath pH was adjusted to 0.8 with KCN. Although the BF image appears to be a polycrystalline film, the corresponding electron diffraction pattern still contains single-crystal spots superimposed with an arc. This arching is a strong

282

Nano-Plating

indication that new grains with the orientations slightly rotated around the , 100 . zone axis are nucleated. These rotated grains are bounded by low-angle sub-boundaries, which apparently gave the impression that this Au film was polycrystalline. The structure of a thin Au film grown from the I bath is shown in Figure 7.4.1.2.7. The Au film was essentially a {100} single-crystal film containing a high density of growth twins (see Figure 7.4.1.2.7(a)). The presence of growth twins is indicated by the twinrelated spots in the electron diffraction spots (Figure 7.4.1.2.7(b)) and the twin images were revealed using the standard dark-field imaging as seen in Figure 7.4.1.2.7(c). These growth twins are seen to align along the two , 110 . directions. A 72-nm Au film obtained from the I bath was stripped off the substrate and subsequently heated at 250 8C for 1 h inside a microscope. Figure 7.4.1.2.8(a) and (b) represent the film before and after the heating, respectively. The locations marked with A and A0 are an identical site in the film. The heating appears to have eliminated certain defects such as dislocations. The rotated grains, indicated by arching in the diffraction spots, have grown as a result of the heating. However, the sub-boundaries surrounding the lattice-matched and rotated grains remained. Figure 7.4.1.2.9 illustrates how the deformed structure of cold-rolled {100} Fe substrates affects the early nucleation and growth stages of the Au deposits obtained from the I bath. The Fe substrates used in this experiment were cold-rolled with the thickness being reduced to 50% of their original values. The plating times were (a) 0.5, (b) 2, (c) 15, and (d) 30 s. These results are very similar to those described in Figure 7.4.1.2.4, where the {100} single-crystal Fe substrates were used directly without any cold-rolling. For the 0.5s plating, the electron diffraction pattern consists of only rings, indicating that the Au deposit is randomly oriented. Au deposits plated for more than 2 s exhibited a strong single-crystal diffraction pattern. The deformation structure induced by cold-rolling of the substrate did not affect the lattice-matching of the Au deposits. Gold films were deposited on tempered Fe– 0.8 wt.% C steel sheets from the I bath. The tempered steel substrate contained martensite grains. TEM micrographs and electron diffraction patterns for the Au films deposited for (a) 2, (b) 4, (b0 ) 4, and (c) 20 s are seen in Figure 7.4.1.2.10. Initially, fine Au grains appear to decorate the underlying acicular martensite grains (see Figure 7.4.1.2.10 (a) and (b)). Such a surface morphology began to disappear after 20 s plating. Nevertheless, the Au films are expected to have grown latticematched to the underlying martensite grains.

7.4.1.2.4. THE SURFACE MORPHOLOGY AND TEXTURE OF PLATED GOLD FILMS

Figure 7.4.1.2.11 shows TEM micrographs for an electrolytic Au strike film grown on a plated amorphous Ni– P film. The film is seen to cover the substrate uniformly and its

Electrolytic Au Figure 7.4.1.2.7. The structure of a {100} gold film obtained from the I bath. A site marked as A in (a) the bright-field (BF) image corresponds to that marked as A0 in (c) the dark-field image. The corresponding electron diffraction pattern is also shown in (b).

283

284

Nano-Plating

Figure 7.4.1.2.8. A structural change in gold electrodeposits observed after annealing at 280 8C for 1 h. (a) Before and (b) after the annealing. Symbols A and A0 indicate an identical position.

Figure 7.4.1.2.9. A structural change in electrolytic gold films grown from the I bath on {100} iron substrates (rolled by 50% in thickness reduction) with increasing plating time. (a) 0.5, (b) 2, (c) 15, and (d) 30 s plating time.

Electrolytic Au

285

Figure 7.4.1.2.10. The structure of electrolytic gold films grown on annealed steel (martensite) substrates. The BF images and its corresponding electron diffraction patterns are shown for gold films plated for (a) 2, (b) 4, and (c) 20 s.

286

Nano-Plating

Figure 7.4.1.2.11. The structure of a gold strike film plated on an amorphous Ni–25 at.% P substrate. The BF image and its electron diffraction pattern are shown in (a) and (b), respectively, whereas the BF image in (a0 ) is taken at higher magnification.

average grain size was 20 nm. The electron diffraction pattern consisted of complete rings and thus no texture was present. According to X-ray analysis, however, the film had a weak 111 texture with a texture index of 1.2. Figures 7.4.1.2.12.1– 7.4.1.2.12.3 illustrate the effect of bath temperature (40, 60, and 80 8C), film thickness, and current density on the surface morphology of Au films plated in the cyanide bath (cf. Table 7.4.1.2.2). For all bath temperatures and current densities used here, the film surface became progressively rougher with increasing film thickness. The change in the surface roughness did not appear to be affected by current density or bath temperature. Each surface irregularity has a well-defined geometrical form and thus is most likely to consist of single-crystal Au grains. The textures of these Au films were further studied using an X-ray diffraction method (see Figure 7.4.1.2.13). One common feature, illustrated in Figure 7.4.1.2.13 is that no texture initially developed in thin (0.5 mm) Au films, but that the 111 texture appeared suddenly at the film thickness of

Electrolytic Au

287

Figure 7.4.1.2.12.1. SEM micrographs showing the surface morphology of gold films plated on amorphous substrates from a cyanide bath, whose solution temperature was 40 8C. These micrographs are arranged as a function of film thickness and current density.

0.7 mm. The 111 texture was maintained even when the film thickness increased further. The use of an amorphous substrate was responsible for the absence of texture in the thin films. Figures 7.4.1.2.14.1 – 7.4.1.2.14.3 show the effect of bath temperature (40, 60, and 80 8C), film thickness, and current density on the surface morphology of Au films plated from the sulfite bath (cf. Table 7.4.1.2.2). Except for the 40 8C bath, the film surface

288

Nano-Plating


Electrolytic Au

289


290 Nano-Plating Figure 7.4.1.2.13. A change in the texture index of gold films plated on amorphous substrates from a cyanide bath. These plots are arranged as a function of film thickness, solution temperature, and current density.

Electrolytic Au

291

Figure 7.4.1.2.14.1. A surface morphology change in gold films grown on amorphous substrates from a sulfite bath (408C). These SEM micrographs are arranged as a function of film thickness and current density.

292

Nano-Plating

Figure 7.4.1.2.14.2. A surface morphology change in gold films grown on amorphous substrates from a sulfite bath (608C). The SEM micrographs are arranged as a function of film thickness and current density.

Electrolytic Au

293

Figure 7.4.1.2.14.3. A surface morphology change in gold films grown on amorphous substrates from a sulfite bath (80 8C). The SEM micrographs are arranged as a function of film thickness and current density.

294 Nano-Plating Figure 7.4.1.2.15. A change in the texture index of gold films plated on amorphous substrates from a sulfite bath. These plots are arranged as a function of film thickness, solution temperature, and current density.

Electrolytic Au

295

gradually became rougher with increasing film thickness. In the 40 8C bath, the surface roughness increased with increasing film thickness, but the morphology exhibited a complex change. At a current density of 20 A/m2, semi-spherically shaped features suddenly appeared on the surface at a film thickness of 50 mm. At 40 A/m2, the surface roughness increased uniformly with increasing film thickness. Finally, at 80 A/m2, dendritic Au crystals formed on the surface at a thickness of 1.0 mm. For the 60 8C (Figure 7.4.1.2.14.2) and 80 8C baths (Figure 7.4.1.2.14.3), no anomalous growth was observed as the surface roughness increased smoothly with film thickness. The surface roughness of the Au films from the cyanide bath consisted of sharply pointed protrusions, whereas the sulfite bath caused rounded protrusions. For both 60 and 80 8C baths, the surface roughness was very fine for the films grown at low current densities but became coarse for those obtained at high current densities. Figure 7.4.1.2.15 summarizes the texture of the Au films studied above. One common feature in these graphs is that, similar to the case of the cyanide bath, the film did not possess any texture initially but exhibited the 111 texture temporarily at the thickness of 0.7 mm, followed by the development of various textures. The final texture depends on plating conditions, but it is difficult to deduce a systematic trend from these graphs. As described previously, the Au films obtained from the 40 8C bath (cf. Figure 7.4.1.2.14.1) displayed very different surface morphologies as their textures changed with film thickness. Conversely, the surface roughness of the Au films from the 60 and 80 8C baths increased smoothly with increasing film thickness. The surface morphologies of the Au films from the 60 and 80 8C baths were very similar; however, the behavior of the texture was very different.

REFERENCES Inoue, K., Nakata, T. & Watanabe, T. (2002) J. Jpn Inst. Metals, 66, 400. Kamasaki, S. (1974) A TEM study on the crystal growth & fine structure of electrolytic gold films, PhD Dissertation, Tokyo Metropolitan University. Kamasaki, S. & Tanabe, Y. (1973) J. Metal Finish. Soc. Jpn, 24, 276. Kamasaki, S. & Tanabe, Y. (1974a) J. Metal Finish. Soc. Jpn, 25, 275. Kamasaki, S. & Tanabe, Y. (1974b) J. Metal Finish. Soc. Jpn, 25, 476. Kamasaki, S. & Tanabe, Y. (1974c) J. Metal Finish. Soc. Jpn, 25, 482. Kamasaki, S. & Tanabe, Y. (1974d) J. Metal Finish. Soc. Jpn, 25, 528. Kamasaki, S. & Tanabe, Y. (1974e) J. Metal Finish. Soc. Jpn, 25, 588. Kamasaki, S. & Tanabe, Y. (1974f) J. Metal Finish. Soc. Jpn, 25, 647. Kamasaki, S. & Tanabe, Y. (1974g) J. Metal Finish. Soc. Jpn, 25, 653. Kamasaki, S. & Tanabe, Y. (1975) J. Metal Finish. Soc. Jpn, 26, 15. Tanabe, Y. & Kamasaki, S. (1970) J. Metal Finish. Soc. Jpn, 21, 281. Tanabe, Y. & Kamasaki, S. (1971) J. Metal Finish. Soc. Jpn, 22, 54. Tanabe, Y. & Kamasaki, S. (1972) J. Metal Finish. Soc. Jpn, 23, 572.

7.4.1.3. Electrolytic Cd (Miyazaki, 1998) 7.4.1.3.1. CADMIUM METAL

The structure of cadmium metal is hexagonal close-packed (hcp). The lattice constants are a ¼ 0:2987 nm and c ¼ 0:5617 nm ðc=a ¼ 1:886Þ: If the atoms forming the hcp structure are assumed to be spherical, the ratio of c=a will be 1.663. Compared with this ideal ratio (1.663), the hcp structure of cadmium has a larger c=a value, thus it is envisioned that the cadmium lattice is stretched along the , 001 . direction. The melting point is 320.9 8C and the resistivity is 7:3 £ 1026 V=cm:


Three types of baths were prepared for plating Cd films (see Table 7.4.1.3.1). The current density, bath temperature, film thickness, and bath pH were varied in this experiment. Amorphous Fe –Co – Si – B foils (manufactured by Nippon Amorphous Metals Co., Ltd.) obtained by a splat-quenching method were used as the substrates. The film thickness, which ranged from 0.17 to 25.2 mm, was determined by weighing the plated films. The bath was agitated with a magnetic stirrer during plating. Figure 7.4.1.3.1 plots the current efficiency change with current densities in electrolytic Cd films grown from (A) sulfate and (B) chloride baths. For both baths, the current efficiency is high at low current densities but decreases at high current densities.


Figure 7.4.1.3.2 shows a surface morphology change with increasing film thickness in Cd films obtained from (A) sulfate and (B) chloride baths. In both baths, isolated granular Cd crystals with a dimension of several mm were initially formed on the substrate surface. With increasing film thickness these crystals grew and linked together to cover the whole surface. The change in texture index with increasing film thickness in the Cd films, described in Figure 7.4.1.3.2, is plotted in Figure 7.4.1.3.3. Since the amorphous foils were used as substrates, the initial Cd nuclei took random orientations and consequently the thin Cd films did not display any texture. With increasing film thickness, the film rapidly took the 001 texture. 296

297

Electrolytic Cd Table 7.4.1.3.1. Bath compositions for plating cadmium films. Bath name

Bath composition

(A) Sulfate bath (B) Chloride bath (C) Nitrate bath

CdSO4·8/3H2O (1 mol/l) CdCl2·5/2H2O (1 mol/l) Cd(NO3)2·4H2O (1 mol/l)

The surface morphology of Cd films obtained from (A) sulfate and (B) chloride baths were studied by varying the current density while keeping the total consumed electric charge constant (see Figure 7.4.1.3.4). In both baths, at low current densities, the surface morphology can be characterized as large scattered Cd crystals covering the substrate surface. At high current densities, the whole surface was uniformly covered with fine Cd crystals. The surface morphology of Cd films from a nitric bath is seen in Figure 7.4.1.3.5 and is quite different from those obtained from the sulfate or chloride baths. A texture index change with increasing current density in Cd films from (A) sulfate and (B) chloride baths is shown in Figure 7.4.1.3.6. For both baths, the films did not exhibit any texture at low current densities but showed a strong 001 texture at high current densities. The 001 texture is equivalent to many hexagonal Cd crystals standing on their basal plane.

Figure 7.4.1.3.1. Relationship between current density and current efficiency for electrodeposited cadmium films obtained from (A) sulfate and (B) chloride baths. The film thickness ¼ 6:7 mm; pH ¼ 1:8; the solution temperature ¼ 25 8C; and the substrate ¼ amorphous:

298

Nano-Plating

Figure 7.4.1.3.2. A surface morphology change in electrolytic cadmium films as a function of film thickness. The cadmium films were obtained from (A) sulfate and (B) chloride baths. The current density ¼ 1000 A=m2 ; pH ¼ 1:8; the solution temperature ¼ 25 8C; and the substrate ¼ amorphous:

Electrolytic Cd

299

Figure 7.4.1.3.3. A change in the texture index of electrolytic cadmium films with film thickness. The cadmium films were obtained from (A) sulfate and (B) chloride baths. The pH ¼ 1:8; the solution temperature ¼ 25 8C, and the substrate ¼ amorphous:

A surface morphology change with bath temperature in Cd films plated at 1000 A/m2 from (A) sulfate and (B) chloride baths ðpH ¼ 1:8Þ is shown in Figure 7.4.1.3.7. In both cases, the Cd grain size is larger in the high-temperature baths than that from the lowtemperature baths. Figure 7.4.1.3.8 shows a texture index change with bath temperature in 6.73 mm-thick Cd films deposited at 1000 A/m2 from (A) sulfate and (B) chloride baths ðpH ¼ 1:8Þ: In the sulfate bath, the texture did not change with bath temperature but remained as 001 (Figure 7.4.1.3.9). If the plating was performed at 750 A/m2, the texture was 001 at low bath temperatures, but changed to 103 at a bath temperature of 40 8C (see Figure 7.4.1.3.10). In the chloride bath, even at the current density of 1000 A/m2, the texture changed from 001 to 103 at 40 8C. A surface morphology change with a change of bath pH in Cd films plated from sulfate baths is shown in Figure 7.4.1.3.10. It is clear that the crystal size did not change significantly with the bath pH. The texture of these crystals, however, changed from 001 to 103 with increasing pH as seen in Figure 7.4.1.3.11. Figure 7.4.1.3.12 displays SEM micrographs showing the surface morphology of electrolytic Cd films grown on the {100), {110}, and {111} surfaces of a single-crystal copper. The Cd crystals grown on the {100} and {111} surfaces are elongated, whereas those deposited on the {110} surface are rounded and granular. Nevertheless, these crystals are expected to be lattice-matched to the substrate.

300

Nano-Plating

Figure 7.4.1.3.4. A surface morphology change with current density in electrolytic cadmium films grown on amorphous substrates from (A) sulfate and (B) chloride baths. The film thickness ¼ 6:7 mm; pH ¼ 1:8; and the solution temperature ¼ 25 8C:

Electrolytic Cd

301

Figure 7.4.1.3.5. Surface morphologies in electrolytic cadmium films plated on amorphous substrates at current densities of 200 and 1000 A/m2 from a nitric acid bath.

Figure 7.4.1.3.6. A texture index change with current density in electrolytic cadmium films grown amorphous substrates from (A) sulfate and (B) chloride baths. The film thickness ¼ 6:7 mm; pH ¼ 1:8; and the solution temperature ¼ 25 8C:

302

Nano-Plating

Figure 7.4.1.3.7. A surface morphology change with solution temperature in electrolytic cadmium films grown on amorphous substrates from (A) sulfate and (B) chloride baths. The current density ¼ 1000 A=m2 ; the film thickness ¼ 6:7 mm; and pH ¼ 1:8:

Electrolytic Cd

303

Figure 7.4.1.3.8. A texture index change with solution temperature in electrolytic cadmium films grown on amorphous substrates from (A) sulfate and (B) chloride baths. The current density ¼ 1000 A=m2 ; the film thickness ¼ 6:7 mm; and pH ¼ 1:8:

Figure 7.4.1.3.9. A texture index change with solution temperature in electrolytic cadmium films grown on amorphous substrates from a sulfate bath. The current density ¼ 1000 A=m2 ; the film thickness ¼ 6:7 mm; and pH ¼ 1:8:

304

Nano-Plating

Figure 7.4.1.3.10. A surface morphology change with solution pH in electrolytic cadmium films grown on amorphous substrates from a sulfate bath, whose pH was adjusted using H2SO4 as a buffer. The current density ¼ 1000 A=m2 ; the film thickness ¼ 6:7 mm; and the solution temperature ¼ 25 8C:

Electrolytic Cd

305

Figure 7.4.1.3.11. A texture index change with solution pH in the of electrolytic cadmium films grown on amorphous substrates from a sulfate bath. The current density ¼ 1000 A=m2 ; the film thickness ¼ 6:73 mm; and the solution temperature ¼ 25 8C:

306 Nano-Plating Figure 7.4.1.3.12. A surface morphology change with film thickness in cadmium films plated on the {100}, {110}, and {111} planes of copper single-crystal substrates.

Electrolytic Cd

307

REFERENCE Miyazaki, S. (1998) Undergraduate Thesis, Dept. of Industrial Chemistry, Tokyo Metropolitan University.

FURTHER READING Aotani, K. (1953) Denki Kagaku, 21, 21.

7.4.1.4. Electrolytic Co (Imafuji, 1997) 7.4.1.4.1. COBALT METAL

Cobalt metal undergoes a hcp ! fcc phase transformation point at 417 8C. The lattice constants of the hcp phase are a ¼ 0:25071 nm and c ¼ 0:40686 nm and the ratio of c=a is 1.6228. If the atoms forming the hcp structure are assumed to be spherical, the ideal ratio of c=a is 1.663. Compared with this ideal ratio the hcp structure of cobalt is smaller, indicating that the cadmium lattice is slightly compressed along the , 001 . direction. The lattice constant of the fcc phase is a ¼ 0:25071 nm: The melting point is 1495 8C, the specific gravity is 8.85 g/cm3, and the resistivity is 6:24 £ 1026 V=cm:

7.4.1.4.2. PLATING CONDITIONS

Cobalt sulfate and cobalt chloride baths were used in this study. In both baths, 0.5 mol/l cobalt sulfate and 0.5 mol/l cobalt chloride were used. The bath pH was adjusted to 2.0 with the addition of dilute sulfuric acid or dilute hydrochloric acid. Additionally, a number of Co films were plated in a sulfamate bath. The plating was performed under the condition of constant current density. The film thickness for all films was set to 5 mm, which was obtained by adjusting the plating time while considering the current efficiency. Amorphous Fe –Co – Si – B foils (manufactured by Nippon Amorphous Metals Co., Ltd.) obtained by a splat-quenching method were used as the substrates.


A surface morphology change with a current density change in electrolytic Co films obtained from (A) sulfate and (B) chloride baths is seen in Figure 7.4.1.4.1. For both baths, the film surface was rough at low current densities, but became smoother at high current densities. The surface roughness of the Co film from the chloride bath is smaller than the results from the sulfate bath, especially at low current densities. The crystal size (grain size) of Co films from (A) sulfate and (B) chloride baths is plotted in Figure 7.4.1.4.2 as a function of current density. The crystal size was calculated from the width of an X-ray diffraction peak using Scherrer’s equation. The crystal size of the Co films from the sulfate bath is small at low and high current densities but shows a maximum value (100 nm) at 1000 A/m2. The crystal size of the Co films from the chloride bath is large at low current densities but decreases toward high current densities. This trend 308

Electrolytic Co

309

Figure 7.4.1.4.1. SEM micrographs showing how the surface morphology of electrolytic cobalt films grown on amorphous substrates from (A) sulfate and (B) chloride baths changes with current density. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; and pH ¼ 2:

310

Nano-Plating

Figure 7.4.1.4.2. A change in the crystal size of electrolytic cobalt films with current density. The cobalt films obtained from (A) sulfate and (B) chloride baths are compared. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; pH ¼ 2; and the substrate ¼ amorphous:

in the grain size can be seen in Figure 7.4.1.4.3. Figure 7.4.1.4.3(A) shows BF and DF TEM micrographs that display the grain structure of Co films plated at current densities of 10, 1000, and 2000 A/m2 from the sulfate bath. BF and DF TEM micrographs are placed at the upper and lower sides together with the associated electron diffraction pattern in the middle. According to TEM measurements, the grain sizes of the Co films grown at 100 A/m2 from the sulfate bath were in the order of several tens of nm, which is comparable to the results given in Figure 7.4.1.4.2, whereas the film grown at 1000 A/m2 was , 100 nm. At 2000 A/m2, the grain sizes again decreased to several tens of nm. TEM micrographs in Figure 7.4.1.4.3(B) indicate that the grain sizes of Co films from the chloride bath are large at low current densities and small at high current densities. SEM and TEM cross-sections for Co films plated at 1000 A/m2 from (A) sulfate and (B) chloride baths are shown in Figure 7.4.1.4.4. Figure 7.4.1.4.4(a) –(c) are SEM, BF TEM, and DF TEM images, respectively. The growth direction is indicated with an arrow in Figure 7.4.1.4.4(a) and (b). The SEM cross-section images indicate the columnar growth and rough surface in the Co films. Cross-section TEM micrographs revealed that the size and shape of Co grains from the sulfate bath is indeed different from those obtained from the chloride bath. It is clear that individual column or surface roughness observed by the SEM is not due to a single grain, but rather an assembly of fine grains. Co films from the chloride bath contained a fiber axis along the growth direction.

Electrolytic Co Figure 7.4.1.4.3. The effect of current density on the structure of electrolytic cobalt films grown on amorphous substrates from (A) chloride and (B) sulfate baths. TEM BF and DF images are shown together with the corresponding electron diffraction patterns for each current density. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; and pH ¼ 2:

311

312 Nano-Plating

Figure 7.4.1.4.3 (continued )

Electrolytic Co Figure 7.4.1.4.4. Cross-sectional images of electrolytic cobalt films grown on amorphous substrates from (A) sulfate and (B) chloride baths. (a) SEM, (b) TEM BF, and (c) TEM DF images. The current density ¼ 1000 A=m2 ; the film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; and pH ¼ 2:

313

314

Nano-Plating

Figure 7.4.1.4.5. A texture index change with current density in electrolytic cobalt films grown on amorphous substrates from (A) sulfate and (B) chloride baths. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; and pH ¼ 2:

The texture change with increasing current density is plotted in Figure 7.4.1.4.5. For both the sulfate and chloride baths, the texture was initially 100, but changed to 110 with increasing current density. Note that at low current densities, the 100 texture from the chloride is stronger than the texture from the sulfate bath. The change in grain size with increasing bath temperature is plotted in Figure 7.4.1.4.6. For both sulfate and chloride baths, the grain size increases with increasing bath temperature. The grain size from the sulfate bath increases linearly, however, the grain size from the chloride bath suddenly becomes large at 60 8C.

Figure 7.4.1.4.6. A crystal size change with solution temperature in electrolytic cobalt films grown on amorphous substrates from (A) sulfate and (B) chloride baths. The film thickness ¼ 5 mm; and pH ¼ 2:

Electrolytic Co

315

Figure 7.4.1.4.7. The effect of the addition of NaCl to the sulfate bath on the surface morphology of electrolytic cobalt films. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; pH ¼ 2; and the substrate ¼ amorphous:

316

Nano-Plating

Figure 7.4.1.4.8. A change in the crystal size of electrolytic cobalt films with the addition of NaCl to the sulfate bath. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; pH ¼ 2; and the substrate ¼ amorphous:

To study the effect of Cl2 concentration on the surface morphology of Co films, various quantities of NaCl were added to the basic sulfate bath. Current densities of (A) 100 and (B) 1000 A/m2 were used and the results are shown in Figure 7.4.1.4.7. The surface roughness is small for a small amount of NaCl (see Figure 7.4.1.4.7(a)), but becomes large with a large amount (see Figure 7.4.1.4.7(d)). Characteristic pointed features seen in the films from the chloride bath (cf. Figure 7.4.1.4.1(B)) appeared on the surface (see Figure 7.4.1.4.7(A) (d)). The grain sizes of Co films from the sulfate bath were plotted in Figure 7.4.1.4.8 as a function of the quantity of NaCl and current density. For all sulfate baths containing various quantities of NaCl, the grain size is small at low current densities, then becomes temporarily large at 200 A/m2, and finally decreases at higher current densities. Thus, the addition of NaCl to the sulfate bath affects the Co films by changing the grain size. Figure 7.4.1.4.9 depicts a surface morphology change with respect to current density in electrolytic Co films grown in a sulfamate bath. The surface morphology is markedly different from those grown in sulfate or chloride baths (see Figure 7.4.1.4.1). The surface morphology change with respect to pH in a sulfamate bath is shown in Figure 7.4.1.4.10.

Electrolytic Co

317

Figure 7.4.1.4.9. A surface morphology change in electrolytic cobalt films grown at (a) 500 and (b) 2000 A/m2 from the sulfamate bath. The film thickness ¼ 5 mm; the solution temperature ¼ 30 8C; pH ¼ 4; and the substrate ¼ amorphous:

Figure 7.4.1.4.10. A surface morphology change in electrolytic cobalt films grown on amorphous substrates from the sulfamate bath, whose pH were (a) 1.5 and (b) 4.0. The current density ¼ 500 A=m2 ; the film thickness ¼ 5 mm; and the solution temperature ¼ 30 8C:

The pH was adjusted using a H2SO2 buffer. pH is seen to have a dramatic effect on the surface morphology.

REFERENCE Imafuji, K. (1997) Graduation Thesis, Microstructure of electrodeposited Co film, Tokyo Metropolitan University.

7.4.1.5. Electrolytic Cr 7.4.1.5.1. CHROMIUM METAL

The structure of chromium metal is body-centered cubic (bcc) and the lattice constant is a ¼ 0:2884 nm: The melting point is 1875 8C and the resistivity is 12:9 £ 1026 V=cm:


A Sargent solution containing 2.5 mol/l CrO3 and 0.025 mol/l H2SO4 is generally used for plating Cr films. The effect of adding halogens (Cl, Br, and I) on the structure of a plated Cr metal will also be briefly discussed in this experiment. Amorphous 25 m-thick Fe –Co – Si – B foils (manufactured by Nippon Amorphous Metals Co., Ltd.) obtained by a splatquenching method were used as substrates.


Figure 7.4.1.5.1 shows surface morphologies in electrolytic Cr films obtained from the 50 8C Sargent bath as a function of current density. Although all the films contained cracks, the film surface obtained at low current densities was smooth, but the film grown at high current densities was rough. A change in the grain size with increasing current density is plotted in Figure 7.4.1.5.2, where the grain size was determined using Scherrer’s equation (see Eq. (6.2) in Section 6.3.2). Although the plot exhibits a data scatter, all the films consist of an assembly of 150-nm grains regardless of the magnitude of the current density. The texture index change as a function of current density is plotted in Figure 7.4.1.5.3. This graph also shows a data scatter. It is, however, reasonable to conclude that the films possess the 111 texture independent of the current density. A texture index change with increasing film thickness in the Cr films grown at 20 A/m2 is shown in Figure 7.4.1.5.4. The thin Cr films do not display any texture, probably because their grains initially randomly nucleated on the amorphous substrate. The Cr films eventually exhibited the 111 texture with increasing film thickness. A surface morphology change with bath temperature in the Cr films plated at 20 A/m2 is shown in Figure 7.4.1.5.5. The surface is rough at low temperatures but becomes smooth at high temperatures. The grain size does not change with bath temperature, as seen in Figure 7.4.1.5.6. Finally, the texture is 100 at low temperatures but changes to 111 at 40 8C. We also investigated the effect of halogen ions on the structure of the Cr films. No halogen ion effect was observed in the grain size, but the addition of halogen ions to the 318

Electrolytic Cr

319

Figure 7.4.1.5.1. A surface morphology change with current density in electrolytic chromium films grown on amorphous substrates.

320

Nano-Plating

Figure 7.4.1.5.2. A plot showing a change in the crystal size of electrolytic chromium films with current density.

Figure 7.4.1.5.3. A texture index change with current density in electrolytic chromium films grown on amorphous substrates.

Figure 7.4.1.5.4. Atextureindex change with film thickness inelectrolytic chromiumfilms grown on amorphoussubstrates.

Electrolytic Cr

321

Figure 7.4.1.5.5. A surface morphology change with solution temperature in electrolytic chromium films grown on amorphous substrates.

322

Nano-Plating

Figure 7.4.1.5.6. A crystal size change with increasing solution temperature in electrolytic chromium films grown on amorphous substrates.

Figure 7.4.1.5.7. A texture index change with increasing solution temperature in electrolytic chromium films grown on amorphous substrates.

bath caused the surfaces to become rougher. The texture index tends to become smaller with the addition of halogen ions. In particular, the addition of iodine ions resulted in a loss of texture (Figure 7.4.1.5.7). REFERENCES Yoshioka, S. (1999) Undergraduate Thesis, Dept. of Industrial Chemistry, Tokyo Metropolitan University.

7.4.1.6. Electrolytic Cu 7.4.1.6.1. COPPER METAL

The structure of copper metal is fcc and its lattice constant is a ¼ 0:3615 nm: The melting point is 1083 8C and its resistivity is 1:6730 £ 1026 V=cm:


Table 7.4.1.6.1 lists the composition of sulfate and pyrophosphate baths used for plating copper films in this study. Copper sulfate is soluble in water and the bath pH was adjusted with dilute sulfuric acid. Copper pyrophosphate is insoluble in water but can become soluble by forming complex ions if potassium pyrophosphate is present in larger amounts than copper pyrophosphate. Plating was performed under constant-current conditions with the plating time being controlled through the use of a coulomb meter. The bath was agitated with a stirrer. The pH of the fresh pyrophosphate bath used for plating was 8.7. The substrates were amorphous Fe –Co – Si – B foils (manufactured by Nippon Amorphous Metals Co., Ltd.) obtained by a splat-quenching method or Cu sheets covered with an electrolytic Ni – 32 at.% P amorphous film. The purpose of using these substrates was to obtain electrolytic Cu films, whose structure was not affected by the substrate structure.


7.4.1.6.3.1 Plating in a sulfate bath Figure 7.4.1.6.1 displays the change in surface morphology with respect to increasing film thickness in Cu films plated at current densities of (a) 500 and (b) 1000 A/m2 in a sulfate bath containing 0.5 mol/l CuSO4·H2O. For both current densities, the surface irregularity was small when the films were thin, but increased for the thicker films. The texture index change in the Cu films described in Figure 7.4.1.6.1 is plotted in Figure 7.4.1.6.2 as a function of film thickness. At 500 A/m2, the Cu film displayed a strong 110 texture from the beginning. The high texture index (close to 10) indicates that almost all the grains in the film are oriented along the , 110 . direction. At 1000 A/m2, on the other hand, the film initially did not exhibit any particular texture but suddenly displayed the 110 texture with an increase in the film thickness. Motivated by a previous report that bath aging often causes a change in the microstructure and properties of plated films, we incorporated 323

324

Nano-Plating Table 7.4.1.6.1. Bath compositions used for plating copper films. Bath name

Bath composition

Sulfate bath Pyrophosphate bath

CuSO4·H2O (1 mol/l) CuP2O7·3H2O (0.2 mol/l) K4P2O7·3H2O (0.7 mol/l)

the use of an aged bath in this experiment. The change in surface morphology with increasing current density is shown in Figure 7.4.1.6.3 for electrolytic Cu films grown from (a) freshly prepared and (b) aged (1 month) sulfate baths. The surface appears to be rough at low current densities but becomes smoother at higher current densities. This trend is the same for both fresh and aged baths. The texture index was measured for the same set of Cu films shown in Figure 7.4.1.6.3 (see Figure 7.4.1.6.4). In Figure 7.4.1.6.2, we found that the film thickness, at which the current density affects the texture most sensitively, was 5 mm. The same film thickness was used to construct Figure 7.4.1.6.4. In both fresh and aged baths, the 110 texture is strong at low current densities but becomes weaker at high current densities. Here, the 111 texture becomes stronger. In detail, we note that there was a small difference in the texture index plots between the fresh and aged baths. Compared with the 110-texture curve of the aged bath, the curve of the fresh bath is shifted slightly more toward the top right side (the high current density side), suggesting that the 110 texture appeared more readily from the freshly prepared bath than from the aged bath. A change in surface morphology with bath temperature in 5 mm-thick electrolytic Cu films grown at (a) 500 and (b) 1000 A/m2 from a sulfate bath is shown in Figure 7.4.1.6.5. At 1000 A/m2, the bath temperature appears to have very little effect on the surface morphology, whereas at 500 A/m2, large surface irregularities developed with increasing bath temperature. In other words, the effect of the bath temperature on the surface morphology is small at high current densities but is large at low current densities. The texture index of the Cu films studied in Figure 7.4.1.6.5 was plotted as a function of bath temperature in Figure 7.4.1.6.6. For both current densities (500 and 1000 A/m2), a weak 111 texture appeared at low bath temperatures, but the 110 texture appeared at , 30 8C. The 110 texture rapidly increased with increasing bath temperature. 7.4.1.6.3.2 Plating in a pyrophosphate bath The effect of film thickness and current density on the change in surface morphology of electrolytic copper films grown from a pyrophosphate bath is depicted in Figure 7.4.1.6.7. It is clear that the films deposited at a current density of 1000 A/m2 exhibit anomalously large surface roughness. The development of such a rough surface can be attributed to the high diffusivity of metal ions. At current densities below 500 A/m2, the surface roughness increased slowly with increasing film thickness, which is not as prominent as the films

Electrolytic Cu

325

Figure 7.4.1.6.1. A surface morphology change with increasing film thickness in electrolytic copper films grown on amorphous substrates from a 60 8C sulfate bath.

326

Nano-Plating

Figure 7.4.1.6.2. A texture index change with increasing film thickness in electrolytic copper films grown on amorphous substrates from a 60 8C sulfate bath.

grown from a sulfate bath. For the copper films grown from a pyrophosphate bath, the effect of current density on the surface morphology is very small. A texture index change with respect to film thickness and current density in electrolytic copper films shown in Figure 7.4.1.6.7 was plotted in Figure 7.4.1.6.8. For all current densities, the films initially took the 111 texture. Furthermore, at low current densities, the films maintained the 111 texture regardless of film thickness. When the current density was increased, however, the 111 texture of the initial films became weaker, and suddenly switched to the 110 texture as the film thickness was increased. A texture index contour map for the electrolytic copper films is plotted in Figure 7.4.1.6.9 as a function of film thickness and current density. When the films are thin, they take the 111 texture at the low current density side. As the films became thicker, the texture changes to 110 at the high current density side. A texture index change with respect to solution temperature and current density in the 5 mm-thick electrolytic copper films is plotted in Figure 7.4.1.6.10. For all current densities, the texture changed from 100 to 111 at , 60 8C. The 100 ! 111 texture change was also observed in the Cu films obtained from a sulfate bath, although the transition bath temperature was different, as illustrated in Figure 7.4.1.6.6.

REFERENCES Tomita, M. (1999) Microstructure of electrodeposited Cu film, Graduation thesis, Tokyo Metropolitan University. Miura, Y. (2000) Microstructure of electrodeposited Cu film, Graduation thesis, Tokyo Metropolitan University.

Electrolytic Cu

327

Figure 7.4.1.6.3. A surface morphology change with current density in 5 mm thick electrolytic copper films grown on amorphous substrates from a 30 8C sulfate bath. Two types of baths were used, i.e. (a) as-prepared and (b) aged for 1 month.

328

Nano-Plating

Figure 7.4.1.6.4. A texture index change with current density in 5 mm thick electrolytic copper films grown on amorphous substrates from a 30 8C sulfate bath ðpH ¼ 2Þ: Two types of baths were used, i.e. (a) as-prepared and (b) aged for 1 month.

Figure 7.4.1.6.5. A surface morphology change with solution temperature and current density in 5 mm thick electrolytic copper films grown on amorphous substrates from a sulfate bath ðpH ¼ 2Þ:

Electrolytic Cu

329

Figure 7.4.1.6.6. A texture index change with solution temperature and current density in 5 mm thick electrolytic copper films grown on amorphous substrates from a sulfate bath ðpH ¼ 2Þ:

330

Nano-Plating

Figure 7.4.1.6.7. A surface morphology change with film thickness and current density in electrolytic copper films grown on amorphous substrates from a 55 8C pyrophosphate bath ðpH ¼ 8:7Þ:

Electrolytic Cu

331

Figure 7.4.1.6.8. A texture index change with film thickness and current density in electrolytic copper films grown on amorphous substrates from a 55 8C pyrophosphate bath ðpH ¼ 8:7Þ:

332

Nano-Plating

Figure 7.4.1.6.9. A texture index contour map plotted as a function of film thickness and current density for electrolytic copper films grown on amorphous substrates from a 55 8C pyrophosphate bath.

Figure 7.4.1.6.10. A texture index change with solution temperature and current density in electrolytic copper films grown on amorphous substrates from a pyrophosphate bath ðpH ¼ 8:7Þ:

7.4.1.7. Electrolytic Fe 7.4.1.7.1. IRON METAL

Pure iron metal undergoes two phase transformations; one at 910 8C from the bcc phase (called a-Fe or a-ferrite) to the fcc phase (called g-Fe or austenite) and the other at 1390 8C from the g-Fe to the bcc phase called d-ferrite, which remains up to the melting point, 1534 8C. The lattice constants, a, of a-ferrite and austenite are 0.28664 and 0.363 nm, respectively. The resistivity is 9:71 £ 1026 V=cm: Iron metal also undergoes a ferromagnetic-to-paramagnetic transformation at 770 8C (Curie point).


In this study, sulfate, chloride and sulfamate baths were prepared to plate Fe films. An iron salt concentration of 1 mol/l was used for all baths. Since iron sulfamate alone does not dissolve in water, 0.1 mol/l HFNH was added to the sulfamate bath as a complexing agent. Sulfate and chloride baths contained 0.5 mol/l H3BO3 as a buffer. The pH of the sulfate and sulfamate baths were adjusted to 2.0 and 2.5, respectively, with H2SO4. The pH of the chloride bath was adjusted to 2.5 with HCl. The plating was done at current densities of 500, 1000, and 2000 A/m2 and the effect of film thickness (1,5, and 10 mm) on the surface morphology and texture was studied. The bath temperature was kept at 50 8C. The substrate materials were rolled Cu sheets, which were electropolished to a mirror finish and then coated with a 2 mm-thick amorphous Ni– P film obtained by an electroless plating method. The purpose of using amorphous substrates is to obtain plated films whose structures are not affected by the substrate structure. Single-crystal {100}, {110}, and {111} substrates were also used to study the crystallographic lattice-matching (epitaxial) relationship between plated Fe films and the single-crystal Cu. A platinum plate was used as an anode and bath agitation was performed with a magnetic stirrer. The film thickness was calculated by weighing the plated film. This film thickness value should be regarded as the average thickness, in particular, for films having severe surface irregularities. The current efficiency was 80% for all baths used.


Figure 7.4.1.7.1(A) displays the surface morphology of 1 mm-thick Fe films grown on (a) {100}, (b) {110}, and (c) {111} single-crystal Cu substrates at 1000 A/m2 from a 333

334 Nano-Plating Figure 7.4.1.7.1. Electrolytic iron films (1 mm thick) grown from a sulfate bath at the current density of 1000 A/m2 on (a) single-crystal {100} copper, (b) {110} copper, and (c) {111} copper substrates. (A) The surface morphology of plated iron films and (B) indexed electron diffraction patterns from the iron film/copper substrate composite films.

Electrolytic Fe

335

sulfate bath. Characteristic geometrical patterns are seen in these SEM micrographs. The appearance of these well-defined geometrical patterns originates from the growth facets developed on the surface of the plated Fe films, which have grown epitaxially on the Cu substrates. Based on these geometrical patterns, we inferred the lattice-matching relationships as seen in Figure 7.4.1.7.1(B). A deduction of the exact relationship between the Fe films and the single-crystal Cu substrates requires TEM examinations of plated Fe film/Cu substrate composite specimen, followed by an analysis of the electron diffraction pattern. Figure 7.4.1.7.2 shows a change in surface morphology with respect to an increase in the thickness of Fe films grown on single-crystal {110} Cu substrates from (a) sulfate, (b) chloride, and (C) sulfamate baths. The Fe films from the sulfate bath display the geometrical pattern characteristic of epitaxial growth up to a thickness of 5 mm, but those from the chloride bath had lost the geometrical pattern by 1 mm. From the sulfamate bath, both geometrical and random patterns are mixed in a 5 mm-thick Fe film, suggesting that epitaxial and non-epitaxial Fe grains coexist. It is clear from these micrographs that the critical thickness, at which epitaxy or oriented growth breaks down, depends on the type of plating baths used. In this experiment, the critical thickness is 7.5 mm for the sulfate bath, , 1 mm for the chloride bath, and 5 mm for the sulfamate bath. The change in surface morphology with respect to the change in current density and film thickness for Fe films plated on amorphous Ni– P substrates from a sulfate bath is seen in Figure 7.4.1.7.3. These micrographs represent the surface morphologies of those plated Fe films that are not affected by the structure of the substrate. An increase in the current density is seen to slightly reduce the surface roughness. For all current densities, the surface roughness increased at a thickness of 1 – 5 mm, but remained unchanged between 5 , 10 mm: The change in surface morphology for a chloride bath is shown in Figure 7.4.1.7.4. This bath slightly reduced the surface roughness at high current densities also. In addition, severe surface roughness developed with increasing film thickness. The change in surface morphology for a sulfamate bath is shown in Figure 7.4.1.7.5. This bath slightly reduced the surface roughness at high current densities also. In this bath, the surface roughness increased with increasing film thickness, but only a small change was noticed with increasing current density. A texture index change with respect to current density and film thickness in Fe films grown on amorphous Ni– P substrates from a sulfate bath is shown in Figure 7.4.1.7.6. For all current densities used, the films did not initially display any texture, suggesting that the Fe nuclei were randomly oriented on the amorphous substrate. With increasing film thickness, the 211 texture developed. The 211 texture was strong at low current densities but became weaker at high current densities. A texture index change with film thickness and current density in Fe films obtained from a chloride bath is shown in Figure 7.4.1.7.7. For all current densities, the thin Fe films

336 Nano-Plating Figure 7.4.1.7.2. A surface morphology change with increasing film thickness in electrolytic iron films grown on {110} single-crystal copper substrates from (a) sulfate, (b) chloride, and (c) sulfamate baths at the current density of 1000 A/m2.

Electrolytic Fe Figure 7.4.1.7.3. A surface morphology change with increasing film thickness and current density in electrolytic iron films grown on amorphous Ni –P films from a sulfate bath.

337

338 Nano-Plating Figure 7.4.1.7.4. A surface morphology change with increasing film thickness and current density in electrolytic iron films grown on amorphous Ni –P films from a chloride bath.

Electrolytic Fe

339

Figure 7.4.1.7.5. A surface morphology change with increasing film thickness and current density in electrolytic iron films grown on amorphous Ni –P films from a sulfamate bath.

340 Nano-Plating

Figure 7.4.1.7.6. A texture index change with film thickness and current density in iron films grown on amorphous Ni–P films from a sulfate bath.

Electrolytic Fe

Figure 7.4.1.7.7. A texture index change with film thickness and current density in iron films grown on amorphous Ni –P films from a chloride bath.

341

342 Nano-Plating

Figure 7.4.1.7.8. A texture index change with film thickness and current density in iron films grown on amorphous Ni –P films from a sulfamate bath.

Electrolytic Fe

343

are randomly oriented and do not show any characteristic texture. In general, the Fe films from the chloride bath do not exhibit any particular texture even with increasing film thickness. A texture index change with film thickness and current density in Fe films obtained from a sulfamate bath is shown in Figure 7.4.1.7.8. Similar to the chloride bath, the thin Fe films do not show any characteristic texture for all current densities. With increasing film thickness, however, the films showed the 211 texture, which became more prominent at higher current densities. A similar result was also reported by Ohno (1988).

REFERENCES Inoue, K., Nakata, T. & Watanabe, T. (2001) J. Jpn Inst. Metals, 65, 229. Inoue, K., Nakata, T. & Watanabe, T. (2002) Mater. Trans., 43, 1318.

FURTHER READING Ohno, I. (1988) J. Metal Finish. Soc. Jpn, 39, 149.

7.4.1.8. Electrolytic Ni (Kitagawa, 1993; Mizushima, 1994; Takasaka, 1995; Doi, 2001) 7.4.1.8.1. NICKEL METAL

The structure of nickel metal is fcc and the lattice constant is a ¼ 0:35238 nm: The melting point is 1453 8C and the resistivity is 12:5 £ 1026 V=cm:


Four types of Ni plating baths were used in this study. These baths include sulfate, chloride, Watts (0.9 kmol/m3 sulfate salt þ 0.09 kmol/m3 chloride salt), and sulfamate. In addition, five modified Watts baths were prepared by changing the mixing ratio of NiSO4 and NiCl2 (see Table 7.4.1.8.1(b)). All the baths contained 0.5 kmol/m3 H3BO3. The pHs of all the baths were set to 2.8. A magnetic stirrer (350 rpm) was used to agitate the bath during plating. The bath temperature was kept at 40 8C. Amorphous splat-quenched Fe –Co – Si – B foils (manufactured by Nippon Amorphous Metals Co., Ltd.) were used as the substrates. The foils were first ultrasonically cleaned in acetone and then etched in a solution of dilute sulfuric acid immediately before plating. A platinum plate ð1:5 £ 2:0 cm2 Þ was used as an anode and the inter-electrode distance was , 5.0 cm. The plating was conducted at a constant current density of 1500 A/m2. The thickness of the plated films were calculated by measuring the weight gain. Most of the Ni films were

Table 7.4.1.8.1. Solution compositions for plating nickel films in this study. (a) Four types of baths (kmol/m3) (a) Sulfate bath NiSO4 NiCl2 Ni(S2NH2)2 H3BO3

0.99 – – 0.5

(b) Watts bath

(c) Chloride bath

(d) Sulfamate bath

0.9 0.99 – 0.5

– 0.99 – 0.5

– – 0.99 0.5

(b) Modified Watts baths (the modification is done by changing the mixing ratio of sulfate and chloride baths) (kmol/m3) A bath B bath C bath D bath E bath NiSO4 NiCl2 H3BO3

0.1 – 0.5

0.75 0.25 0.5

0.5 0.5 0.5

344

0.25 0.75 0.5

– 1.0 0.5

Electrolytic Ni

345

plated until the total 72 C was consumed. The film thickness was varied by changing the electric charge. Figure 7.4.1.8.1 is a graph plotting film thickness with increasing electric charge in electrolytic Ni films grown at current densities of 100 and 1000 A/m2 from sulfate and chloride baths. It is seen that the deposition rate from the sulfate bath is different from that from the chloride bath. This graph also represents current efficiency under various plating conditions. We found that the current efficiency changed with plating conditions but remained in the range of 60 –80% for all the films. In Figure 7.4.1.8.1, the film thickness is seen to change linearly with respect to the electric charge.


7.4.1.8.3.1 Surface morphology and crystal size Figure 7.4.1.8.2 displays SEM micrographs showing surface morphologies as a function of current density and quantity of electric charge (or film thickness) in electrolytic Ni films grown on the amorphous substrates from sulfate (Figure 7.4.1.8.2.1), chloride (Figure 7.4.1.8.2.2), and Watts’ baths (Figure 7.4.1.8.2.3). All the specimens were tilted by 308 to enhance the image of the surface irregularities. It is important to remember that the micrographs in the vertical direction appear to be compressed, while those in the horizontal direction appear without any compression. The surface roughness (surface irregularity

Figure 7.4.1.8.1. Relationship between film thickness and coulomb in electrolytic nickel films obtained from sulfate and chloride baths.

346 Nano-Plating

Figure 7.4.1.8.2.1. A surface morphology change with current density and coulomb in electrolytic nickel films grown from a sulfate bath.

Electrolytic Ni

Figure 7.4.1.8.2.2. A surface morphology change with current density and coulomb in electrolytic nickel films grown from a chloride bath.

347

348 Nano-Plating

Figure 7.4.1.8.2.3. A surface morphology change with current density and coulomb in electrolytic nickel films grown from a Watts bath.

Electrolytic Ni

349

along the film normal direction) of the films (0.2 mm thick) obtained from all baths is very similar and small (about 0.1 mm). The roughness increases with increasing film thickness. Surface morphologies for Ni films obtained from five types of baths (see Table 7.4.1.8.1) are shown in Figure 7.4.1.8.3(a) – (e). The angle of the surface asperities appears to become sharper with increasing nickel chloride content. A TEM was also used to determine how the grain size changes with increasing film thickness in electrolytic Ni films grown from (A) sulfate and (B) chloride baths. These plan-view TEM samples were prepared by electropolishing 0.2, 1.8, and 8 mm-thick films from both surfaces. From these samples, the grain structure of the films at a height of 0.1, 0.9, and 4 mm from the substrate surface were examined. The location of the films taken for TEM examinations is schematically illustrated in Figure 7.4.1.8.4(d). From Figure 7.4.1.8.4, the grain size was less than 0.05, 0.1, and 0.1 mm at the height of 0.1, 0.5, and 4 mm, respectively. Figure 7.4.1.8.5 shows surface morphologies versus bath temperature in electrolytic Ni films grown at current density of 150 A/m2 from (A) sulfate, (B) Watts, and (C) chloride baths. In the films from the sulfate bath, the lateral dimension of the surface roughness remained the same but the vertical dimension increased with increasing thickness.

Figure 7.4.1.8.3. Surface morphologies of electrolytic nickel films obtained from sulfate, chloride, and mixed baths (see Table 7.4.1.8.1(b)).

350 Nano-Plating Figure 7.4.1.8.4. Plan-view TEM micrographs showing the grain structure of electrolytic nickel films observed at three different depths (0.12, 1.3, and 8.0 mm) along the film thickness direction. The nickel films were grown at 100 and 1000 A/m2 from (A) sulfate and (B) chloride baths.

Electrolytic Ni

351

Figure 7.4.1.8.5. A surface morphology change with increasing solution temperature in electrolytic nickel films grown from (A) sulfate, (B) Watts, and (C) chloride baths.

The films from the chloride bath showed a similar trend, in which not only the height increased with increasing film thickness, but also a change in shape, toward a triangular pyramid. The change in surface morphology for the Watts bath appears to be the intermediate between those from the sulfate and chloride baths. Figure 7.4.1.8.6 shows the effect of solution agitation speed on the surface morphology of the films grown at a current density of 150 A/m2 from a chloride bath. Both the lateral and vertical dimensions of the surface roughness increased with increasing agitation speed. 7.4.1.8.3.2 Texture Figure 7.4.1.8.7 demonstrates the change in texture index with respect to current density and film thickness in Ni films obtained from a sulfate bath. At a current density of 100 A/m2, the films did not initially exhibit any texture, but displayed the 100 texture with

352 Nano-Plating Figure 7.4.1.8.6. Effect of solution agitation on the surface morphology of electrolytic nickel films grown on amorphous substrates from a chloride bath. The bath was agitated using a magnetic stirrer, whose agitation rate was varied by changing the rotation speed.

Electrolytic Ni

353

Figure 7.4.1.8.7. A texture index change with increasing coulomb and current density in electrolytic nickel films obtained from a sulfate bath.

increasing film thickness (or consumed electric charge). At current densities above 500 A/m2, the texture became 110. Figure 7.4.1.8.8 depicts the change in texture index versus current density for Ni films grown from four types of baths. The texture is weak at low current densities (or low deposition rates), but starts exhibiting a strong texture at high current densities (or high deposition rates). The texture is 110 for (A) the sulfate bath, 311 for (C) the chloride bath, 110 for (B) the Watts bath, and 100 for (D) the sulfamate bath. The result from the sulfamate bath agrees with the results reported by Verma and Wilman (1971). Figure 7.4.1.8.9 depicts the change in texture index versus film thickness (or consumed electric charge) for Ni films grown in Ni plating baths with various mixing ratios of NiSO4 and NiCl2 (see Table 7.4.1.8.1(b)). The films from the A bath, which is the simple sulfate

354

Nano-Plating

Figure 7.4.1.8.8. A texture index change with current density in electrolytic nickel films obtained from (A) sulfate, (B) Watts, (C) chloride, and (D) sulfamate baths.

bath, developed the 110 texture. On addition of Cl ions, such as in the B bath, where the ratio of NiSO4: NiCl2 was 3:1, this texture was lost completely. A further increase in the concentration of Cl ion, such as in the D bath, where the ratio of NiSO4: NiCl2 was 1:3 brought the 311 texture. Figure 7.4.1.8.10 depicts the effect of bath temperature on the texture index of Ni films plated from (a) sulfate, (b) Watts, and (c) chloride baths. The dependency of the texture on bath temperature is quite complex. In (A) the sulfate bath, the texture changed from 110 to 100 with increasing bath temperature, whereas in (C) the chloride bath, the texture moved from 311 to 110. Finally, in (B) the Watts bath, the texture began with 110, changed to 100 temporarily, and again returned to 110 at high temperatures. The effect of bath agitation on the film texture was also investigated, but no effect was found.

Electrolytic Ni

Figure 7.4.1.8.9. A texture index change with coulomb in electrolytic nickel films obtained from five mixed baths (see Table 7.4.1.8.1(b)).

355

356 Nano-Plating Figure 7.4.1.8.10. A texture index change with solution temperature in electrolytic nickel films obtained from (A) sulfate, (B) Watts, and (C) chloride baths.

Electrolytic Ni

357

REFERENCES Doi, S. (2001) Microstructure of electrodeposited Cu film, Graduation University. Kitagawa, Y. (1993) Microstructure of electrodeposited Cu film, Metropolitan University. Mizushima, I. (1994) Microstructure of electrodeposited Cu film, Metropolitan University. Takasaka, T. (1995) Microstructure of electrodeposited Cu film, Metropolitan University. Verma, S.K. & Wilman (1971) J. Phys., D: Appl. Phys., 4, 2051.

thesis, Tokyo Metropolitan Graduation thesis, Tokyo Graduation thesis, Tokyo Graduation thesis, Tokyo

7.4.1.9. Electrolytic Sn (Teshigawara, 2001) 7.4.1.9.1. TIN METAL

Tin metal has two allotropic forms at normal pressure and undergoes an allotropic transformation at 13.2 8C from gray tin (a-Sn) with a cubic structure to white tin (b-Sn) with a tetragonal structure (bct). When tin is cooled below 13.2 8C, it cannot be easily transformed, thus it is often supercooled to less than 2 40 8C. At a temperature below 2 40 8C, white tin transforms catastrophically to a powdery form of gray tin. The lattice constant for a-Sn is a ¼ 0:6489 nm; whereas those for b-Sn are a ¼ 0:5821 nm and c ¼ 0:3182 nm: The melting temperature is 232 8C and the resistivity is 11 £ 1026 V=cm:


A pyrophosphate bath can be used to plate tin films. The two types of baths used in this study are listed in Table 7.4.1.9.1. The bath temperature was changed from 30 to 80 8C, while the bath pH was varied from 6 to 10. Since the pH of the freshly prepared bath was 9.1, it was lowered with HCl and raised with ammonium water. The plating was conducted at current densities from 50 to 500 A/m2. Tin films with the thickness of 0.8 –62.4 mm were studied. The bath agitation was done using a magnetic stirrer. The anode was a 20 mm £ 25 mm platinum sheet. Copper sheets covered with an electroless amorphous Ni– 24.0 at.% P film were chosen as the substrates. We also used single-crystal {100}, {110}, and {111} Cu plates with a dimension of 10 mm £ 20 mm after electropolishing them to a mirror finish.


7.4.1.9.3.1 Surface morphology The surface morphology of Sn films plated on the {100}, {110}, and {111} faces of singlecrystal Cu plates from the A bath are shown in Figure 7.4.1.9.1. The plating conditions are described in the figure caption. On the {100} and {110} Cu surfaces, the Sn surface is rough when the film is thin and becomes rougher with increasing film thickness. On the {111} Cu surface, the Sn surface was relatively smooth until the film thickness reached 5.1 mm, at which point randomly distributed tetragonal Sn crystals covered the surface. The surface morphology of Sn films is clearly affected by the crystallographic orientation of the substrate surface. 358

359

Electrolytic Sn Table 7.4.1.9.1. Bath composition for tin plating. Composition

A bath (mol/l)

B bath (mol/l)

SnCl2·2H2O K4P2O7

0.15 0.45

0.5 0.15

Copper metal is fcc and Sn metal (a-Sn) is tetragonal at ambient temperatures. Thus, it is difficult for tetragonal Sn to lattice-match to the cubic Cu substrate. This was the case for the {100} and {110} substrates, on which Sn deposits nucleated randomly. On the {111} substrate, Sn metal can lattice-match easily to the {111} Cu substrate with a small misfit, as illustrated in Figure 7.4.1.9.2. The (111) Sn plane is seen to fit to the {111} Cu plane, where the , 112 . direction of Sn lines up along the , 011 . of Cu. This matching

Figure 7.4.1.9.1. A surface morphology change with increasing film thickness in electrolytic tin films grown on (a) {100}, (b) {110}, and (c) {111} single-crystal copper substrates from a 40 8C bath ðpH ¼ 8Þ containing 0.15 mol/l SnCl2·2H2O at a current density of 200 A/m2.

360

Nano-Plating

Figure 7.4.1.9.2. An expected interfacial matching between an electrolytic tin film and a {111} Cu substrate. The matching relationship is shown to follow (111) Sn//(111) Cu and ,112 . Sn//,011 . Cu.

relationship can be expressed in a compact form, i.e. (111) Sn//{111} Cu and , 112 . Sn//, 011 . Cu. The formation of smooth Sn films on the {111} Cu substrate can be attributed to good lattice-matching properties in this orientation. Our explanation is still theoretical and thus more definitive experiments that verify the lattice-matching relationship will be necessary. A change in surface morphology with current density and film thickness in Sn films from the A bath is shown in Figure 7.4.1.9.3, where the plating conditions are described in the figure caption. Below a thickness of 2.6 mm, the film exhibits relatively fine surface irregularities, whereas films thicker than 2.6 mm produced a very rough surface, on which block-shaped crystals formed. Prior to the formation of block-shaped crystals, the surface of thin films is rough at low current densities and smooth at high current densities. This trend is consistent with our theory that the state of current density concentration depends on the current density. At this moment, it is difficult to consider a formation mechanism for the block-shaped crystals with a variety of shapes.

Electrolytic Sn Figure 7.4.1.9.3. A surface morphology change with increasing film thickness and current density in electrolytic tin films grown on amorphous substrates from the 40 8C A bath ðpH ¼ 8Þ:

361

362

Nano-Plating

A change in the surface morphology of 0.8 mm-thick Sn films from the A bath is displayed in Figure 7.4.1.9.4 as a function of bath pH and current density. A film thickness of 0.8 mm was chosen, so that no block-shaped crystals, as observed in Figure 7.4.1.9.3, will appear in this film thickness range. As described above, the bath pH was lowered with HCl and raised with ammonium water. The surface morphology did not change when the bath pH was less than 8 but appeared different at pH . 8: Although this morphological change is due to a change in pH (hydrogen ion concentration), it can be considered to originate from a difference in the type of anions used as a buffer. For all pH used, the surface became rough at low current densities and smooth at high current densities. The change in surface morphology with respect to a change in bath temperature is shown in Figure 7.4.1.9.5. The film thickness was 0.8 mm. The surface is smooth from the low-temperature bath and rough from the high-temperature baths. The effect of Sn salt concentration and current density on the surface morphology is shown in Figure 7.4.1.9.6. For the low Sn salt concentrations, the surface is rough if the current densities are low, but is smooth for the high Sn salt concentration regardless of the magnitude of the current density. 7.4.1.9.3.2 Texture The texture change with increasing film thickness produced scattered results, primarily due to the appearance of block-shaped crystals (see Figure 7.4.1.9.3). To obtain more dependable data on the effect of bath temperature on texture index, we chose two film thickness values ((A) 0.8 and (B) 2.6 mm), below which no block-shaped crystals were formed (see Figure 7.4.1.9.7). For both film thickness values, the texture was 110 at low bath temperatures and was 101 at high bath temperatures. We note that this texture variation with bath temperature is more prominent for the thicker (2.6 mm) films. The weaker texture in the thinner films can be understood from the fact that the structure of the thinner films is affected by the structure of the substrate. Here, we used the amorphous Ni –P substrates, on which randomly oriented Sn grains are likely to have nucleated. Conversely, the thicker films, which were less affected by the substrate structure, should develop a stronger intrinsic film texture. Figure 7.4.1.9.8 describes a texture change for two film thickness values ((A) 0.8 and (B) 2.6 mm), with varying current density, and bath pH. For both thickness values, the bath, whose pH was adjusted to 6 with HCl, exhibited the 100 texture at low current densities and the 110 texture at high current densities. For the pH ¼ 8 bath (HCl as a buffer), the films exhibited a strong 110 texture for both current densities. For the bath, whose pH was adjusted to 10 with NH4, on the other hand, a strong 110 texture appeared at low current densities, but no texture showed up at high current densities. The texture change with current density is plotted in Figure 7.4.1.9.9 as a function of bath temperature and film thickness. For both film thickness values (0.8 and 2.6 mm), as

Electrolytic Sn

363

Figure 7.4.1.9.4. A surface morphology change with pH and current density in 0.8 mm-thick electrolytic tin films grown on amorphous substrates from the 40 8C A bath.

Figure 7.4.1.9.5. A surface morphology change with solution temperature in 0.8 mm-thick electrolytic tin films grown on amorphous substrates at the current density of 500 A/m2 from the A bath ðpH ¼ 6Þ:

364

Nano-Plating

Figure 7.4.1.9.6. A surface morphology change with the concentration of SnCl2·2H2O and current density in 0.8 mm-thick electrolytic tin films grown on amorphous substrates from the 20 8C A bath ðpH ¼ 6Þ:

indicated in Figure 7.4.1.9.7, the texture changed suddenly as the bath temperature increased from 40 to 60 8C. In the low-temperature baths, the texture was not too strong at low current densities, but became 110 at high current densities. In the high-temperature baths, the texture became weak at both current densities.

Figure 7.4.1.9.7. A texture index change with pH in electrolytic tin films grown on amorphous substrates at 200 A/m2 from the B bath ðpH ¼ 6Þ:

Electrolytic Sn

365

Figure 7.4.1.9.8. A texture index change with pH and current density in electrolytic tin films grown on amorphous substrates at 200 A/m2 from the 40 8C A bath.

Figure 7.4.1.9.10 shows the texture change as a function of bath temperature, current density, and two film thickness values ((A) 0.8 and (B) 2.6 mm). For both film thickness values, the texture was weak at low current densities. At high current densities, on the other

366

Nano-Plating

Figure 7.4.1.9.9. A texture index change with solution temperature and current density in (A) 0.8 mm- and (B) 2.6 mm-thick electrolytic tin films grown on amorphous substrates from the B bath ðpH ¼ 6Þ:

Electrolytic Sn

367

Figure 7.4.1.9.10. A texture index change with solution temperature and current density in (A) 0.8 mm- and (B) 2.6 mm-thick electrolytic tin films grown on amorphous substrates from the B bath ðpH ¼ 6Þ:

368

Nano-Plating

Figure 7.4.1.9.11. A texture index change with the concentration of SnCl2·2H2O and current density in (A) 0.8 mm- and (B) 2.6 mm-thick electrolytic tin films grown on amorphous substrates from the 20 8C A bath ðpH ¼ 6Þ:

hand, the texture of the films from the low-temperature bath became 110. From the high-temperature bath, the film lost the texture or displayed the 101 texture. Texture index versus current density is plotted in Figure 7.4.1.9.11 in terms of Sn salt concentration and film thickness. For the low Sn salt concentration bath, the film took a weak texture at low current densities, but exhibited a strong 110 texture at high current densities. For the high Sn salt concentration bath, the film displayed a strong 100 texture at low current densities and the 101 texture at high current densities.

REFERENCES Aotani, K. (1953) Denki Kagaku, 21, 21. Teshigawara, T., Nakata, T. & Watanabe, T. (2001) Scr. Mater., 44, 2285.

7.4.1.10. Electrolytic Zn (Watanabe et al., 2000) 7.4.1.10.1. ZINC METAL

Zinc metal is hexagonal-close-packed (hcp). The lattice constants are a ¼ 0:26649 nm and c ¼ 0:49470 nm: The ratio, c=a ¼ 1:856; is larger than the ideal packing ratio of 1.633, indicating that the Zn lattice is more stretched along the c axis. The melting temperature is 419.5 8C and the resistivity is 5:916 £ 1026 V cm:


The five types of Zn plating baths used are listed in Table 7.4.1.10.1. These are (A) sulfate, (B) chloride, (C) mixed (sulfate: chloride ¼ 1 : 1Þ; (D) acetic acid, and (E) nitric acid baths. For all baths, the Zn metal ion concentration was kept constant (0.83 mol/l). The effect of additives was studied by adding 0.128 mol/l cresol sulfonic acid or 0.625 mol/l citric acid to the sulfate bath. The plating conditions are listed in Table 7.4.1.10.2. A Zinc sheet was used as an anode.


7.4.1.10.3.1

Films plated on single-crystal copper substrates

Figure 7.4.1.10.1 shows the surface morphology of electrolytic Zn films grown at 1000 A/m2 on single-crystal (a) {100}, (b) {110}, and (c) {111} Cu in (A) sulfate and (B) chloride baths. The films are arranged according to film thickness. Zinc films from the sulfate bath display characteristic geometrical (crystallographic) patterns up to a film thickness of 28 mm. With a further increase in the thickness, however, the geometrical pattern was destroyed. Zinc films from the chloride bath had lost the geometrical pattern at a thickness of 5 mm. In all the cases, there are characteristic surface morphologies associated with the orientation of the substrate. For example, a hexagonal pattern is present on the {100} surface. Similarly, a pattern covered by periodic long rectangular bars appeared on the {110}, and a hexagonal pattern on the {111}. The development of the orientation-dependent surface morphology strongly suggests that the Zn films are latticematched to the Cu substrates. In Figure 7.4.1.10.2, probable lattice-matching relationships between Zn and variously orientated Cu are given. 369

370

Nano-Plating Table 7.4.1.10.1. Bath composition for zinc electrodeposits. Bath name

Composition

(A) Sulfate bath (B) Chloride bath (C) Mixed bath (D) Acetic acid bath (E) Nitric acid bath

ZnSO4·7H2O (0.83 kmol/m2) ZnCl4 (0.83 kmol/m2) ZnSO4·7H2O (0.415 kmol/m2) þ ZnCl4 (0.415 kmol/m2) Zn(CH3COOH)2·2H2O (0.83 kmol/m2) Zn(NO3)2·6H2O (0.83 kmol/m2)

7.4.1.10.3.2 Plating conditions, surface morphology, and texture 7.4.1.10.3.2.1. Bath type and current density change. The change in surface morphology with respect to current density in 5.6 mm-thick Zn films plated from the (A) sulfate, (B) chloride, and (E) nitric acid baths ðpH ¼ 4:66Þ are shown in Figure 7.4.1.10.3. In this experiment, amorphous alloy foils were used to obtain Zn films, which are not affected by the crystal structure of the substrate. Zinc films from the (A) sulfate bath displayed a geometrical pattern at both low and high current densities and the magnitude of the surface irregularity did not appear to depend on current density. Zinc films from the (B) chloride bath showed a surface morphology covered by fine-scale particles at low current densities but exhibited a characteristic geometrical pattern unique to the hexagonal crystal with increasing current density. Zinc films from the (E) nitric acid bath showed surface irregularities outlined by spherical features. The size of these spheres was small at low current densities but became larger at high current densities. All the Zn films studied above were subjected to a texture analysis using an X-ray diffraction method. Figure 7.4.1.10.4 is the texture index of electroplated Zn films plotted as a function of current density and type of bath ((A), (B), and (E)). At low current densities, no particular texture developed in the films grown from any bath. In Zn films from the sulfate and chloride baths the 101 texture emerged with increasing current

Table 7.4.1.10.2. Zinc plating condition. Plating condition

Condition

Current density Solution temperature Anode material Agitation Additives

50 , 5000 A/m2 5 , 80 8C Zinc sheet Magnetic stirrer (A) p-Cresol sulfonic acid (B) Citric acid

Electrolytic Zn

371

Figure 7.4.1.10.1. A surface morphology change with increasing film thickness in electrolytic zinc films grown on (a) {100}, (b) {110}, and {111} single-crystal copper substrates from 30 8C (A) sulfate and (B) chloride baths at the current density of 1000 A/m2.

372 Nano-Plating

Figure 7.4.1.10.2. An estimated diagram for a lattice-matching between an electrolytic zinc film and single-crystal copper substrate.

Electrolytic Zn

373

Figure 7.4.1.10.3. A surface morphology change with current density in 5.6 mm-thick electrolytic zinc films grown on amorphous substrates from 30 8C (A) sulfate, (B) chloride, and (E) nitric acid baths ðpH ¼ 4:66Þ:

density. In Zn films from the nitric acid bath, the 101 texture appeared weakly at 1000 A/m2 but was lost altogether with a further increase in current density. 7.4.1.10.3.2.2. Film thickness change. Figure 7.4.1.10.5 depicts the change in surface morphology with increasing film thickness for Zn films plated at 1000 A/m2 from the (A) sulfate, (B) chloride, (C) mixed, (D) acetic acid, and (E) nitric acid baths ðpH ¼ 4:66Þ: For all baths used, the surface of the Zn films initially displayed fine-scale irregularities, which increased with increasing film thickness. The shape of the convex part of the film

374

Nano-Plating

Figure 7.4.1.10.4. A texture index change with current density in 5.6 mm-thick electrolytic zinc films grown on amorphous substrates from 30 8C (A) sulfate, (B) chloride, and (E) nitric acid baths ðpH ¼ 4:66Þ:

Electrolytic Zn Figure 7.4.1.10.5. A surface morphology change with increasing film thickness in electrolytic zinc films grown on amorphous substrates from (A) sulfate, (B) chloride, (C) mixed, (D) acetic acid, and (E) nitric acid baths ðpH ¼ 4:66Þ at the current density of 1000 A/m2.

375

376

Nano-Plating

surface depends on the type of bath used. Similar to the case of Figure 7.4.1.10.3(A) and (B), the surface of Zn films obtained from the (A) sulfate and (D) acetic acid baths showed a pattern of hexagonal symmetry, which was defined by the {001} facet of the hexagonal Zn crystal. The geometrical pattern, which is generally defined by the crystallographic facets of Zn crystal, did not appear in Zn films grown in the other baths. In particular, the films from the nitric acid bath exhibited a spherical morphology, as shown in Figure 7.4.1.10.3(E). The surface morphology of the films from the (C) mixed bath appeared similar to the surface from the (B) chloride bath, indicating that the chloride component in the bath affected the morphology more than the sulfate component. All the Zn films shown in Figure 7.4.1.10.5 were analyzed using an X-ray diffraction method and their texture index determined using Willson’s equation. The results are summarized in Figure 7.4.1.10.6. All Zn films from the (A) sulfate, (B) chloride, (C) mixed, and (D) acetic acid baths exhibited the 001 texture from the beginning and kept that texture to a thickness of more than 50 mm. Zinc films from the (E) nitric acid bath showed a different texture behavior namely the 100 texture started weakly, but disappeared completely when the thickness reached more than 10 mm, leaving no characteristic textures.

7.4.1.10.3.2.3. Effect of pH. The surface morphology of Zn films grown on amorphous alloy foils from the (A) sulfate and (B) nitric acid baths are shown in Figure 7.4.1.10.7 as a function of bath pH. The current density was chosen from Figure 7.4.1.10.4, where a characteristic texture appeared to emerge at 100 A/m2. Similarly, the film thickness was selected from Figure 7.4.1.10.5, where a change in the surface morphology occurred at a thickness of 5.6 mm. The pH of the freshly prepared (A) sulfate and (B) nitric acid baths was 4.66 and 4.33, respectively. They were adjusted using sulfuric and nitric acid, respectively. For both baths, the pH values appear to have no effect on the surface morphology change. All the Zn films shown in Figure 7.4.1.10.7 were analyzed using an X-ray diffraction method and their texture index determined from the intensity of their diffraction peaks (see Figure 7.4.1.10.8). Zinc films from the sulfate bath exhibited a strong 001 texture for all pH values. Zinc films from the nitric acid bath displayed a weak 001 texture at pH values less than 2 but showed no particular texture for pH values more than 2.

7.4.1.10.3.2.4. Effect of solution temperature. A surface morphology change with bath temperature in 5.6 mm-thick Zn films grown at 100 A/m2 from the (A) sulfate bath is shown in Figure 7.4.1.10.9. In the low-temperature bath, a characteristic symmetry pattern of hexagonal crystal appeared, but in the high-temperature bath, no recognizable pattern showed up. The texture change with respect to bath temperature is plotted in Figure

Electrolytic Zn

377

Figure 7.4.1.10.6. A texture index change with increasing film thickness in electrolytic zinc films grown on amorphous substrates from 30 8C (A) sulfate, (B) chloride, (C) mixed, (D) acetic acid, and (E) nitric acid baths ðpH ¼ 4:66Þ at the current density of 1000 A/m2.

378

Nano-Plating

Figure 7.4.1.10.7. A surface morphology change with solution pH in 5.6 mm-thick electrolytic zinc films grown on amorphous substrates from 30 8C (A) sulfate and (B) chloride baths at the current density of 100 A/m2.

7.4.1.10.10. A strong 001 texture appeared for all bath temperatures except for small data fluctuations, indicating that the bath temperature did not affect the texture. 7.4.1.10.3.2.5. Effect of additives. The effect of additives on the surface morphology and texture of Zn films from the (A) sulfate bath was studied using cresol sulfonic acid and citric acid as additives. The surface morphologies of Zn films from the sulfate bath containing (a) no additives, (b) 0.128 kmol/m3 cresol sulfonic acid, (c) 0.625 kmol/m3 citric acid are shown in Figure 7.4.1.10.11 as a function of film thickness. Similar to the case of Figure 7.4.1.10.3(A), the surface of Zn films from the additive-free sulfate bath exhibited a characteristic hexagonal symmetry. The addition of cresol sulfonic acid changed the surface morphology completely. The addition of citric acid did not significantly change the surface morphology. X-ray diffraction patterns from the Zn films studied in Figure 7.4.1.10.11 are shown in Figure 7.4.1.10.12. In Figure 7.4.1.10.12(a), we displayed the standard X-ray diffraction

Electrolytic Zn

379

Figure 7.4.1.10.8. A texture index change with solution pH in 5.6 mm-thick electrolytic zinc films grown on amorphous substrates from 30 8C (A) sulfate and (E) nitric acid baths at the current density of 1000 A/m2.

pattern taken from the JCPDS file for a comparative purpose. X-ray diffraction patterns from Zn films obtained in the sulfate bath containing (b) no additives, (c) cresol sulfonic acid, (d) citric acid are shown in Figure 7.4.1.10.11 as a function of film thickness. Similar to the Zn films shown in Figure 7.4.1.10.3(A), Zn films from the additive-free sulfate bath displayed a strong 001 texture, which was absent or weak in Zn films from the baths containing the additives. It is interesting to note that the addition of cresol sulfonic acid increased the intensity of the (101) and (110) diffraction peaks, but reduced the intensity of the (103) peak. When citric acid was added to the bath, the (100) and (110) peaks did not appear, but instead, the (101), (002) and (103) peaks appeared strongly. From the above results, it can be concluded that there is no relationship between the texture and surface morphology in electroplated Zn films.

380

Nano-Plating

Figure 7.4.1.10.9. A surface morphology change with solution temperature in 5.6 mm-thick electrolytic zinc films grown on amorphous substrates from a sulfate bath ðpH ¼ 4:66Þ at the current density of 100 A/m2.

Figure 7.4.1.10.10. A texture index change with solution temperature in 5.6 mm-thick electrolytic zinc films grown on amorphous substrates from a sulfate bath at the current density of 1000 A/m2.

Electrolytic Zn

381

Figure 7.4.1.10.11. A surface morphology change with increasing film thickness in electrolytic zinc films grown on amorphous substrates at the current density of 1000 A/m2 from a 30 8C sulfate bath ðpH ¼ 4:66Þ:

382

Nano-Plating

Figure 7.4.1.10.12. A change in X-ray diffraction patterns with increasing film thickness and the type of an additive used in electrolytic zinc films grown on amorphous substrates at 1000 A/m2 from a 30 8C sulfate bath ðpH ¼ 4:66Þ containing (b) no additive, (c) cresol sulfonic acid as an additive, and (d) citric acid. For comparison, the standard X-ray diffraction pattern from unoriented zinc powder listed in the JCPDS file is shown in (a).

Electrolytic Zn REFERENCES Aotani, K. (1953) Denki Kagaku, 21, 21. Watanabe, T. & Minami, S. (2000) J. Jpn Inst. Metals, 64, 67.

383

7.4.2 Pure Alloys 7.4.2.1. Electrolytic Ag–Cd 7.4.2.1.1. PLATING METHOD

The mixing ratio of AgNO3 and CdSO4 salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 30 g/l. NaCN (20 g/l) was added to the bath as a complexing agent. The Ag content of the plated Ag – Cd alloys was determined by extracting Ag as AgCl. Structural analysis was performed using an X-ray diffraction method.

7.4.2.1.2. PLATING CONDITIONS AND FILM COMPOSITION

Unfortunately precise data were not provided in the references namely Aotani (1951), Raub (1953), and Brenner (1963).


Figure 7.4.2.1.1(a) is the equilibrium phase diagram for Ag –Cd alloys, and (b) and (c) are the phase diagrams for electrolytic Ag – Cd alloy films obtained by Aotani (1951) and Raub (1953), respectively. Electrodeposited Ag –Cd alloys form the a phase in the composition range 44– 100 at.% Ag, the g phase in the range of 38 –40 at.% Ag, and the h phase in the range of 12 –38 at.% Ag. The a phase is a Ag solid solution with a fcc structure, the b phase is (AgCd) an CsCl-type intermetallic compound with a bcc structure, the g phase (Ag5Cd8) is an intermetallic with a bcc structure, the 1 phase (AgCd8) is hcp, and the h phase is a Cd solid solution with a hcp structure. Note that the composition range of each phase observed (Aotani, 1951; Raub, 1953; Brenner, 1963) in plated Ag – Cd alloy films, is slightly different from the range indicated in the equilibrium phase diagram.

384

Electrolytic Ag – Cd

385

Figure 7.4.2.1.1. (a) The equilibrium phase diagram of Ag–Cd binary alloys. The combined phase diagram and lattice constants of plated Ag–Cd alloys obtained by Aotani (1951) are shown in (b) and the phase diagram by Raub (1953) in (c).

386

Nano-Plating

REFERENCES Aotani, K. (1951) J. Jpn Inst. Metals, B15, 52. Brenner, A. (1963) Electrodeposition of Alloys, Principle and Practice, vol. I, Academic Press, New York, p. 194. Raub, E. (1953) Metalloberfla¨che, 7A, 17.

FURTHER READING Horbison, R.W. (1933) Deut. Goldschmiede-Ztg., 36, 525. Jayakrisknen, S. (2000) TIMF, 78(Part 3), 124. Mateescu, M., Prada, M. & Blejoiu, S. (2000) TIMP, 78(2), 53. Raub, E. & Wallhorst, B. (1947) Metallforsh., 2, 33. Stillwell, C.W. (1933) Metal Ind (NY), 31, 47. Stillwell, C.W. & Feinburg, H.I. (1933) J. Am. Chem. Soc., 55, 1864. Stillwell, C.W. & Stout, L.E. (1932) J. Am. Chem. Soc., 54, 2583.

7.4.2.2. Electrolytic Ag –Co


The compositions of the Ag –Co alloy-plating bath used in this study are listed in Table 7.4.2.2.1. The sources of Ag and Co were silver iodide and cobalt sulfate, respectively. Potassium iodide was added to the bath as a complexing agent to make silver iodide soluble in water. Amorphous Fe –Ni – Si – B alloy foils obtained by a splat-quenching method were used as substrates. Although this particular bath produced silver – cobalt alloy films, it generated cracks or rough surfaces in the films. The addition of 0.01 mol/l cresol sulfonic acid to the bath, however, eliminated the problem. Similar improvements were achieved by adjusting the bath pH to 3 with dilute sulfuric acid (Igarashi and Watanabe, 2000). The change in bath pH with the addition of cresol sulfonic acid is plotted in Figure 7.4.2.2.1. Cresol sulfonic acid appears to lower the pH. The composition of plated Ag – Co films was analyzed using an energy-dispersive X-ray (EDX) spectrometer attached to an SEM.


A change in the Co content of the Ag –Co films is plotted against the current density in Figure 7.4.2.2.2. Three types of baths were prepared by changing the bath pH to (a) 2.0 and (b) 3.0, or by adding (c) 0.01 mol/l cresol sulfonic acid. For all the baths, the Co content in the films was zero or very small at low current densities. At current densities above 100 A/m2, the Co content increased with increasing current density and thus it became possible to grow various compositions of Ag –Co alloy films. The relationship between the current efficiency and the Co content of the Ag –Co alloy films is shown in Figure 7.4.2.2.3. Although the bath pH affected the current efficiency,

Table 7.4.2.2.1. Bath composition for plating Ag–Co alloys. Composition

Concentration (mol/l)

AgI CoSO4·7H2O KI

0.01 0.1 1

387

388

Nano-Plating

Figure 7.4.2.2.1. A solution pH change with the addition of cresol sulfonic acid.

Figure 7.4.2.2.2. A change in the cobalt content of electrolytic Ag–Co alloys with current density.

Electrolytic Ag – Co

389

Figure 7.4.2.2.3. Relationship between current efficiency and the cobalt content of electrolytic Ag–Co alloys.

the current efficiency was generally low for the low Co content and high for the high Co content. 7.4.2.2.3. FILM MICROSTRUCTURE

X-ray diffraction patterns from different electroplated Ag –Co alloy films are shown in Figure 7.4.2.2.4. Diffraction peaks from pure Ag metal are broadened with increasing Co content, which indicates that the grain size is getting smaller with alloying. In the composition range of 0 – 60 at.% Co, no diffraction peaks from Co metal appeared. Although the Co peaks were present in the alloy composition range above 60 at.% Co, they were very weak and broad. The structure of the Co metal was hcp and no shift in the Co diffraction peaks was observed with alloying. X-ray diffraction patterns for Ag –Co alloy films plated from the bath containing 0.01 mol/l cresol sulfonic acid are shown in Figure 7.4.2.2.5. Similar to Figure 7.4.2.2.4, no shift in the diffraction peaks of both Ag and Co metals was seen. Based on these X-ray diffraction data, we calculated the lattice constants of Ag and Co metals and plotted them as a function of alloy composition in Figure 7.4.2.2.6(b). For comparison, the equilibrium phase diagram of Ag –Co alloys is displayed in Figure 7.4.2.2.6(a). This alloy system is of a mutually insoluble type at ambient temperatures. As seen in Figure 7.4.2.2.6(b), the lattice constants of Ag and Co are

390

Nano-Plating

Figure 7.4.2.2.4. X-ray diffraction patterns from various electrolytic Ag–Co alloy films. The bath used was additive-free and the pH was adjusted to 3.0 with dilute sulfuric acid.

Figure 7.4.2.2.5. X-ray diffraction patterns from various electrolytic Ag–Co alloy films grown from the bath containing 0.01 mol/l cresol sulfonic acid.


391

Figure 7.4.2.2.6. (a) The equilibrium phase diagram of Ag–Co alloys and (b) a lattice constant change in electrolytic Ag–Co alloys.

392

Nano-Plating

constant over the entire composition range, indicating that these alloys are phase-separated into pure Ag and pure Co metals. Cross-section TEM micrographs and their diffraction patterns from (A) Ag – 44 at.% Co and (B) Ag –69 at.% Co alloy films are seen in Figure 7.4.2.2.7. In the (A) film, spherical

Figure 7.4.2.2.7. Cross-section TEM micrographs and their electron diffraction patterns of electrolytic Ag– Co alloy films.


393

Figure 7.4.2.2.8. A cross-section TEM micrograph of a Ag–44 at.% Co alloy film plated at a current density of 200 A/m2, showing alloy composition values at the six sites.

crystals are dispersed throughout the film and were identified by an electron diffraction analysis as Ag crystals. High-resolution TEM observations were performed in an attempt to understand the structure of these spherical Ag crystals and the surrounding matrix (see Figures 7.4.2.2.8 and 7.4.2.2.9). In Figure 7.4.2.2.8, the spherical Ag crystals and the matrix are clearly delineated. An EDX analysis revealed that the spherical crystal was Ag – rich and the matrix region was Co-rich. The reason why we did not obtain 100 at.% Ag in the spherical crystals is that there was an object overlap within the film thickness and thus it was impossible to avoid an X-ray signal from the matrix material. The Ag – Co alloy film

394

Nano-Plating

Figure 7.4.2.2.9. A cross-section TEM micrograph of an electrolytic Ag–69 at.% Co alloy film and electron diffraction patterns from the two sites, A and B.

shown in Figure 7.4.2.2.9 was obtained from the bath containing no cresol sulfonic acid. The structure is basically the same as the one shown in Figure 7.4.2.2.8. An electron diffraction pattern from the sphere showed spotty pattern, which confirmed that these spherical crystals are Ag. An electron diffraction pattern from the matrix showed rings, indicating that it is an assembly of fine grains. An X-ray diffraction analysis was performed


395

Figure 7.4.2.2.10. Structural changes upon heating at different temperatures in a Ag– 60 at.% Co alloy deposit.

on the Ag – 60 at.% Co film after heat-treating at various temperatures (200 – 700 8C). These results are displayed in Figure 7.4.2.2.10. Diffraction peaks from Co metal were very weak from the as-deposited films or films heat-treated at low temperatures, but increased with increasing temperatures. The peaks from Ag metal also increased. The angles (or peak positions) of these diffraction peaks remained the same throughout the heat treatment. From the above observations, it can be concluded that the spherical crystals are Ag metal and the fine-grained matrix material is Co metal. It is not clear why the Co metal became fine-grained. Thus, more experiments are needed to clarify the mechanism of producing fine-grained Co metal. In all alloy compositions, few 10 nm-diameter spherical Ag crystals are imbedded inside the fine-grained Co metal matrix. Their volume ratio changed with the composition. The surface morphology of various compositions of Ag – Co alloy films obtained from the pH ¼ 3 bath is shown in Figure 7.4.2.2.11. The surface of the Ag – rich films is rough but became smooth with alloying. The surface of Co films obtained from a Ag – free bath

396

Nano-Plating

Figure 7.4.2.2.11. A surface morphology change with alloy composition in electrolytic Ag– Co alloy films.


397

Figure 7.4.2.2.12. A surface morphology change with current density and film thickness in electrolytic pure cobalt films.

(0.1 mol/l CoSO4·7H2O þ 1 mol/l KI) had severe roughness as seen in Figure 7.4.2.2.12. This roughness changed with current density and film thickness.

REFERENCE Igarashi, N. & Watanabe, T. (2000) J. Surf. Finish. Soc. Jpn, 51, 414.

7.4.2.3. Electrolytic Ag –Cu 7.4.2.3.1. PLATING METHOD

The compositions of the Ag –Cu alloy plating bath used in this study are listed in Table 7.4.2.3.1. The mixing ratio of AgCN and CuCN salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant (3 g/l). A graphite sheet was used as an anode. The plating was conducted at 100 A/m2 on a polycrystalline Fe sheet in a 20 8C bath for 20 s. The bath was agitated with a magnetic stirrer during plating. Thin, electron-transparent, Au films were also prepared as a substrate, so that one can observe the initial stages of Ag – Cu alloy electrodeposits in situ using TEM (see the preparation method in Figure 6.10 of Chapter 6). For plating the thick alloy films, polycrystalline Fe sheets were used as substrates. Here the alloy film was stripped from the Fe substrate in 0.2% HCl and then subjected to a detailed structural analysis using TEM, electron diffraction and high-resolution reflection electron diffraction methods. An atomic absorption method was used to determine the composition after dissolving the alloy films in nitric acid.


The relationship between the Cu concentration in the bath and the Cu content in the plated Ag – Cu alloy films is illustrated in Figure 7.4.2.3.1. The plating was conducted at 150 A/m2 for 10 and 180 s. The film composition appears to depend markedly on the film thickness. This phenomenon is often observed in alloy plating systems that contain metal ions with a large difference in deposition potential. In the initial stage of

Table 7.4.2.3.1. Bath composition for plating Ag– Cu alloys. Composition (g/l)

AgCN CuCN KCN K2CO3

Bath name A

B

C

D

E

F

G

H

I

2.33 1.69

1.86 2.11

1.62 2.39

1.12 2.96

0.93 3.17 10 5

0.56 3.59

0.18 4.03

0.07 4.41

0.04 4.19

398

Electrolytic Ag – Cu

399

Figure 7.4.2.3.1. Relationship between the Cu content of plated Ag–Cu alloy films and the Cu concentration of the plating baths. The current density was 100 A/m2 and the bath temperature was 20 8C.

the Ag – Cu alloy deposition, a larger amount of noble metal (Ag) tends to deposit preferentially, forming Ag – rich alloys known as the initial layer (Watanabe, 1990). Thus, the composition of the thin alloy film plated for 10 s is expected to be Ag – rich as seen in Figure 7.4.2.3.1.


High-resolution reflection electron diffraction patterns from various compositions of plated Ag –Cu alloy films are shown in Figure 7.4.2.3.2. Diffraction lines from pure Ag and Cu metals are indexed in the figure. The lattice constant of the Ag film was 0.409 nm, which was slightly larger than the value (0.4086 nm) quoted in the JCPDS file (JCPDS

400

Nano-Plating

Figure 7.4.2.3.2. High-resolution reflection electron diffraction patterns from plated Ag, Ag–Cu alloy, and Cu films.

(Ag)). Similarly, the lattice constant of the Cu film was 0.362 nm, which was also slightly larger than the JCPDS value (0.36150 nm) (JCPDS (Cu)). Lattice constants for the other alloy compositions were also measured from the diffraction patterns and are plotted in Figure 7.4.2.3.3(b) over the entire composition range. The equilibrium phase diagram of Ag – Cu alloys is of a typical eutectic type (Massalski, 1990) as seen in Figure 7.4.2.3.3(a). Up to the alloy composition of 21 at.% Cu, diffraction rings from Cu metal (or Cu-rich


401

Figure 7.4.2.3.3. (a) The equilibrium binary phase diagram of Ag–Cu alloys and (b) the lattice constants of Ag–Cu alloy films obtained by an electroplating method.

402

Nano-Plating

alloy) did not appear. Above 90 at.% Cu, diffraction rings from Ag metal (or Ag – rich alloy) were not present. In the composition range of 21 –90 at.% Cu, diffraction rings from both Ag and Cu metals were present, indicating the presence of a two-phase alloy containing Ag – rich (a) and Cu-rich (b) phases. The absence of diffraction rings from Cu and Ag metals in plated Ag – Cu alloy films below 21 at.% Cu and above 90 at.% Cu was attributed to the formation of supersaturated Ag – Cu alloys in the Ag – rich and Cu-rich sides, although it is also believed to be possible that the diffraction intensity was reduced by the refined grain size. According to Aotani (1951), the formation of a supersaturated alloy is possible up to 33 at.% Cu in the Ag side and up to 15 at.% Ag in the Cu side. Fedotev and Vyacheslavov (1970) reported 20 at.% Cu for the former and 20 at.% Ag for the latter. Kimata and Nishi (1967a) reported that the two phases were present in the composition range of 30 –97 at.% Cu but did not discuss the formation of supersaturated alloys (Kimata et al., 1967a,b). We have emphasized throughout this book that the structure of plated metals and alloys is analogous to the structure of a solid obtained by splat-quenching from high temperatures. According to this theory, the formation of supersaturated alloys in plated Ag – Cu alloys should be possible up to 14.1 at.% Cu in the Ag side and up to 9.51 at.% Ag in the Cu side. The grain size of these plated Ag – Cu films can be estimated qualitatively from the width of the reflection electron diffraction rings shown in Figure7.4.2.3.2. Diffraction rings in the Ag and Cu sides were sharp, whereas rings from the alloy films having the composition close to the 50:50 ratio were very broad. The grain size apparently decreased with increased alloying. Figure 7.4.2.3.4 contains bright-field (BF)/dark-field (DF) TEM micrographs and their associated electron diffraction pattern, which show the structure of a plated Ag – Cu alloy film, in which both Ag and Cu crystals coexist. Here both the Ag and Cu crystals appear to lattice-match to the thin {001} Au substrate as indicated by the electron diffraction pattern. DF images taken from the (200) Ag and the (200) Cu spots are seen in Figure 7.4.2.3.4(b) and (c), respectively. From these images, the film consists of mixed phases of 5 – 20 nm Ag and Cu crystals, which have grown epitaxially on the Au substrate. Another example for simultaneous observations of deposit/substrate composites is presented in Figure 7.4.2.3.5, which shows BF TEM micrographs and electron diffraction patterns for plated Ag – 15 at.% Cu, Ag – 70 at.% Cu, and Ag – 90 at.% Cu alloy films grown on a thin {110} Au substrate. For all alloy compositions, the films appear to have grown epitaxially, although they do contain a high density of defects. In the electron diffraction patterns, diffraction spots from both the deposit and the Au substrate are visible, although these patterns are complicated by the presence of double diffraction spots. The effect of 1-h annealing at 400 8C in a vacuum on the structure of plated Ag – 20 at.% Cu and Ag – 90 at.% Cu alloy films is seen in Figure 7.4.2.3.6. Before annealing, only the diffraction spots from the Ag – and Cu-rich solid solution alloys were present. After


403

Figure 7.4.2.3.4. (a) A plan-view TEM micrograph and its electron diffraction pattern from an electrolytic Au– Cu alloy film grown on a {100} Au film substrate. Micrographs (b) and (c) represent dark-field images taken from satellite spots, A (a phase) and B (b phase), present around the Au {002} substrate spot. Sites marked with d and d0 are identical locations in the specimen. Similarly, sites marked with p, p0 , and p00 are identical.

annealing, however, diffraction spots from the pure Ag and Cu metal appeared and the crystal size increased to 200 nm. The surface morphology of various compositions of plated Ag –Cu alloy films grown for 3 min (except for (e), 30 min) is shown in Figure 7.4.2.3.7. The films from the high Ag concentration baths formed dendrites but those from the high Cu concentration baths produced a smooth surface.

404 Nano-Plating Figure 7.4.2.3.5. Plan-view TEM micrographs and their electron diffraction patterns from composite films of plated Ag–Cu alloy/thin single-crystal Au. (a) An electrolytic Ag–15 at.% Cu alloy film grown on a {110} single-crystal Au substrate, (b) an electrolytic Ag–70 at.% Cu alloy film grown on a {111} singlecrystal Au substrate, and (c) an electrolytic Ag–90 at.% Cu alloy film grown on a {110} single-crystal Au substrate. All the Ag–Cu alloy films were deposited for 10 s.


405

Figure 7.4.2.3.6. The effect of 1-h vacuum annealing at 400 8C on the structure of electrolytic Ag–Cu alloy films. The structure of a Ag–20 at.% Cu alloy film (a1) before and (a2) after the annealing. The structure of a Ag– 90 at.% Cu alloy film (b1) before and (b2) after the annealing.

Figure 7.4.2.3.7. Surface morphologies in various Ag–Cu alloy films electroplated for 3 min except for the film shown in (a), which was plated for 30 min. (a) Ag–17 at.% Cu, (b) Ag–26 at.% Cu, (c) Ag–33 at.% Cu, (d) Ag– 50 at.% Cu, (e) Ag–84 at.% Cu, and (f) Ag–90 at.% Cu.

406

Nano-Plating

REFERENCES Aotani, K. (1951) J. Jpn Inst. Metals, B15, 52. Fedotev, N.P. & Vyacheslavov, P.M. (1970) Plating, 57, 700. JCPDS File #4-786 (Ag). JCPDS File #4-836 (Cu). Kimata, T. & Nishi, S. (1967a) J. Metal Finish. Soc. Jpn, 18, 268. Kimata, T. & Nishi, S. (1967b) J. Metal Finish. Soc. Jpn, 18, 293. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, pp. 94 – 97. Tanabe, Y., Kakekawa, M. & Shimizu, Y. (1976) J. Mater. Soc. Jpn, 13, 255. Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 652.

7.4.2.4. Electrolytic Ag –Sn 7.4.2.4.1. PLATING METHOD

The compositions of the Ag – Sn alloy plating bath used in this study are listed in Table 7.4.2.4.1. To obtain various Ag. and Sn alloy compositions, two types of baths were prepared. The plating was carried out at current densities of 10– 100 A/m2 from roomtemperature (25 8C) baths, without agitation, until 130 C was consumed. The film composition was analyzed using inductively coupled plasma (ICP) atomic emission spectroscopy.


The Ag content of Ag – Sn alloy films grown from the two baths, A and B, is plotted in Figure 7.4.2.4.1 as a function of current density. The Ag content is high at low current densities but decreases with increasing current density. Conversely, the Sn content is low at low current densities but increases with increasing current density. We could not obtain Ag – Sn alloy films, with a Sn content less than 25 at.%. Consequently, we investigated the structure of plated Ag –Sn alloy films containing in excess of 25 at.% Sn.


Figure 7.4.2.4.2 displays X-ray diffraction patterns for plated Ag – Sn alloy films with various alloy compositions. All the films exhibited strong diffraction peaks, indicating the presence of crystalline phases. The films with a low Ag content displayed strong peaks from the b-Sn phase as well as from intermetallic compounds (1 or z phase JCPDS File #4-800; JCPDS File #29-1151). With increasing Ag content, however, diffraction peaks

Table 7.4.2.4.1. Bath composition for plating Ag–Sn alloys (g/l). Composition

A

B

K4P2O7 KI SnCl2·2H2O AgI

0.54 2.0 0.18 0.02

0.54 2.0 0.195 0.005

407

408

Nano-Plating

Figure 7.4.2.4.1. The Ag content (at.%) of electrolytic Ag– Sn alloy films versus current density (A/m2).

Figure 7.4.2.4.2. X-ray diffraction patterns from electrolytic Ag–Sn alloys having various Ag contents.

Electrolytic Ag –Sn

409

Figure 7.4.2.4.3. (a) The equilibrium phase diagram of Ag–Sn alloys and (b) the phase diagram of plated Ag–Sn alloys.

from the b-Sn phase decreased, accompanied by an increase in the diffraction peaks of the 1 or z phase. For the 75 at.% Ag alloy film, the peaks of the b-Sn finally disappeared and only the 1 or z phase peaks remained.

410

Nano-Plating

Figure 7.4.2.4.4. The (111) lattice spacing of the 1 phase and the (101) lattice spacing of the z phase in electrolytic Ag–Sn alloy films.

As seen in the equilibrium phase diagram in Figure 7.4.2.4.3(a) (Massalski, 1990), the Ag – Sn binary alloy system is a peritectic type that forms two kinds of intermetallic compounds. It is difficult to determine which diffraction peaks belong to the 1 or z phase because of the proximity of their lattice spacing. For this phase determination, the strongest unknown peak, whose diffraction angle (near 408) was close to that of the intermetallic compounds (the 1 or z phase), was chosen and plotted against Ag content as seen in Figure 7.4.2.4.4, where the diffraction angle was converted to the lattice spacing. The horizontal line at 0.2280 nm represents the (111) and (021) lattice spacing of the 1 phase (Ag3Sn), which is listed in the JCPDS file. Similarly, the line at 0.2263 nm represents the (101) spacing of the z phase (Ag4Sn). Near the alloy composition of 20 at.% Ag, the lattice spacing was close to that of the 1 phase but with a further increase in Ag content, the lattice spacing approached that of the z phase. This indicates that with increasing Ag content, the lattice spacing changed continuously from that of the 1 phase to that of the z phase, or the mixing ratio of the 1 and z phases varied successively with the Ag content. However, as the lattice spacing of the (111) in the 1 phase is very close to that of the (021), more work is necessary to fully determine the phases in the plated Ag – Sn alloy films. The phases observed in plated Ag –Cu alloys are summarized in Figure 7.4.2.4.3(b). For the film with 25 at.% Sn, only the 1 phase (z phase) was detected and no b-Sn was found.


411

Figure 7.4.2.4.5. Cross-section TEM micrographs and their electron diffraction patterns for different plated Ag– Sn alloy films.

412

Nano-Plating

Figure 7.4.2.4.6. Part of the image seen in Figure 7.4.2.4.5(c) was magnified in (a). EDX spectra taken from two locations, (b) and (c), in (a) are shown in Figure 7.4.2.4.6 (b0 ) and (c0 ), respectively.

Tanabe et al. (1983) used a cyanide-type bath for plating Ag – Sn alloy films and found that the Ag-phase, z phase, and b-Sn grew lattice-matched on Fe substrates. These phases were also found to lattice-match to each other. These observed phases, however, were different from those listed in Figure 7.4.2.4.3(b). Cross-section TEM specimens of plated Ag –Sn alloy films with various compositions were prepared using a microtome technique. TEM micrographs for the microtomed samples are seen in Figure 7.4.2.4.5 together with the electron diffraction patterns. It is well known that mechanical damage can be introduced into TEM samples during microtoming. Specially, the film containing 75 at.% Ag was very brittle and thus was easily lost during preparation. None of the films exhibited a layer structure, which is often formed by compositional or structural fluctuations along the film thickness direction. Electron diffraction patterns yielded results similar to the X-ray results. Part of the image in Figure 7.4.2.4.5(c) was magnified in Figure 7.4.2.4.6. Here EDX spectra were taken from two locations, (b) and (c). The results are shown in Figure 7.4.2.4.6(b0 ) and (c0 ), respectively. Note that the Cu peaks originated from the Cu mesh used to position the microtomed TEM samples. The composition at (b) was 1 at.% Ag –Sn, whereas the composition at (c) was 70 at.% Ag –Sn. From this analysis, the former can be identified as


413

Figure 7.4.2.4.7. SEM micrographs showing the surface morphologies of electrolytic Ag– Sn alloy films with various compositions.

b-Sn and the latter as the 1 phase or the z phase. These phases were distributed uniformly in the alloy films. Figure 7.4.2.4.7 contains SEM micrographs showing the surface morphologies of plated Ag – Sn alloy films with various compositions. The surface appears to be rough for the Agrich and Sn-rich alloys but becomes smooth near the alloy composition of 10– 20 at.% Ag. In Figure 7.4.2.4.8, we show (a) a bright-field image of an electrolytic Ag – Sn film, in which the a, b, and z phases coexist, (b) its electron diffraction pattern, (c) a dark-field

414

Nano-Plating

Figure 7.4.2.4.8. (a) A bright-field image of an electrolytic Ag–Sn film, in which the a, b, and z phases coexist, (b) its electron diffraction pattern, (c) a dark-field (DF) image of the a/z phases, and (d) a DF image of the z phase.

Figure 7.4.2.4.9. Electron diffraction patterns from electrolytic Ag– Sn alloy films having the compositions of (a) 1.0 (b) 58.2 and (c) 78.5 at.% Sn.


415

(DF) image of the a/z phases, and (d) a DF image of the z phase. The grain size and distribution of each phase can be seen in the TEM micrographs Tanabe et al. (1983). Figure 7.4.2.4.9 displays electron diffraction patterns for electrolytic Ag – Sn alloy films grown on a single-crystal copper. The composition of Sn in the alloys were (a) 1.0 at.% Sn, (b) 58.2 at.% Sn, and (c) 78.5 at.% Sn. All the alloy films have a lattice matching relationship to the copper substrate.

REFERENCES Arai, S. & Watanabe, T. (1996) J. Jpn. Inst. Metals, 60, 1149. Tanabe, Y., Hasegawa, N. & Odaka, M. (1983) J.Metal Finish. Soc. Jpn, 34, 456.

FURTHER READING Harada, M. & Satoh, R. (1990) IEEE Trans. CPMT, 13, 736. JCPDS File #4-800. JCPDS File #29-1151. Kawaguchi, T. & Kojima, M (1996a). The Second Symposium on Microjoining and Assembly in Electronics (February). p. 180. Kawaguchi, T. & Kojima, M (1996b). The Second Symposium on Microjoining and Assembly in Electronics (February). p. 186. Kubota, N., Horikoshi, T. & Sato, E. (1983) J. Metal Finish. Soc. Jpn, 34, 217. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, p. 402. Matsushita, S. (1971) J. Metal Finish. Soc. Jpn, 22, 60. Puippe, J.-C. & Fluehmann, W.F. (1983) Plat. Surf. Finish., 70, 46. Suganuma, K. & Nakamura, Y. (1995) J. Jpn. Inst. Metals, 59, 1299. Takao, H., Hasegawa, H., Tsukada, T., Mizuno, M., Yamada, K. & Yamamoto, S (1996). The Second Symposium on Microjoining and Assembly in Electronics (February), p. 104. Takemoto, T., Hiraishi, M., Miyazaki, M., Matsunawa, A. & Ohsawa, Y (1996). The Second Symposium on Microjoining and Assembly in Electronics (February), p. 239.

7.4.2.5. Electrolytic Ag –Zn


The structure of electrolytic Ag – Zn alloy films have been described by Aotani (1951) and Brenner (1963). These results are summarized in Figure 7.4.2.5.1 together with the equilibrium phase diagram of Ag – Zn alloys.

Figure 7.4.2.5.1. (a) The equilibrium phase diagram of Ag–Zn alloys and (b) the phase diagram of electrolytic Ag–Zn alloy films.

REFERENCES Aotani, K. (1951) J. Jpn Inst. Metals, B15, 52. Brenner, A. (1963) Electrodeposition of Alloys, Principle and Practice, vol. I, Academic Press, New York, p. 194.

416

7.4.2.6.

Electrolytic Al– Mn


Al –Mn alloy deposits can be obtained in a fuse-salt bath. The following steps were taken to prepare the plating bath. A mixture of AlCl3, NaCl, and KCl, whose mixing mass ratio was 77:14:9, was first put into a plating cell and then melted by heating while in a nitrogen atmosphere. The removal of impurities was achieved by adding 5 mass % Al powder to the mixture, which was then stored at 200 8C for 3– 7 days. The solution was colorless and transparent. NaCl and KCl were dried under reduced pressure at 350 8C for 48 h, but anhydrous AlCl3 was used directly without any treatment. According to Karl Fischer titration, the initial oxygen ion concentration, which accounts for water content in the solution, was 0.2 mass %. Super-high grade MnCl4·H2O was dried under reduced pressure at 259 8C for 24 h. The desired amount of the anhydrous MnCl4 salt was added to the bath as a source of the alloying element. A cold-rolled steel sheet was used as the substrate and a 99.8% pure Al plate was employed as an anode. For further details of this plating method, readers should refer to the original papers (Grushko and Stafford, 1990; Uchida et al., 1991).


Figure 7.4.2.6.1 is a graph that plots an experimental curve defining a boundary between smooth and powdery Al –Mn alloy deposits as a function of current density and solution flow rate. The deposits tend to become powdery at low solution flow rates, high current densities, and low Mn concentration. At high solution flow rates, low current densities, and high Mn concentrations, films with smooth surfaces were obtained. Figure 7.4.2.6.2 shows a relationship between the Mn content in the plated films and the concentration of Mn ions in the bath. The Mn content in the plated films appears to increase with increasing Mn concentration in the bath regardless of the plating conditions. The Mn content in the plated films was also plotted against the current density, the solution flow rate, the amount of AlCl3, and the bath temperature in Figure 7.4.2.6.3(a) – (d), respectively. A change in alloy composition with bath temperature and current density in electroplated Al –Mn alloy films is shown in Figures 7.4.2.6.4 and 7.4.2.6.5. This was determined by Grushko and Stafford (1990). 417

418

Nano-Plating

Figure 7.4.2.6.1. A graph illustrating how the Mn2þ concentration and the solution flow rate affect the magnitude of current density for obtaining smooth Al–Mn alloy deposits.

Figure 7.4.2.6.2. The Mn content of plated Al–Mn alloy films as a function of the Mn2þ ions concentration of the bath.

Electrolytic Al – Mn

419

Figure 7.4.2.6.3. The Mn content of plated Al– Mn alloy films plotted against (a) the current density, (b) the solution flow rate, (c) the amount of AlCl3, and (d) the bath temperature.

Figure 7.4.2.6.4. A change in alloy composition with bath temperature in electroplated Al–Mn alloy films.

420

Nano-Plating

Figure 7.4.2.6.5. A change in alloy composition with current density in electroplated Al– Mn alloy films.


X-ray diffraction patterns from electrodeposited Al– Mn films grown under various plating conditions are shown in Figure 7.4.2.6.6. For the Al –Mn alloys containing low Mn content, strong {111} and {200}diffraction peaks from Al crystals appeared, but these pattern broadened for the 25– 40 mass % Mn alloys, suggesting that the films became amorphous. With a further increase in Mn content, the (202) peak of an orthorhombic Al6Mn crystal and the {035} peak of Al2Mn crystals appeared. Highresolution lattice imaging in combination with the conventional TEM technique indicated that an amorphous phase is already locally present, even for the 15 mass % Mn alloy films. Figure 7.4.2.6.7(a) is the equilibrium phase diagram (Massalski, 1990) of Al –Mn alloys over the Mn composition range of 1– 40 at.%. In Figure 7.4.2.6.7(b), the composition range of various phases found in electroplated Al – Mn alloy deposits is plotted as a function of bath temperature (Grushko and Stafford, 1990). In addition, Grushko and Stafford (1990) discussed the structure of these plated alloy films using TEM micrographs and electron diffraction patterns.

Electrolytic Al – Mn

421

Figure 7.4.2.6.6. X-ray diffraction patterns from electrodeposited Al – Mn films having various alloy compositions grown under different plating conditions.

422

Nano-Plating

Figure 7.4.2.6.7. (a) The equilibrium phase diagram of Al–Mn alloys over the limited Mn composition range of 1 , 40 at:% and (b) the phase diagram of electroplated Al –Mn alloys plotted as a function of bath temperature.

REFERENCES Grushko, B. & Stafford, G.R. (1990) Met. Trans. A, 21A, 2869. Uchida, J., Tsuda, T., Yamamoto, Y., Seto, H., Abe, K. & Shibuya, A. (1991) Tetsu-to-Hagane, 77, 931. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, p. 171.

7.4.2.7.

Electrolytic Au–Cu (Kawai et al., 1968a,b)


The compositions of the Au – Cu alloy plating bath used in this study are listed in Table 7.4.2.7.1 (Kawai et al., 1968b). KCN was added at 7 g/l to all the baths. A Ni sheet was used as the substrate. The composition of the alloy deposits was determined by both chemical and X-ray fluorescence methods. For the chemical analysis, a Cu sheet was used as the substrate. Here the plated films were dissolved in an aqua regia solution, which was subsequently dried. The dried residue was then dissolved in warm water, followed by an Au reduction using a thiosulfurous acid solution. After the reduced Au was filtered and rinsed in water, it was burned to Au metal, and then weighed for an alloy composition analysis. For the X-ray fluorescence analysis, a Ni sheet was chosen as the substrate. The alloy composition by the chemical method agreed with the composition by the X-ray fluorescence method within an accuracy of 5%. Cu was a target for an X-ray diffractometer used in this study. The surface morphology was studied using a TEM replica technique.


In Figure 7.4.2.7.1, the Au content of plated Au –Cu alloy films was plotted against the Au concentration in the bath, together with the effect of the CN/Cu ratio in the bath.


The equilibrium phase diagram of Au – Cu binary alloys (Massalski, 1990) is shown in Figure 7.4.2.7.2(a). According to the phase diagram, the Au – Cu alloy system exhibits mutual solubility over the entire composition range at high temperatures, but form intermetallic compounds such as Cu3Au, CuAu, and CuAu3 at temperatures below 400 8C. In Figure 7.4.2.7.2(b), we plot lattice constants against the alloy composition for electroplated Au – Cu alloy films. This data indicates that only the Au crystal appeared in the Au-rich region. Similarly, only the Cu crystal appeared in the Cu-rich region, except for the Au-67.6 at.% Cu alloy, where both Au and Cu crystals coexisted. An increase in the Au content in the Cu-rich side did not change the lattice constant. For the Au-rich alloys, the lattice constant decreased slightly with increasing Cu content. According to Raub and Sautter (1956, 1963), if the bath pH is adjusted to 7.0– 7.6, the lattice constant of Au – Cu alloy films decreases gradually from 0.407 nm for the pure Cu to 0.358 nm for the alloy 423

424

Nano-Plating

Table 7.4.2.7.1. The composition of Au– Cu alloy plating baths. Gold concentration in plating bath (Au/(Au þ Cu)) (%)

KAu(CN)2 (g/l)

K2Cu(CN)3 (g/l)

KCN CN/Cu

Free KCN (g/l)

50

7.5

17.5

3.00 3.25 3.50

0 1.3 2.6

40

7.5

26.0

3.00 3.25 3.50

0 2.0 3.8

30

7.5

40.5

3.00 3.25 3.50

0 3.1 5.6

20

7.5

69.1

3.00 3.25 3.50

0 5.1 10.2

10

7.5

155.6

3.00 3.25 3.50

0 11.5 23.0

Figure 7.4.2.7.1. The Au content of plated Au–Cu alloy films plotted as a function of the Au concentration in the bath and the CN/Cu ratio in the bath.

Electrolytic Au – Cu

425

Figure 7.4.2.7.2. (a) The equilibrium phase diagram of Au–Cu binary alloys and (b) the measured lattice constants of electroplated Au–Cu alloy films versus the alloy composition.

426

Nano-Plating

films containing 30 at.% Au. Furthermore, they reported that the addition of 10 g/l sodium sulfate to the bath changes the lattice constant of Cu with Au alloying.

REFERENCES Kawai, K., Kuroda, T. & Ishiguro, I. (1968a) J. Metal Finish. Soc. Jpn, 19, 6. Kawai, K., Kuroda, T. & Ishiguro, I. (1968b) J. Metal Finish. Soc. Jpn, 19, 356.

FURTHER READING Massalski, T.B. (1990). Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, p. 362– 368. Raub, E. & Sautter, F. (1956) Metalloberfla¨che, 10, 65. Raub, E. & Sautter, F. (1963) Metalloberfla¨che, 17, 17.

7.4.2.8.

Electrolytic Au– Ni (Shimizu, 1976; Shimizu et al.,

1976) 7.4.2.8.1. PLATING METHOD

The compositions of the Au – Ni alloy plating bath used in this study are listed in Table 7.4.2.8.1. Various proportions of chemicals (HAuCl4, NiCl2·6H2O, KCN, KOH, Na3C6H5O7·2H2O, C6H8O7·H2O) were mixed to prepare 10 baths. The concentration of NiCl2·6H2O was changed to obtain Au –Ni alloy deposits with different compositions, while keeping the concentration of HAuCl4 constant (1.0 g/l) (see Table 7.4.2.8.1). The plating was carried out on polycrystalline or single-crystal Cu substrates at a current density of 100 A/m2, in a 50 8C bath, using a graphite plate as an anode. The bath pH was adjusted to 3 –3.5 by adding KOH. To determine the alloy composition, the Au – Ni alloy deposits were first dissolved in aqua regia, and then analyzed using an atomic absorption method.


The Ni content of a plated Au – Ni alloy film against the Ni concentration of the bath is plotted in Figure 7.4.2.8.1, where the plating time was chosen to be (a) 40 and (b) 360 s. The plating time (or the film thickness) appears to affect the alloy composition significantly. The thickness dependence of the alloy composition is a well-known phenomenon that occurs in some binary alloy plating systems, where two metal ions have markedly different deposition potentials. In such systems, the nobler metal ions deposit preferentially at the beginning of the plating and form a so-called initial layer. Thus, the composition of the initial layer is richer in the nobler metal than the composition of the subsequently deposited alloy layer, which has a constant bulk composition. The result shown in Figure 7.4.2.8.1 reflects the effect of solution agitation, where solution agitation promoted deposition of the nobler Au metal.


Figure 7.4.2.8.2(a) is the equilibrium phase diagram of Au – Ni alloys (Massalski, 1990), which indicates mutual solubility over the whole composition range at high temperatures, but has a miscibility gap at low temperatures. To determine the alloy composition of plated Au – Ni alloy films, we analyzed the X-ray diffraction patterns as seen in Figure 7.4.2.8.3. In addition, electron diffraction measurements were taken from 10 different Au – Ni alloy films grown on {001} Cu substrates (see Figure 7.4.2.8.4). In the alloy deposits containing 427

428

Nano-Plating

Table 7.4.2.8.1. The composition of Au– Ni alloy plating baths. Composition (g/l)

HAuCl4 NiCl2·6H2O KCN KOH Na3C6H5O7·2H2O C6H8O7·H2O

Bath name A

B

C

D

E

F

G

H

I

J

1.0 5 2.0 – 100 100

1.0 15 2.0 – 100 100

1.0 30 2.0 – 100 100

1.0 60 2.0 10 100 100

1.0 100 2.0 20 100 100

1.0 150 2.0 30 100 100

1.0 200 2.0 40 100 100

1.0 300 2.0 60 100 100

0.5 300 1.0 60 100 100

0.3 300 0.5 60 100 100

less than 40 at.% Ni, Au crystals appeared and their lattice constant (JCPDS File #4-0784) decreased with increasing Ni content. This trend suggests that Ni dissolved substitutionally into the Au host metal and formed a solid-solution alloy up to 40 at.% Ni. According to the equilibrium phase diagram, however, 40 at.% Ni cannot be dissolved in Au at ambient temperatures (see Figure 7.4.2.8.2(a)). Therefore, these alloys were super-saturated with Ni and are considered to be a meta-stable phase, which will be called a1 (M). In the composition range of 52 – 94 at.% Ni, we discovered another meta-stable phase (AuNi3), which was also observed by Fedotev et al. (1967) in Au – Ni films plated from a pyrophosphate –cyanide bath. This meta-stable phase will be called Au –Ni (M). For the alloys with more than 76 at.% Ni, the X-ray diffraction patterns were found to originate from the Ni crystals. The lattice constant of the Ni crystals (JCPDS File #4-0850) increased with increasing Au content. These Ni crystals were also super-saturated solid-solution alloys, which will be called a2 (M). A change in the lattice constants of these phases is plotted in Figure 7.4.2.8.2(b) as a function of alloy composition.

Figure 7.4.2.8.1. The Ni content of plated Au–Ni alloy films plotted as a function of the Ni concentration of the bath.

Electrolytic Au – Ni

429

Figure 7.4.2.8.2. (a) The equilibrium phase diagram of Au–Ni alloys and (b) the lattice constants of plated Au– Ni alloy deposits versus the alloy composition.

430

Nano-Plating

Figure 7.4.2.8.3. X-ray diffraction patterns from various compositions of plated Au–Ni alloy films.


Figure 7.4.2.8.4. Electron diffraction patterns from various compositions of plated Au–Ni alloy films grown on {001} Cu substrates.

431

432

Nano-Plating

The gold – nickel binary alloy system has a miscibility gap at ambient temperatures and thus plated Au – Ni alloy films should have formed two phases of Au- and Ni-rich alloys. As demonstrated in Figure 7.4.2.8.2(b), the films developed supersaturated solid-solution alloys by extending the solubility limit, resulting in the formation of three meta-stable phases, a1 (M), Au – Ni (M), and a2 (M). Super-saturated Au – Ni alloys similar to the a1 (M) phase were obtained using a quenching technique from the liquid state by Koster and Dannhl (1936), who reported that the lattice constants decreased with increasing Ni content, while obeying the Vegard law. Although the lattice constant change in the present experiment was slower than the change obtained by Koster and Dannhl (1936), it was consistent with results by Raub (1953) and Kawai (1968). When the Au – Ni (M) started appearing, the lattice constant of the a1 (M) phase no longer changed with increasing Ni. Conversely, the lattice constant of the Au – Ni (M) phase, decreased from 0.38 to 0.372 nm with increasing Ni. Finally, the lattice constant of the a2 (M) phase increased from 0.35 to 0.358 nm with increasing Au. Fedotev et al. (1967) estimated the structure of the meta-stable phase, Au – Ni (M), as AuNi3 because the composition range of the phase formation was 23– 26 at.% Au. Sanadze and Gulyaev (1960a,b) reported that several meta-stable phases in bulk Au –Ni alloys, obtained by metallurgical processes involving melting and solidification, were formed upon annealing. They observed three phases with an ordered structure, AuNi3, AuNi, and AuNi3. The lattice constants were a ¼ 0:3948 nm for AuNi3, a ¼ 0:3829 nm and c ¼ 0:3712 nm for AuNi, and a ¼ 0:3712 nm for AuNi3. The lattice constants of the Au – Ni (M) phase observed in the present study appears to be close to the constants of the AuNi phase, whereas the lattice constants in the Ni-rich side are close to that of AuNi3. The structure of the meta-stable Au –Ni (M) phase is thus considered to take a structural form similar to AuNi and AuNi3. Figure 7.4.2.8.5 contains TEM micrographs showing the initial stages of Au – 76 at.% Ni alloy films plated on single-crystal {001} Cu substrates for (a) 0.5 and (b) 1 s. Islandshaped crystals with dimension of 5– 10 nm are distributed randomly over the substrate surface. Electron diffraction patterns from these samples showed spotty rings, which indicate that these islands were randomly oriented. It is apparent that the islands did not form a strong metallic bond to the substrate and thus were not latticed-matched. After a 1-s plating, the number of islands increased, covering the whole surface, and at the same time started lattice-matching to the substrate as indicated by the development of {001} singlecrystal electron diffraction patterns. These patterns contained diffraction spots from both the a1 (M) and the Au – Ni (M) phases. The islands seen in the bright-field image of Figure 7.4.2.8.5(b) represent the crystal size. Figure 7.4.2.8.6(a) is a bright-field TEM micrograph showing Au – 65 at.% Ni alloy films plated on a single-crystal {110} Cu substrate. Figure 7.4.2.8.6(b) and (c) are dark-field images taken from the diffraction spots of the a1 (M) and the Au – Ni (M) phases, respectively. Both the a1 (M) and the Au – Ni (M) crystals are granular with


433

Figure 7.4.2.8.5. TEM micrographs showing the initial stages of Au–65 at.% Ni alloy films plated on singlecrystal {001} Cu substrates for (a) 0.5 and (b) 1 s.

Figure 7.4.2.8.6. Bright-field (BF) and dark-field (DF) TEM micrographs showing the structure of Au– 65 at.% Ni alloy films plated on single-crystal {011} Cu substrates.

434

Nano-Plating

size ranging from 5 to 200 nm. These samples exhibited a {001} single-crystal electron diffraction pattern, indicating that they are latticed-matched to the substrate. The effect of bath temperature ((a) 50 8C and (b) 2 4 8C) on the structure of Au76 at.% Ni alloy films grown on single-crystal {001} Cu substrates is shown in Figure 7.4.2.8.7. Dark-field images ((a0 ) and (b0 )) were taken from the circled spots in the electron diffraction patterns. An electron diffraction pattern from the film plated in the 50 8C bath shows sharp spots, whereas the film grown in the 2 4 8C bath displays broad rings. All three meta-stable phases, a1 (M), Au –Ni (M), and a2 (M) can be easily identified from electron diffraction patterns of the films plated in the 50 8C bath, but not from the 2 4 8C bath. The bath temperature therefore affects the size of sub-grains and their lattice-matching behavior. Figure 7.4.2.8.8 is a TEM micrograph showing the effect of annealing on the structure of Au –56 at.% Ni alloy films plated from the 2 4 8C bath. A heating

Figure 7.4.2.8.7. The effect of the bath temperature ((a) 50 8C and (b) 24 8C) on the structure of Au–76 at.% Ni and Au–78 at.% Ni alloy films grown on single-crystal {001} Cu substrates. (a) and (b) are BF images, whereas (a0 ) and (b0 ) are the corresponding DF images.


435

Figure 7.4.2.8.8. A TEM micrograph showing the effect of annealing on the grain growth of Au–56 at.% Ni alloy films plated on single-crystal {001} Cu substrates.

436

Nano-Plating

Figure 7.4.2.8.9. SEM micrographs showing the surface morphologies of Au–Ni alloy films with various compositions.

experiment was conducted inside a TEM on the alloy film stripped off the substrate. The sub-grains appear to increase upon heating, accompanied by a sharpening of the electron diffraction pattern. Phase separation became obvious with annealing. Figure 7.4.2.8.9 contains SEM micrographs that show the surface morphologies of Au – Ni alloy films with various compositions. All the films displayed smooth surfaces, although some films contained nodules.

REFERENCES Fedotev, N.P., Vyacheslavov, P.M., Lokshtanova, O.G. & Kruglova, E.G. (1967) J. Appl. Chem. USSR, 40, 2167. JCPDS File #4-0784. JCPDS File #4-0850. Kawai, K. (1968) J. Metal Finish. Soc. Jpn, 19, 487. Koster, W. & Dannhl, W. (1936) Z. Metallkunde, 28, 248. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, p. 402. Raub, E. (1953) Metalloberfla¨che, 7, 17. Sanadze, V.V. & Gulyaev, G.V. (1960a) Soviet Phys. Cryst., 4, 464. Sanadze, V.V. & Gulyaev, G.V. (1960b) Soviet Phys. Cryst., 4, 496. Shimizu, Y (1976) PhD. dissertation, A TEM study on the crystal growth and fine structure of electrolytic binary alloys Tokyo Metropolitan University. Shimizu, Y. & Tanabe, Y. (1976) J. Metal Finish. Soc. Jpn, 27, 20.

7.4.2.9.

Electrolytic Au–Pd (Shimizu et al., 1976a,b)


A mixture of the chelate neutral cyanide bath of Au containing EDTA – Na2 and the ammonium phosphate complex bath of Pd was used for plating the Au – Pd alloy films (see Table 7.4.2.9.1). EDTA – Na2 was added as a chelating agent to stabilize the gold cyanide complex in the mixed bath. If EDTA – Na2 was not added to the mixed bath, white precipitates, which are most likely to be AuCN, were formed. The main ingredients of the bath were HAuCl4, PdCl2, HCl, (NH4)2HPO4, 30 g/l Na2HPO4·12H2O, and 10 g/l EDTA – Na2. The alloy composition of the plated Au –Pd alloy films was changed by varying the relative amount of HAuCl4/PdCl2 salts, and by varying the current density. The bath pH was 7 and the bath temperature was 50 8C (Table 7.4.2.9.2). After the plated Au –Pd alloy films, together with the substrates, were dissolved in aqua regia, the Au and Pd compositions were analyzed using an atomic absorption method. Polycrystalline and single-crystal ({100}, {110} and {111}) Cu sheets were used as the substrates.


The relationship between the Pd content of plated Au – Pd alloy films and the Pd concentration of the bath is plotted in Figure 7.4.2.9.1, which also displays the current efficiency of the bath. The plating was conducted for 15 and 50 s to obtain thin and thick Au – Pd alloy deposits. In the Au –Pd bath, Au and Pd metal ions have a large difference in deposition potential. Consequently, an alloy layer containing the more noble metal (Au), known as the initial layer [6], is deposited preferentially at the beginning of the plating until it reaches the equilibrium condition. The composition of the alloy films plated for 15 s was indeed richer in Au than the one plated for 50 s (see Figure 7.4.2.9.1) (Table 7.4.2.9.3).


Figure 7.4.2.9.2(a) is the equilibrium phase diagram of the Au –Pd alloys (Massalski, 1990). The phase diagram represents a solid-solution alloy system, which extends over the entire composition range. High-resolution electron diffraction patterns from various plated Au – Pd alloy films are shown in Figure 7.4.2.9.3. It is clear from these diffraction patterns that all the films are fcc. Based on this result, we calculated the lattice constants, which are 437

438

Nano-Plating

Table 7.4.2.9.1. The composition of Au– Pd alloy plating baths. Composition (g/l)

HAuCl4 (g/l) PdCl2 KCN (NH4)2HPO4 Na2HPO4·12H2O EDTA·Na2

Bath name Au

B

C

D

E

Pd

3.4 0 4.0 0 30 10

1.7 0.9 2.0 7.5 30 10

1.1 1.2 1.3 10.0 30 10

0.7 1.5 0.8 12.0 30 10

0.3 1.6 0.4 14.0 30 10

0 1.8 0 15.0 30 10

displayed over the entire composition range in Figure 7.4.2.9.2(b). Lattice constant data from both the thin and thick alloy films, which were obtained by depositing for 15 and 50 s, are included in this graph. The data follows a straight line connecting the lattice constant of pure Au (0.40786 nm (JCPDS File #4-0784)) and the constant of pure Pd (0.38902 nm (JCPDS File #46-1043)), in agreement with Vegards law. The early stages of Au –Pd alloy deposition were examined by TEM. Consistent with the results above, the initial layer was found to contain a greater proportion of Au. A lattice constant change with respect to deposition time in Au –Pd alloy films grown from four baths (B, C, D, and E) is observed in Figure 7.4.2.9.4. For all alloy deposits obtained from the four baths, the lattice constant was initially closer to that of Au because of the presence of the Au-rich layer, and then decreased with plating time (or film thickness), until the equilibrium conditions were reached. Figure 7.4.2.9.5 contains plan-view TEM micrographs showing the initial stage of (a) Au, (b) Au-(50 – 60) at.% Pd, and (c) Pd electrodeposits grown for 0.5 s on single-crystal {001} Cu substrates. In all cases, granular crystals from 5 to 20 nm in size are randomly distributed. The initial deposition morphology does not appear to depend on the alloy composition, although the size and density of the granular crystals changed. Electron diffraction patterns for these films exhibited rings, indicating that the granular crystals were randomly oriented without lattice-matching to the substrate. It is most likely that no chemical bonding was established with the substrate because the crystals were still very small. Similar results were obtained for Au – Pd alloys deposited on the {110} substrates. Table 7.4.2.9.2. The composition of Au– Pd alloy plating baths. Composition

HAuCl4 (g/l)

PdCl2 (g/l)

HCl (g/l)

a b c

0.8 0.5 0.2

0.2 0.5 0.8

5.0 5.0 5.0

Electrolytic Au –Pd

439

Figure 7.4.2.9.1. Relationship between the Pd content of plated Au–Pd alloy films and the Pd concentration of the bath.

The initial stages of nucleation and growth processes in Au – Pd alloy deposits proceeded with the appearance of granular crystals, followed by their growth, coalescence, and finally leading to the film formation. Here each crystal (or grain) began lattice-matching to the substrate. This growth sequence was the same for all alloy compositions but the grain size depended on the alloy composition. The grain size was large for pure metals (Au and Pd), but became smaller with alloying, as seen in Figure 7.4.2.9.6. Since all grains were lattice-matched to the substrate, they were single crystals, but each lattice-matched grain was connected through the sub-grain boundaries. Therefore, the grain size in these films can be defined as the distance between the subgrain boundaries.

Table 7.4.2.9.3. Plating conditions. Bath temperature Current density Plating time (s) Substrate

25 8C 30, 50, 100, 200, and 300 A/m2 0.5, 1.0, 5.0, 15.0, 40.0 Single-crystal Cu{001}, Cu{110}, Cu{111}

120 Polycrystalline Cu sheet

440

Nano-Plating

Figure 7.4.2.9.2. (a) The equilibrium phase diagram of Au–Pd alloys and (b) the lattice constants of plated Au–Pd alloys over the entire composition range.


441

Figure 7.4.2.9.3. High-resolution electron diffraction patterns from various plated Au–Pd alloy films.

Figure 7.4.2.9.4. A lattice constant change with deposition time in Au–Pd alloy films grown from four baths (B, C, D, and E).

442 Nano-Plating Figure 7.4.2.9.5. Plan-view TEM micrographs showing the initial stage of (a) Au, (b) Au –ð50 60Þ at.% Pd, and (c) Pd electrodeposits grown for 0.5 s on single-crystal {001} Cu substrates.


443

Figure 7.4.2.9.6. TEM micrographs and electron diffraction patterns from various compositions of Au–Pd alloy films grown on {001} Cu substrates for 50 s.

444

Nano-Plating

Figure 7.4.2.9.7. SEM micrographs showing the surface morphologies of (a) Au, (b) Au–8 at.% Pd, (c) Au–39 at.% Pd, (d) Au–52 at.% Pd, (e) Au–74 at.% Pd, and (f) Pd films plated for 15 min on polycrystalline Cu substrates.

Figure 7.4.2.9.7 displays SEM micrographs showing the surface morphologies of (a) Au, (b) Au –8 at.% Pd, (c) Au – 39 at.% Pd, (d) Au – 52 at.% Pd, (e) Au –74 at.% Pd, and (f) Pd films plated for 15 min on polycrystalline Cu substrates. The surface morphology for all the films appears random and does not reflect the grain structure of the substrate, except for films (a) and (b). Judging from these morphologies, we can conclude that the tendency of the alloy films to lattice-match to the substrate decreases with increasing Pd content.

REFERENCES Shimizu, Y. & Tanabe, Y. (1976a) J. Metal Finish. Soc. Jpn, 27, 19. Shimizu, Y. & Tanabe, Y. (1976b) J. Metal Finish. Soc. Jpn, 27, 574.


445

FURTHER READING Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, p. 409. JCPDS File #4-0784 (Au). JCPDS File #46-1043 (Pd).

Electrolytic Au– Sn (Tanabe, 1983; Kubota, 1983, 1985; Leidheiser, 1973; Matsushita, 1971; Parker, 1952) 7.4.2.10.


The compositions of the Au –Sn alloy plating bath used in this study are listed in Table 7.4.2.10.1. An 18-8 stainless steel plate was used as an anode. Plating was performed on polycrystalline and single-crystal ({001}, {110}, and {111}) Cu substrates at current densities of 100 and 5000 A/m2 in unstirred baths held at 65 8C. To prepare the TEM specimens, a plated Au –Sn film was stripped off by dissolving the Cu substrate in a mixed chromic acid/sulfuric acid solution (50 kg/m3 CrO3 þ 5 kg/m2 H2SO4).


Figure 7.4.2.10.1 compares the Sn content of plated Au – Sn alloy films with the Sn concentration of the baths. In this study, two current density values (100 and 5000 A/m2) and three plating times (15, 300, and 900 s) were chosen. In all cases, the Sn content of the films increased with increasing Sn concentration in the baths. For alloy plating systems containing metal elements such as Au and Sn, which have a markedly different deposition potential, the nobler metal (Au) tends to deposit preferentially at the start of the plating, followed by the formation of an alloy film with a bulk composition value. Therefore, the Au content will be higher for the thinner film plated for a shorter time (compare Figure 7.4.2.10.1(a) with (b)). The transient layer formed at the start of the plating is called the initial layer (Watanabe, 1983).


Figure 7.4.2.10.2 displays X-ray diffraction patterns taken from plated Au – Sn alloy films with various Sn contents. As suggested above, each film is expected to contain the initial later, which is richer in Au content. Despite the added complexity due to the presence of the initial layer, Tanabe et al. (1983) were able to perform a detailed structural analysis of plated Au – Sn alloy films and constructed the phase diagram seen in Figure 7.4.2.10.3(b) and (c). The equilibrium phase diagram (Massalski, 1990) of Au – Sn alloys is shown in Figure 7.4.2.10.3(a) for comparison. At the top of Figure 7.4.2.10.2, we listed X-ray diffraction line templates for Au (JCPDS File #4-784 (Au)), AuSn (JCPDS File #8-463 (AuSn)), AuSn2 (JCPDS File #28-440 (AuSn2)), AuSn4 (JCPDS File #28-441 (AuSn4)), b-Sn (JCPDS File #4-673 (b-Sn)), and Cu (JCPDS File #4-836 (Cu)), which were taken 446

447

Electrolytic Au –Sn Table 7.4.2.10.1. The composition of Au– Sn alloy plating baths. Composition (g/l)

HAuCl4·4H2O K2SnO3·3H2O KCN KOH

Bath name A

B

C

D

E

F

G

H

2.50 0.60 4.17 0.83

2.50 1.00 4.17 1.38

2.50 1.81 4.17 2.50

2.50 4.23 4.17 5.84

2.50 12.7 4.17 10.0

2.50 34.5 4.17 10.0

1.20 150 2.00 15.0

0.30 150 0.50 15.0

from the JCPDS file. The lattice constants (a ¼ 0:2923 nm; c ¼ 0:4782 nm; c=a ¼ 1:636Þ for the z-phase (14 at.% Sn) are due to Schubert and co-workers (1959). An X-ray diffraction analysis of AuSn2 and AuSn4 obtained by metallurgical processes of melting/ solidification can be found in publications by Schubert et al. (1950, 1959) and Kubota et al. (1983, 1985). Data in Figure 7.4.10.3(b) and (c) were obtained for the Au – Sn alloy films plated at 100 and 5000 A/m2, respectively. Although the phase diagram for the plated Au – Sn alloy films is different from the equilibrium phase diagram, it is not affected by the magnitude of the current density. Compared with the equilibrium phase diagram, the a-phase (Au solid-solution) region in the Au side became narrower (see Figure 7.4.2.10.3). In addition, the composition at

Figure 7.4.2.10.1. Relationship between the Sn content of plated Au–Sn alloy films and the Sn concentration of the baths.

448

Nano-Plating

Figure 7.4.2.10.2. X-ray diffraction patterns taken from plated Au–Sn alloy films having various Sn contents.

Electrolytic Au –Sn

449

Figure 7.4.2.10.3. (a) The equilibrium phase diagram of Au–Sn alloys; (b) and (c) are the phase diagrams of Au–Sn alloy films plated at current densities of 100 and 5000 A/m2, respectively.

which the z-phase appeared, moved toward the higher Sn side and changed to , 28 at.% Sn. The compositions at which AuSn, AuSn2, and AuSn4 emerged, agrees with the equilibrium phase diagram. The presence of the z-phase up to 67 at.% Sn and AuSn2 up to 90 at.% Sn, however, does not agree with the equilibrium phase diagram. It is interesting to note in Figure 7.4.2.10.3(b) and (c) that there are composition ranges containing three phases simultaneously. It is possible for two phases to co-exist in the equilibrium binary alloys. The presence of three phases in one alloy at the same time is very rare. Such a phenomenon, however, is often encountered in plated binary alloy films. It is important to remember that plated Au – Sn alloys do not form any meta-stable phases. Figure 7.4.2.10.4(a) and (b) are TEM micrographs for plated Au –19 at.% Sn alloy films grown on {001} Cu substrates for (a) 1 and (b) 3 s. In both cases, the film did not cover the

450

Nano-Plating

Figure 7.4.2.10.4. TEM micrographs for plated Au–19 at.% Sn alloy films grown at a current density of 100 A/m2 from the D bath on {001} Cu substrates for (a) 1 and (b) 3 s.

substrate surface completely. The electron diffraction patterns indicated that the films consisted of a mixture of the a- and z-phases, which were lattice-matched to the {001} Cu substrate. Figure 7.4.2.10.5(a) is a bright-field TEM micrograph showing the initial stage of a plated Au – 20 at.% Sn alloy film grown at 100 A/m2 on a {001} Cu substrate. The film appears to cover the substrate surface almost completely. The electron diffraction pattern displayed the 001 zone axis pattern, indicating that the film is lattice-matched to the {001} Cu substrate. Figure 7.4.2.10.5(b) and (c) are dark-field TEM micrographs taken from the (200) spot of the a-phase and the ð1100Þ spot of the z-phase, respectively. From these micrographs, we can estimate the size of the a-phase to be less than 10 nm and the z-phase to be , 10 nm. Figure 7.4.2.10.6 shows electron diffraction patterns from Au-58.2 at.% Sn films grown at 100 A/m2 for 8 s on (a) {001}, (b) {110}, and (c) {111} single-crystal Cu substrates. These films contained diffraction patterns from the z-phase, AuSn, and the Cu substrate. From these diffraction patterns, we obtained the following crystallographic


451

Figure 7.4.2.10.5. A bright-field TEM micrograph showing the initial stage of a plated Au–20 at.% Sn alloy film grown at 100 A/m2 on a {001} Cu substrate. (a) A BF image with its electron diffraction pattern, (b) a DF image taken from the (200) spot of the a phase, and (c) a DF image from the (1100) spot of the j phase.

relationships. {001} Cu==ð0001Þ z; ½110 Cu==½1100 z or {001} Cu==ð0001Þ z; ½110 Cu==½1100 z {011} Cu==ð0111Þ z; ½011 Cu==½21; 10 z or {011} Cu==ð0112Þ z; ½011 Cu==½21; 10 z {111} Cu==ð0001Þ z; ½011 Cu==½21; 10 z From Figure 7.4.2.10.6(a), we deduced the following epitaxial relationship. {001} Cu==ð2110Þ AuSn; , 110 . Cu==½0001 AuSn Figure 7.4.2.10.7 shows a SEM micrographs for plated Au –Sn alloy films, showing characteristic surface morphologies, which are seen to vary with the alloy composition. Kubota et al. (1983, 1985) plated Au – Sn alloy films with various alloy compositions using a mixed bath of gold cyanide and stannous pyrophosphate salts. They compared these results with the equilibrium phase diagram. According to their results, the eutectic

452 Nano-Plating

Figure 7.4.2.10.6. Electron diffraction patterns from Au –58.2 at.% Sn films grown from the E bath at 100 A/m2 for 8 s on (a) {001}, (b) {110}, and (c) {111} single-crystal Cu substrates.


Figure 7.4.2.10.7. Relationship between surface morphologies and various compositions of plated Au–Sn alloy films.

453

454

Nano-Plating

point coincided, but the hexagonal AuSn phase appeared at a composition of 40 at.% Sn rather than at 50 at.% Sn.

REFERENCES JCPDS File #4-784 (Au). JCPDS File #8-463 (AuSn). JCPDS File #28-440 (AuSn2). JCPDS File #28-441 (AuSn4). JCPDS File #4-673 (b-Sn). JCPDS File #4-836 (Cu). Kubota, N., Horikoshi, T. & Sato, E. (1983) J. Metal Finish. Soc. Jpn, 34, 37. Kubota, N., Yoshimura, S. & Sato, E. (1985) J. Metal Finish. Soc. Jpn, 36, 355. Leidheiser, H. & Ghuman, A.R.P. (1973) J. Electrochem. Soc., 120, 484. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 1, ASM International, pp. 433– 434. Massalski, T.B., Breimer, H., Gohle, R., Schubert, K., Breimer, H. & Gohle, R. (1959) Z. Metallkunde, 50, 146. Matsushita, S. (1971) J. Metal Finish. Soc. Jpn, 22, 60. Parker, E.A. (1952) Plating, 39, 43. Schubert, K., Breimer, H. & Gohle, R. (1950) Z. Metallkunde, 41, 298. Schubert, K., Breimer, H. & Gohle, R. (1959) Z. Metallkunde, 50, 146. Tanabe, Y., Hasegawa, N. & Odaka, M. (1983) J. Metal Finish. Soc. Jpn, 34, 452. Watanabe, T. (1983) J. Metal Finish. Soc. Jpn, 41, 652.

7.4.2.11.

Electrolytic Cd– Sn (Aotani, 1952)


A mixture of sodium stannate (Na2Sn) and cadmium cyanide (Cd(CN)2) was used to make a Sn – Cd alloy plating bath. Electrolytic Cd –Sn alloy films having various compositions were obtained by changing the mixing ratio of these two metal salts. The plating was conducted on 2 £ 1 cm2 iron plate substrates at a current density of 200 A/m2 in an unstirred 70 8C bath. A structural analysis of the plated films was performed using a Debye camera.


Unfortunately precise data were not provided in the references. 7.4.2.11.2.1 Film microstructure Figure 7.4.2.11.1(a) is the equilibrium phase diagram of Sn – Cd binary alloys, whereas (b) shows the phases/lattice constants of electrolytic Cd –Sn alloys. For the lattice constants, only the a axis value was plotted for convenience. Although the b phase forms at high temperatures as seen in Figure 7.4.2.11.1(a), this alloy system is considered to be primarily a two-phase separation type at ambient temperatures. Indeed, the b phase did not appear in the plated films and only the Sn and Cd crystals indicative of a two-phase separation type were present. The lattice constants for these phases were constant over the whole composition range.

455

456

Nano-Plating

Figure 7.4.2.11.1. (a) The equilibrium phase diagram of Cd –Sn alloys and (b) the lattice constants of Cd/Sn phases in electrolytic Cd –Sn alloys.

REFERENCE Aotani, K. (1952) Denki Kagaku, 20, 611.

7.4.2.12.

Electrolytic Cd– Zn (Aotani, 1951, 1953)


A Cd – Zn alloy plating bath was prepared as a complex mixture solution from CdSO4 and ZnSO4 with a total metal salt concentration of 30 g/l. The composition of plated Cd – Zn alloy films was varied by changing the mixing ratio of these metal salts. NaOH (20 g/l) was also added to the solution. The plating was conducted on 2 £ 1 cm2 iron substrates at a current density of 2000 A/m2 in the unstirred 70 8C bath. The Cd content was determined by separating Cd from the alloy films as CdSO4. A structural analysis of the plated films was performed using a Debye camera.


Unfortunately precise data were not provided in the references.


Figure 7.4.2.12.1 is (a) the equilibrium phase diagram of Cd – Zn binary alloys and (b) phases/lattice constants observed in electroplated Cd – Zn alloys. Both Cd and Zn have a hcp structure and form a typical eutectic system. Both metals have a small degree of mutual solubility at ambient temperatures. In the plated films, two phases co-exist in the composition range 19 –74 at.% Zn. Outside this composition range, only single phases appeared. It is important to note that the lattice constants do not change over the entire composition range.

REFERENCES Aotani, K. (1951) J. Jpn Inst. Metals, B-15, 21. Aotani, K. (1953) Denki Kagaku, 21, 21.

457

458

Nano-Plating

Figure 7.4.2.12.1. (a) The equilibrium phase diagram of Cd–Zn binary alloys and (b) phases/lattice constants observed in electroplated Cd –Zn alloys.

Electrolytic Co – Cu (Subrahamanyam, 1967; Morral, 1967; Shimizu, 1978; Miyazaki, 1999)

7.4.2.13.


The compositions of the Co – Cu alloy plating bath used in this study are listed in Table 7.4.2.13.1 (Subrahamanyam and Rama Char, 1967). The mixing ratio of Co/Cu salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant (0.2 mol/l). Thin electrolytic Au films and electropolished pure polycrystalline Cu foils were used as the substrates. Co – Cu alloy films were electrodeposited directly onto electron-transparent thin Au films for , 1– 20 s. The alloy film and the substrate were examined simultaneously using TEM. For the study of alloy phases and surface morphologies, the alloy films were electrodeposited for 5 min. A structural analysis of the alloy phases was performed using a high-resolution reflection electron diffraction unit attached to the TEM. For a composition analysis of the alloy films, the films were first dissolved in nitric acid and then studied using an atomic absorption method.


The Co content of the deposits and the current efficiency are plotted as a function of Co concentration in Co –Cu alloy plating baths (see Figure 7.4.2.13.1). The Co content of the deposits is also plotted in terms of plating time (film thickness). The thin alloy films deposited for a short time tended to exhibit a higher Cu content, due to the formation of an initial layer.


Reflection electron diffraction patterns from various Co –Cu alloy films are shown in Figure 7.4.2.13.2. These patterns, in combination with the conventional electron diffraction patterns obtained in TEM, were used to determine the lattice constants of the alloy films. Figure 7.4.2.13.3(a) is the equilibrium phase diagram of Co –Cu alloys (Massalski, 1990) and (b) shows lattice constants for the various phases of the corresponding alloy films obtained by a plating method. Note that the lattice constant data reported by Subrahamanyam and Rama Char (1967) are also plotted in Figure 7.4.2.13.3(b). This alloy system is a peritectic type. It is well known that pure 459

460

Nano-Plating

Table 7.4.2.13.1. The composition of Co–Cu alloy plating baths. Composition (g/l)

Cu2P2O7·6H2O CoCl2·6H2O K4P2O7 (NH4)2HC6H5O7

Bath name Cu

A

B

C

D

E

F

G

H

I

Co

39.1 0 270 10

25.4 16.7 270 10

19.6 23.8 270 10

13.7 30.9 270 10

9.8 35.7 270 10

7.8 38.1 270 10

3.9 42.8 270 10

2.0 45.2 270 10

1.0 46.4 270 10

0.4 47.3 270 10

0 47.6 270 10

Co has an fcc structure (a phase) at high temperatures but becomes hcp (1 phase) at low temperatures. In the Co-rich side, up to 19.7 at.% Cu can dissolve in the fcc phase at 1367 8C but no Cu dissolves at ambient temperatures. In the Cu-rich side, 9.5 and 10.0 at.% Co are shown to dissolve at 1050 and 1112 8C, respectively. No Co dissolves at ambient temperatures. As seen in Figure 7.4.2.13.3(b), fcc phases appeared over the whole composition range in plated Co –Cu alloys. On the Co-rich side, the lattice constant of these fcc phases increased with increasing Cu content, whereas on the Cu-rich side, the lattice constant of pure Cu (0.362 nm) decreased slowly with increasing Co content and reached 0.361 nm

Figure 7.4.2.12.1. The Co content of plated Co –Cu alloy films and the current efficiency plotted as a function of Co concentration in Co –Cu alloy plating baths.

Electrolytic Co –Cu

Figure 7.4.2.13.2. Reflection electron diffraction patterns from various Co –Cu alloy films.

461

462

Nano-Plating

at 31 at.% Co. With increasing alloying element, the lattice constants on both the Co- and Cu-rich sides reached saturation at a particular alloy composition. This alloy composition was 20 at.% Cu on the Co-rich side and 20 at.% Co on the Cu-rich side. According to the equilibrium phase diagram, Cu cannot dissolve in Co at ambient temperatures. Similarly, Co cannot dissolve in Cu. For this reason, the alloys obtained by a plating method are

Figure 7.4.2.13.3. (a) The equilibrium phase diagram of Co –Cu alloys and (b) the lattice constants for various phases of plated Co–Cu alloy films.


463

considered to be supersaturated solid-solution alloys, which are meta-stable phases. The lattice constant of the fcc phase on the Co-rich side found in this experiment agrees with the constant reported by Klement (1963). Shimizu et al. (1978) reported that the fcc phase in the Cu-rich side is entirely different from the Co-rich side. Furthermore, in the composition range of 0– 60 at.% Cu an hcp phase also appeared, forming two fcc/hcp phases. Newman (1950) and Nakahara and Mahajan (1980) reported the appearance of the fcc Co phase stable at high temperatures in electrolytic Co films grown at room temperature. In the Co – Cu alloy films, the lattice constants (a ¼ 0:251 nm and c ¼ 0:410 nm) of a pure electrolytic Co film increased with increasing Cu content, but reached saturation at 30 at.% Cu. TEM examinations of the initial stages of the Co –Cu alloy deposition indicated that the growth mode took a Volmer – Weber type, which proceeds by the formation of islands (see Figure 2.3 in Chapter 2).

Figure 7.4.2.13.4. TEM micrographs and electron diffraction patterns showing the early stages of nucleation and growth of Cu– 85 at.% Co alloy films, which were electrodeposited for (a) 1, (b) 5, and (c) 15 s. (c0 ) is a schematic diagram illustrating the structure and orientation of grains seen in (c).

464 Nano-Plating

Figure 7.4.2.13.5. SEM micrographs showing the surface morphologies of various plated Co–Cu alloy films.


465

Figure 7.4.2.13.4 contains TEM micrographs showing the early stages of nucleation and growth in Cu– 85 at.% Co alloy films, which were electrodeposited for (a) 1, (b) 5, and (c) 15 s. Since the deposition time is different, it is possible that the composition of these films may be different as described in Figure 7.4.2.13.1. In Figure 7.4.2.13.4(a) and (b), the Cu phase (the lattice constant a ¼ 0:361 , 0:362 nm) and the fcc Co crystals were found to co-exist, whereas in Figure 7.4.2.13.4(c), the fcc and hcp Co crystals were present. The alloy film in Figure 7.4.2.13.4(a) consisted of an assembly of 10-nm grains but the electron diffraction pattern showed diffraction spots characteristics of a single-crystal pattern. Therefore, it appears that these fine (10 nm) grains have grown epitaxially on the gold substrate. Furthermore, electron diffraction patterns in Figure 7.4.2.13.4(b) and (c) exhibit orthogonal streaks around the diffraction spots, indicating that these films are textured. The streaks in the [010] and [001] directions correspond to the lattice matching of the alloy crystals along the two orthogonal directions (see a schematic diagram in Figure 7.4.2.13.4(c0 )). Figure 7.4.2.13.5 displays SEM micrographs showing the surface morphologies of various plated Co –Cu alloy films. It is clear that the surface is rough for the Cu-rich alloys and smooth for the Co-rich alloys.

REFERENCES Klement, W. (1963) Trans. AIME, 227, 965. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, vol. 2, 2nd Edition, ASM International, p. 1181. Miyazaki, K., Kainuma, S., Hisatake, K. & Watanabe, T. (1999) Electrochim. Acta, 44, 2713. Morral, F.R. (1967) Plating, 54, 639. Nakahara, S. & Mahajan, S. (1980) J. Electrochem. Soc., 127, 283. Newman, R.C. (1950) Proc. Phys. Soc., 69, 432. Shimizu, Y., Tanabe, Y., Tomita, H. & Kakegawa, M. (1978) J. Metal Finish. Soc. Jpn, 29, 131. Subrahamanyam, D.V. & Rama Char, T.L. (1967) Electroplat. Metal Finish., 20, 44.

7.4.2.14.

Electrolytic Co –Fe (Fukumuro et al., 1999)


Electrodeposition of Co – Fe alloys occurs by a so-called anomalous co-deposition mechanism, whereby an electrochemically less noble Fe metal deposits preferentially, thus making control of the alloy composition difficult (Aotani, 1950; Dahms and Croll, 1965). In practice, different organic additives are mixed in the bath to control the composition. In the present study, however, no such additives were used. Cobalt and iron sulfate salts were a source of metal ions. The mixing ratio of Cobalt sulfate and iron sulfate salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 1 mol/l. The salt mixture was then dissolved in 250 ml deoxidized distilled water. In addition to the metal salts, 0.5 mol/l boric acid was added as a buffer and the bath pH was adjusted to 2.0 with dilute sulfuric acid. A platinum plate was used as a counter electrode and a saturated calomel electrode (SCE) as the reference electrode. The alloy films were electroplated by a controlled potential electrolysis method.


Figure 7.4.2.14.1 is a plot for the Fe content in plated Co –Fe alloys versus the Fe concentration in the baths. For all the baths, the three constant potentials of 2 1.0, 2 1.5, and 2 2.0 V versus SCE were applied to plate the Co – Fe alloy films. As seen in Figure 7.4.2.14.1, any change in the alloy composition due to deposition potential was found to be very small. The alloy composition is primarily determined by the bath composition. The three curves are slightly convex, suggesting a tendency for the less noble metal, Fe, to be deposited preferentially.


Figure 7.4.2.14.2 shows X-ray diffraction patterns from 2 pure metals (Co and Fe) and 10 Co – Fe alloy films plated on polycrystalline Cu substrates. For the films starting from pure Fe down to Co – 28 at.% Fe, the {110} and {211} peaks of the bcc a phase were present. These peaks are seen to shift toward the high-angle side with decreasing Fe content. At a composition of Co –45 at.% Fe, the {200} peak of the fcc g phase emerged. With a further decrease in the Fe content, the {111} and {220} peaks of the g phase appeared at the composition of Co – 19 at.% Fe and became very strong at Co –10 at.% Fe. In addition, 466

Electrolytic Co – Fe

467

Figure 7.4.2.14.1. The Fe content of plated Co– Fe alloys plotted against the Fe concentration of the baths.

a decrease in the Fe content was accompanied by a shift in the {200} and {220} peaks of the g phase toward the low-angle side. This shift corresponds to an expansion in the interplanar spacing of the 200 and 220 planes. In pure Co films, the (100) peak of the hcp 1phase appeared, whereas the {111} and {200} peaks of the g phase almost vanished. Conversely the {220} peak of the g phase overlapped with the {110} peak of the 1-phase, but its intensity further increased. It is well known that a high-temperature phase, fcc a-Co, can be obtained by an electrodeposition method at room temperature. It has been shown (Sasaki and Talbot, 1995) that the formation of a-Co resulted from a simple structural modification of the hcp 1-Co by stacking faults, and that co-deposited hydrogen atoms were responsible for this structural modification. Our TEM observations also confirmed the presence of a-Co and stacking faults. The 1 phase is expected to be present inside the g phase as faulted layers. Based on the above experiment, we summarize the results in Figure 7.4.2.14.3. Here, the phases observed in plated Co – Fe alloy films and their lattice constants are shown in Figure 7.4.2.14.3(b) and (c), respectively, and compared with the equilibrium phase diagram of Co –Fe alloys (see Figure 7.4.2.14.3(a)). In plated Co – Fe alloys, the 1 phase formed in the composition range 0 – 20 at.% Fe, whereas the g phase appeared in the range 0 –45 at.% Fe. In the range 10 –20 at.% Fe, it is possible that the 1 phase exists as the faulted g phase. The fcc a phase was generated in the range 10– 100 at.% Fe. In the range

468 Nano-Plating

Figure 7.4.2.14.2. X-ray diffraction patterns taken from 2 pure metals (Co and Fe) and 10 Co–Fe alloy films plated on polycrystalline Cu substrates.


469

Figure 7.4.2.14.3. (a) The equilibrium phase diagram of Co–Fe alloys, (b) the phase diagram, and (c) the lattice constants of plated Co–Fe alloy films.

470

Nano-Plating

10 –20 at.% Fe, three phases (a, 1, and g phases) existed together, while in the range 20 –45 at.% Fe, two phases (a and g phases) were present. In comparison with the equilibrium phase diagram, we note that the composition range for each phase extended by 10 –20 at.%. The lattice constants for several phases observed in plated Co –Fe alloys are plotted in Figure 7.4.2.14.3(c). The lattice constant of the a phase was calculated from the position of the {110} diffraction peak. The lattice constant of the a phase decreased linearly with decreasing Fe content, from 0.2868 nm of pure Fe to 0.2850 nm of Co –14 at.% Fe. This trend agrees with a previous report (Nakahara and Mahajan, 1980). The lattice constant of the g phase was calculated from the position of the {200}diffraction peak. The lattice constant of the g phase increased linearly with increasing Fe content from the 0.3547 nm of pure Co to the 0.3587 nm of Co – 45 at.% Fe. These lattice constant values lie on a line connecting the lattice constant (0.35447 nm) of fcc a-Co and the constant (0.36394 nm) of fcc g-Fe, thus obeying Vegards law. The lattice constants of the a and g phases in plated Co –Fe alloy films clearly indicate a linear relationship with film composition, although they did exhibit a small scatter in the data. Such a data scatter often arises from hydrogen inclusions (Brenner, 1963; Saga and Miyata, 1976; Raub, 1993) and internal stresses (Fukai, 1985). A change in the lattice constants of the 1-Co with respect to composition was not studied here for the following reasons. First, the {002} and {110} peaks of the 1 phase overlapped with the {111} and {220} peaks of the g phase. Second, the intensity of the {100} peak for the 1 phase was not strong enough for an analysis of the lattice constant.

7.4.2.14.3.1

Crystal size

The grain size of several phases (a, 1, and g phases) in plated Co – Fe alloy films were determined by the full width at half maximum (FWHM) of a diffraction peak using Scherrer’s equation (see Figure 7.4.2.14.4). The grain size of the a phase, which was obtained at deposition potentials of 2 1.0 and 2 2.0 V versus SCE, was calculated using the (110) peak. No significant differences in the grain size of the a phase between the two deposition potentials were observed. The grain size of the a phase started with , 30 nm for pure Fe, then decreased to 20 nm upon alloying with Co, and finally reached the average size of 15 nm in the composition range of 40 – 15 at.% Fe. The grain size of the g phase obtained at 2 2.0 V versus SCE was calculated using the {200} peak. As described above, (1) the 1 phase exists as the faulted g phase and (2) the {002} peak of the 1 phase overlaps with the {111} peak of the g phase. Due to these complications, it was not possible to determine the grain size of the 1 phase uniquely. Consequently, the grain size, which was determined from either the {002} peak of the 1 phase or the {111} peak of the g phase, was taken to be the 1 or g phase. Regardless of the


471

Figure 7.4.2.14.4. The grain size of several phases (a, 1, and g phases) in plated Co– Fe alloy films plotted as a function of the Fe content of the films. The grain size was determined using the Scherrer’s equation.

magnitude of the deposition potential, the grain size of the a/g phases decreased rapidly by alloying with Fe from 55 nm (pure Co) to 15 nm (Co– 15 at.% Fe). An SEM and TEM study was conducted on Co –Fe alloy films plated at a deposition potential of 2 2.0 V versus SCE on polycrystalline Cu substrates and is summarized in Figures 7.4.2.14.5.1 – 5.3. Alloy films having six compositions ((A) Co –10 at.% Fe, (B) Co – 18 at.% Fe, (C) Co –28 at.% Fe, (D) Co –38 at.% Fe, (E) Co – 52 at.% Fe, and (F) Co – 74 at.% Fe) were investigated here. For an SEM study, both the top and fracture surfaces were examined simultaneously (see (a)). For TEM examinations, cross-section TEM specimens were prepared using a microtome technique and the (b) bright-field (BF) and (c) dark-field (DF) images together with (d) the electron diffraction pattern were displayed for each specimen. Although the deposition potential affected the surface morphology, the use of 2 2.0 V versus SCE produced films with uniformly smooth surfaces as seen in Figure 7.4.2.14.5.1 –5.3. An energy-dispersive X-ray analysis of the cross-section specimens indicated that there was no compositional change along the film thickness direction; meaning the alloy composition was uniform throughout the film including the film/Cu interface region. The film with the composition of Co – 18 at.% Fe consists of a mixture of the 1 and g phases and its grain size was determined to be , 30 nm using Scherrer’s equation. Direct measurements from the DF images yielded a grain size of 10 –60 nm. For the film with (B) Co –18 at.% Fe, three phases (g, 1, and a phases) co-existed and their grain size was less than 12 nm, which was the smallest among all the films studied here.

472

Nano-Plating

Figure 7.4.2.14.5.1. (a) An SEM image showing the fracture surface of plated (A) Co –10 at.% Fe and (B) Co–18 at.% Fe alloy films, the cross-section TEM images ((b) BF and (c) DF images), and (d) the electron diffraction pattern.


473

Figure 7.4.2.14.5.2. (a) An SEM image showing the fracture surface of plated (C) Co–28 at.% Fe and (D) Co –38 at.% Fe alloy films, the cross–section TEM images ((b) BF and (c) DF images), and (d) the electron diffraction pattern.

474

Nano-Plating

Figure 7.4.2.14.5.3. (a) An SEM image showing the fracture surface of plated (E) Co–52 at.% Fe and (F) Co – 74 at.% Fe alloy films, the cross– section TEM images ((b) BF and (c) DF images), and (d) the electron diffraction pattern.


475

In the films containing (C) Co – 28 at.% Fe and (D) Co –38 at.% Fe, the presence of the 1 phase was not recognized and the formation probability of the g phase was greatly reduced. For this reason, the grain size of the a phase increased to 20 nm. This change can also be recognized in their DF micrographs. As seen in the SEM micrographs, these films had a columnar structure, in which the width of each column was , 0.5 mm. In the films containing (E) Co –52 at.% Fe and (F) Co– 74 at.% Fe, no g phase was recognized and only the grains of the a phase forming large columns were found.

REFERENCES Aotani, K. (1950) J. Jpn Inst. Metals, B-14, 55. Brenner, A. (1963) Electrodeposition of Alloys, vol. II, Academic Press, New York, p. 304. Dahms, H. & Croll, I.M. (1965) J. Electrochem. Soc., 112, 771. Fukai, Y. (1985) J. Jpn Inst. Metals, 24, 671. Fukumuro, N., Chikazawa, M. & Watanabe, T. (1999) J. Surf. Finish. Soc. Jpn, 50, 448. Nakahara, S. & Mahajan, S. (1980) J. Electrochem. Soc., 127, 283. Raub, Ch.J. (1993) Plat. Surf. Finish., 80(September), 30. Saga, J. & Miyata, S. (1976) J. Jpn Inst. Metals, 40, 1098. Sasaki, K.Y. & Talbot, J.B. (1995) J. Electrochem. Soc., 142, 775.

7.4.2.15.

Electrolytic Co –Mo (Watanabe et al., 1989)


Co – Mo alloy films were electroplated in baths containing a mixture of cobalt sulfate (CoSO4·7H2O) and sodium molybdate (Na2MoO4·4H2O). The mixing ratio of CoSO4·7H2O and Na2MoO4·4H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant (0.26 mol/l). Sodium citrate was added as a complexing agent. The bath pH was adjusted with ammonia water or dilute sulfuric acid. Detailed studies were performed on the alloy films grown at pH ¼ 3 and 5. The bath temperature was set at 33 8C. The composition of these alloy films was determined using an atomic absorption method. A structural analysis of the films was made by an X-ray diffractometer, which used a cobalt Ka line as the irradiation source.


The relationship between Mo concentration in two baths ((a) pH ¼ 3 and (b) pH ¼ 5Þ and Mo content in the plated films is shown in Figure 7.4.2.15.1, in which three current density values (400, 600, and 800 A/m2) were used. In those two baths, the current density does not appear to affect the relationship. The effect of bath pH on Mo content in Co –Mo alloy films plated at 400 A/m2 from the bath containing 60 mol% (0.104 mol/l) sodium molybdate is indicated in Figure 7.4.2.15.2. The Mo content is very sensitive to bath pH and has a peak at pH ¼ 5: At high (. 8) pH, the Mo content depends on the type of buffer used. This result strongly suggests that the Mo content does not change with pH but depends on the type of anions associated with the buffer.


7.4.2.15.3.1 Structure X-ray diffractions patterns for various Co – Mo alloy deposits grown at 400 A/m2 from the pH ¼ 3 baths are shown in Figure 7.4.2.15.3. These patterns are somewhat complicated busy because diffraction peaks from the Cu substrate are also present. Nevertheless, diffraction peaks from the alloy films tend to become broader with increasing Mo content. Peak broadening is due to the refining of the grains. According to the patterns in Figure 7.4.2.15.3, the crystalline grains are still present for the Co –41 at.% Mo film. 476

Electrolytic Co– Mo

477

Figure 7.4.2.15.1. Relationship between the Mo concentration of the two baths ((a) pH ¼ 3 and (b) pH ¼ 5Þ and the Mo content of plated Co –Mo films.

Although an X-ray diffraction method alone cannot absolutely and positively identify the presence of an amorphous phase, it is safe to assume that beyond the 41 at.% Mo, the films became amorphous. Using the X-ray criteria for the presence of an amorphous phase, the structure (crystalline, amorphous, or mixed phases) of Co – Mo alloy films obtained from

Figure 7.4.2.15.2. The effect of bath pH on the Mo content of plated Co –Mo alloy films.

478

Nano-Plating

Figure 7.4.2.15.3. X-ray diffractions patterns from various Co –Mo alloy deposits.

the two baths ((a) pH ¼ 3 and (b) pH ¼ 5Þ is plotted in Figure 7.4.2.15.4 as a function of Mo concentration in the bath and current density. From this plot, it can be concluded that the Co – Mo alloy films become amorphous in the high Mo concentration baths regardless of the current density. A brief sketch for the composition range of an amorphous phase in plated Co –Mo alloy films is drawn in Figure 7.4.2.15.5 (b) together with Figure 7.4.2.15.5 (a) the equilibrium phase diagram of Co –Mo alloys. In the present experiment, the highest attainable Mo content in the films was 55 at.% Mo.


479

Figure 7.4.2.15.4. The structure (crystalline, amorphous, or mixed phases) of Co– Mo alloy films obtained from the two baths ((a) pH ¼ 3 and (b) pH ¼ 5Þ plotted as a function of Mo concentration in the bath and the current density.

Figure 7.4.2.15.5. (a) The equilibrium phase diagram of Co–Mo alloys and (b) a brief sketch for the composition range of an amorphous phase found in plated Co –Mo alloy films.

480

Nano-Plating

Figure 7.4.2.15.6. SEM micrographs showing the surface morphology and fracture surface of three plated Co – Mo alloy films.


481

Figure 7.4.2.15.6 contains SEM micrographs showing the surface morphology and fracture surface of three plated Co –Mo alloy films, which were identified to be amorphous. The fracture surface morphology changes with the Mo content.

REFERENCE Watanabe, T., Naoe, T., Mio, J. & Katsumata, A. (1989) J. Surf. Finish. Soc. Jpn, 40, 458.

7.4.2.16.

Electrolytic Co –Ni (Fukumuro, 1999)


A Co – Ni alloy plating bath consists of cobalt sulfate (CoSO4·7H2O) and nickel sulfate (NiSO4·6H2O) as a source of metal salts. The mixing ratio of CoSO4·7H2O and NiSO4·6H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 1 mol/l. These metal salts were dissolved in deoxidized/distilled water together with 0.5 mol/l H3BO3, and the resulting solution was used as a buffer. The bath pH was adjusted to 2.0 with dilute sulfuric acid. A platinum plate was used as a counter electrode and a saturated calomel electrode (SCE) as the reference electrode. Cobalt –nickel alloy films were electroplated by a controlled potential electrolysis method. The constant potentials ranging from 2 0.8 to 2 2.0 V versus SCE were applied to plate the films. Polycrystalline Cu and stainless steel sheets (SUS-304) were used as the substrates.


Figure 7.4.2.16.1 shows the Ni content of Co – Ni alloy films plated on polycrystalline Cu substrates plotted against the Ni concentration in the bath. As described above, various alloy films were obtained under different constant potentials ranging from 2 0.8 to 2 2.0 V versus SCE. Figure 7.4.2.16.1 shows that the alloy composition was not affected by the deposition potential but rather controlled by the bath composition. Although Ni metal ions are electrochemically nobler than Co metal ions, Ni deposition is suppressed in the present Co – Ni plating system. For this reason, the Co –Ni system is classified as an anomalous codeposition type. For example, to obtain Co – Ni alloy deposits containing 50 at.% Ni, we have to add nickel sulfate at quantities 20 times that of cobalt sulfate to the bath. The downward concave-shaped curve in Figure 7.4.2.16.1 also implies the anomalous codeposition type. Although the above results are based on polycrystalline Cu substrates, similar results were obtained on stainless steel substrates.


7.4.2.16.3.1 Structure Figure 7.4.2.16.2 shows X-ray diffraction patterns from various compositions of Co – Ni alloy films plated at 2 2.0 V versus SCE. The diffraction patterns are seen to change with 482

Electrolytic Co– Ni

483

Figure 7.4.2.16.1. The Ni content of Co –Ni alloy films plated on polycrystalline Cu substrates plotted against the Ni concentration in the bath.

increasing Ni content, i.e. from (a) Co – 15 at.% Ni to (e) Co –74 at.% Ni. The presence of hcp 1 and fcc a phases were identified from these patterns as marked in the figure. It should be noted that the {002} peak of the 1 phase overlaps with the (111) peak of the a phase. Similarly, the {110} and {112} peaks of the 1 phase overlap with the (220) and (311) peak of the a phase, respectively. For the present analysis, therefore, we avoided these three peaks. The appearance of the {100}, {110}, {200}, and {201} peaks of the 1 phase in the (a) 15 and (b) 26 at.% Ni films indicate that the 1 phase was formed. In addition, there is a small (200) peak of the a phase for (b) the 26 at.% Ni film, and thus both the a and 1 phases are considered to coexist in this film. In (c) the 38 at.% Ni film, all the diffraction peaks from the 1 phase markedly decreased, but the strong (200) peak of the a phase emerged. Around this alloy composition, the structure of the film transformed from the 1/a mixed phase to the single a phase. Finally, the structure of (d) the 50 and (e) 74 at.% Ni films became a single a phase. From the intensity of the X-ray diffraction peaks, we found that the 1 phase had a 110 texture, whereas the a phase had a 220 texture. Figure 7.4.2.16.3(a)– (c) shows the equilibrium phase diagram of Co –Ni alloys, phases observed in plated Co – Ni alloy films, and their lattice constants, respectively. The alloy films used to construct the two-dimensional phase diagram in Figure 7.4.2.16.3(b) were plated at 2 2.0 V versus SCE on stainless steel substrates and at 2 0.8, 2 1.6, and 2 2.0 V versus SCE on polycrystalline Cu substrates. As indicated in Figure 7.4.2.16.3(b), the fcc a

484

Nano-Plating

Figure 7.4.2.16.2. X-ray diffraction patterns from various compositions of plated Co –Ni alloy films.

phase was formed over the entire composition range, whereas the hcp 1 phase was produced in the composition range of 0 , 40 at.% Ni. Except for the a phase formed below the composition of 35 at.% Ni, the phase diagram in Figure 7.4.2.16.3(b) is in good agreement with the equilibrium phase diagram shown in Figure 7.4.2.16.3(a). The a phase below the composition of 35 at.% Ni is considered to be a solid-solution Co alloy supersaturated with Ni. Nakahara and Mahajan (1980) reported that the fcc a-Co phase can be formed by an electrolytic method, whereby co-deposited hydrogen atoms transformed the hcp Co into the fcc Co film by introducing stacking faults. Nevertheless, these results further confirm that an electrolytic method can readily generate non-equilibrium phases in plated alloy films. A change in the lattice constant of the a phase was computed from the shift in the {111} peak of the a phase. The lattice constants, a and c, of the 1 phase were calculated from the (100) and (002) diffraction peaks, respectively. The (002) peak of the 1 phase coincided with the {111} peak of the a phase. The lattice constant of the a phase showed a linear relationship with alloy composition, although it was affected slightly by the deposition potential. The lattice constant data appear to follow a dotted line connecting the lattice constant of Ni ða ¼ 0:3524 nmÞ and the constant of fcc Co ða ¼ 0:35447 nmÞ: The lattice constant of the a phase in plated


485

Figure 7.4.2.16.3. (a) The equilibrium phase diagram of Co–Ni alloys, (b) the phase diagram of plated Co –Ni alloy films, and (c) the lattice constants of their phases.

486

Nano-Plating

Co – Ni alloys, therefore, follows Vegards law. It is interesting to note that the lattice constants of the films grown at 2 2.0 V versus SCE are much larger than the ones obtained at 2 0.8 and 2 1.6 V versus SCE. The lattice constant difference can be attributed to a lattice expansion due to the incorporation of hydrogen (Fukai, 1985). The lattice constant, a, of the 1 phase, which was derived from the (100) diffraction peak, increased slightly with increasing Ni content. Similar to our findings (Jartych, 1993), there is a report (Aotani, 1950) claiming that the phase structure of plated Co –Ni alloy films is in good agreement with the equilibrium phase diagram. Aotani found that the bath pH also affected the phase structure and its lattice constants. 7.4.2.16.3.2 Crystal size The grain (crystal) size of the a and 1 phases in Co –Ni alloy films plated on polycrystalline Cu and stainless steel substrates was plotted in Figure 7.4.2.16.4 as a function of the Ni content in the film. The crystal size was determined from the full width at half maximum (FWHM) of an X-ray diffraction peak using Scherrer’s equation. Although three different deposition potentials (2 0.8, 2 1.6, and 2 2.0 V versus SCE) were used, the grain size of the a phase was not affected by the deposition potential. The grain size of the a phase was 30 nm in the composition range of 40– 100 at.% Ni. In the composition range less than 40 at.% Ni, where both the a and 1 phases coexisted, the grain size was 20 nm. The grain size-to-composition relationship for the Co –Ni alloy films grown on stainless steel substrates resembles the relationship with polycrystalline Cu substrates, but the grain

Figure 7.4.2.16.4. Relationship between the grain (crystal) size of the a/1 phases in plated Co–Ni alloy films and the Ni content of the films.


487

size of the a phase grown on the former substrates was generally smaller. This grain size difference will be discussed later. The grain size of the 1 phase was also determined by an X-ray method from the FWHM of the (100) X-ray diffraction peak of the hcp 1 phase using Scherrer’s equation. It should be noted that only the film deposited at 2 2.0 V versus SCE showed the (100) peak. Based on this limited result, it can be said that the grain size of the 1 phase decreased rapidly from the 55 nm of pure Co to 15 , 30 nm with increasing Ni content. In the composition range, where both the a and 1 phases coexisted, the grain size of both phases was very small. Finally, the type of substrates used also affected the grain size. Implications of these results will be further explored using SEM and TEM. 7.4.2.16.3.3 Film microstructure on stainless steel (SUS-304) substrate Figure 7.4.2.16.5 are plan-view TEM micrographs of plated Co – Ni alloy films grown on stainless steel substrates. In preparing TEM specimens, the plated Co –Ni alloy films were first peeled off the stainless substrates and then electropolished. Figure 7.4.2.16.5(A) (a) is the bright-field (BF) image of pure Co, which contains many orderly features displaying banded contrast. The selected-area electron diffraction pattern (cf. Figure 7.4.2.16.5(A) (c)), in which many streaks are seen to emanate around the diffraction spots, has revealed the presence of both a-Co and 1-Co in this film. Various sizes crystals ranging from several 10 to 200 nm were found in the darkfield (DF) images (see Figure 7.4.2.16.5(A) (b)). The banded contrast was found to originate from the alternating thin layers of a-Co and 1-Co within one crystal. The calculated grain size, 22 nm, of a-Co in the pure Co film shown in Figure 7.4.2.16.4 thus represents the average size of each thin layer inside one large grain observed in the DF image (cf. Figure 7.4.2.16.5(A) (b)). The 20 at.% Ni alloy film contained a mixture of a and 1 phases. Both the BF and DF images indicated image contrast features similar to those of pure a-Co (see Figure 7.4.2.16.5(B)). Many alternating thin layers of a-Co and 1-Co crystals were also found in this film. The BF image of the 38 at.% Ni film also showed the orderly banded contrast (see Figure 7.4.2.16.5(C) (a)). The electron diffraction pattern, however, revealed almost no diffraction lines from 1-Co and no streaks around the spots (cf. Figure 7.4.2.16.5(C) (c)). The 50 at.% Ni film contained only the fcc a phase of , 100 nm in size and the orderly banded contrast disappeared (see Figure 7.4.2.16.5(D)). The 88 at.% Ni film showed the same electron diffraction pattern as the 20 and 50 at.% Ni films. The diffraction ring was spottier, indicating that the grains of the a phase were larger. Indeed, both the BF and DF micrographs showed large grains (see Figure 7.4.2.16.5(E) (a) and (b)). Finally, the structure of a plated pure Ni film is seen in the BF/DF images of Figure 7.4.2.16.5(F) (a) and (b). The grain size became even larger than that of the 88 at.% Ni film, and the corresponding electron diffraction pattern was indeed spotty (Figure 7.4.2.16.5(F) (c)).

488

Nano-Plating

Figure 7.4.2.16.6 shows SEM micrographs showing the surface morphology of (a) the 15 and (b) 51 at.% Ni films plated on stainless steel substrates at the deposition potential of 2 2.0 V versus SCE. Both the films displayed a similar granular surface morphology, which consisted of , several hundred-nm particles. As indicated in Figure 7.4.2.16.4, the grain size of these films was 20 , 30 nm and thus did not necessarily reflect the size of the

Figure 7.4.2.16.5. Plan-view TEM micrographs ((a) BF/(b) DF) and (c) electron diffraction patterns for plated (A) Co, (B) Co –20 at.% Ni, (C) Co –38 at.% Ni, and (D) Co –50 at.% Ni (E) Co– 88 at.% Ni alloy films, and (F) Ni film.



489

490

Nano-Plating



491

Figure 7.4.2.16.6. SEM micrographs showing the surface morphology of (a) the 15 and (b) 51 at.% Ni films plated on stainless steel substrates.

surface particles observed here by SEM. Similar surface morphologies were also observed in all other films over the entire composition range. 7.4.2.16.3.4 Film microstructure on polycrystalline copper substrate Figure 7.4.2.16.7(A) and (B) are SEM and TEM micrographs showing the surface morphology and cross-section of 28 and 72 at.% Ni alloy films plated on polycrystalline Cu substrates at 2 2.0 V versus SCE. In contrast to the surface morphologies observed on stainless steel substrates (see Figure 7.4.2.16.6), there are distinctively different smooth and granular regions on these surfaces. The size of particles in the granular region of the 28 at.% Ni film is smaller than the size of the 72 at.% Ni film particles. The presence of smaller surface particles in the 28 at.% Ni film may be related to the fact that the grain size was smaller since both the a and 1 phases co-existed in this film. It should, however, be remembered that the grain size of plated films does not necessarily reflect the surface roughness defined by the size of surface particles such as those seen in this study. TEM examination of the cross-sections has revealed that both the 28 and 72 at.% Ni alloy films grew uniformly from the deposit –substrate interface up to 1 mm (see Figure 7.4.2.16.7(A) (b) and (B) (b)). Black bend fringes running perpendicular to the interface in the Cu substrate side were seen to continue into the deposit side, indicating that these alloy films were oriented in the same direction as the Cu substrate and thus have grown epitaxially on the Cu substrate. The lattice constant, a, of Cu from the JCPDS file is 0.3615 nm while the constant of the a phase is 0.3519– 0.3555 nm, which was obtained from Figure 7.4.2.16.3. The lattice misfit between the Cu and the a phase amounts to 1:7 , 2:7%; which is small enough for the Co – Ni films to grow epitaxially on Cu (Watanabe, 1994). This explains why the grain size of the Co – Ni alloy films plated on Cu substrates was slightly larger than the films grown on

492

Nano-Plating

Figure 7.4.2.16.7. (a) An SEM image showing the surface morphology of plated (A) Co– 28 at.% Ni and (B) Co –72 at.% Ni alloy films, the cross-section TEM images ((b) BF and (c) DF images), and (d) the electron diffraction pattern.


493

Figure 7.4.2.16.8. (a) The equilibrium phase diagram of Co –Ni alloys and (b) the lattice constants of phases found in plated Co –Ni alloy films.

494

Nano-Plating

stainless steel substrates (see Figure 7.4.2.16.4). The surface of stainless steel substrates is generally covered with a layer of amorphous oxide film. Consequently, the film growth on such a surface is initiated by random nucleations without epitaxy, leading to the formation of fine grains. Contrary to the growth on the stainless steel substrates, the film growth on Cu substrates involved epitaxy, which allowed the plated films to replicate the grain structure of the Cu substrate. The grain size of the polycrystalline Cu as noted in the SEM micrographs of Figure 7.4.2.16.7 was very large (. 10 mm). An epitaxial growth on such a large-grained substrate, of course, produced large grains of Co – Ni alloy films (Figure 7.4.2.16.8).

REFERENCES Aotani, K. (1950) J. Jpn Inst. Metals, B-14, 55. Fukai, Y. (1985) J. Jpn Inst. Metals, 24, 671. Fukumuro, N., Chikazawa, M. & Watanabe, T. (1999) J.Surf. Finish. Soc. Jpn, 50, 441. Jartych, E., Olchowik, J., Zurawicz, J.K. & Budzynski, M. (1993) J. Phys.: Condens. Matter, 5, 8921. Nakahara, S. & Mahajan, S. (1980) J. Electrochem. Soc., 127, 283. Watanabe, T. (1994) J. Surf. Sci. Soc. Jpn, 15, 637.

7.4.2.17.

Electrolytic Co– Sn (Imafuji, 1999)


The compositions of the Co – Sn alloy plating bath used in this study, which are based on chloride salts, are listed in Table 7.4.2.17.1. Since SnCl2 salt does not dissolve well in water, potassium pyrophosphate or sodium citrate was added as a complexing agent. The bath pH was adjusted with sodium hydroxide for the citrate bath and with dilute hydrochloric acid for the pyrophosphate bath. Splat-quenched Fe – Ni– B –Mo amorphous foils (manufactured by Nippon Amorphous Metals Co., Ltd.), rolled polycrystalline Cu sheets, and stainless steel foils (SUS304) were used as the substrates.


A change in the Co content of electrodeposited Co – Sn alloy films obtained from the 30 8C citrate bath ðpH ¼ 2:0Þ is shown with respect to current density in Figure 7.4.2.17.1. The Co content appears to increase with increasing current density. The graph in Figure 7.4.2.17.2 is similar to Figure 7.4.2.17.1, except that the bath pH was 12.0 and the Co concentration of the bath varied. The Co content generally increases with increasing current density. However, if the Co concentration is less than 0.095 mol/l, alloying with Co does not occur in the film. On the other hand, if more than 0.195 mol/l Co is present in the bath, the Co content does not change significantly with increasing current density. In the case of a Co concentration of 0.145 mol/l, the Co content changes markedly with current density and thus various Co –Sn alloys can be readily obtained by varying the current density.

Table 7.4.2.17.1. The composition of Co –Sn alloy plating baths. Bath type

Composition

Citrate bath

CoCl2·6H2O SnCl2·2H2O C6H5Na3O7 Bath pH

0.195 mol/l 0.005 mol/l 1.0 mol/l 2.0 , 12.0

Pyrophosphate bath

CoCl2·6H2O SnCl2·2H2O K4P2O7 Bath pH

0.9 mol/l 0.1 mol/l 0.5 mol/l 10.0

495

496

Nano-Plating

Figure 7.4.2.17.1. A change in the Co content of plated Co–Sn alloy films obtained from the 30 8C citrate bath (pH ¼ 2:0 and Co2þ concentration ¼ 0:195 mol=l) with increasing current density.

Figure 7.4.2.17.2. A change in the Co content of plated Co–Sn alloy films obtained from the 30 8C citrate bath ðpH ¼ 12:0Þ with increasing current density. The Co2þ concentration of the bath was also changed.

Electrolytic Co – Sn

497

Figure 7.4.2.17.3. A composition (Co, Sn, and P) change with current density in electrodeposited Co– Sn alloys from the pyrophosphate bath (pH ¼ 10:0 and Co2þ concentration ¼ 0:9 mol=lÞ:

Figure 7.4.2.17.3 shows a composition change with current density in electrodeposited Co – Sn alloys from the pyrophosphate bath ðpH ¼ 10:0Þ (see Table 7.4.2.17.1 for the bath composition). Similar to the case of the citrate bath, the Co content of the film also increases with increasing current density.


The surface morphology of three alloy deposits ((a) Co – 87.4 at.% Sn, (b) Co – 58.8 at.% Sn, and (c) Co – 30.0 at.% Sn) from the citrate bath is shown in Figure 7.4.2.17.4. These

Figure 7.4.2.17.4. The surface morphology of three alloy deposits ((a) Co –87.4 at.% Sn, (b) Co –58.8 at.% Sn, and (c) Co –30.0 at.% Sn) from the citrate bath.

498 Nano-Plating Fig. 7.4.2.17.5. (A) Cross-section TEM micrographs and (B) electron diffraction patterns for plated (a) Co–94.2 at.% Sn, (b) Co –80.0 at.% Sn, (c) Co –60.0 at.% Sn, (d) Co–37.6 at.% Sn, and (e) Co –20.4 at.% Sn alloy films obtained from the citrate bath.

Electrolytic Co – Sn

499

Fig. 7.4.2.17.6. Cross-section TEM micrograph (a) and electron diffraction patterns ((b) & (c)) for a plated Co –Sn alloy film obtained from the citrate bath.

deposits initially started by forming a flat and smooth film. With increasing film thickness, the surface of the deposits became rough through the development of granular surface irregularities or by forming dendrites. Cross-section TEM images and electron diffraction patterns for these deposits are shown in Figure 7.4.2.17.5(A) and (B), respectively. The TEM micrographs indicate that the deposits are an assembly of fine grains. The cross-section of a Co –Sn alloy film from the citrate bath, whose pH was 2.0, is magnified in Figure 7.4.2.17.6. The initial flat and projecting parts can be seen in the micrograph. The grain size of the flat part is extremely fine and the Sn content is high. The grain size of the projecting part (or a protrusion) is larger and the Sn content is low. It is noted that each protrusion is an assembly of fine grains. Figure 7.4.2.17.7 is the same as Figure 7.4.2.17.4 except that the bath pH was 12.0. The surface of these deposits is seen to be smoother than the one shown in Figure 7.4.2.17.4.

500

Nano-Plating

Fig. 7.4.2.17.7. The surface morphology of various compositions of plated Co– Sn alloy films grown from the citrate bath.

REFERENCE Imafuji, K. (1999) Microstructure of electrodeposited Co –Sn alloy film, Master thesis, Tokyo Metropolitan University.

7.4.2.18.

Electrolytic Co– W (Ito et al., 2003)


Co – W alloy plating baths (see Table 7.4.2.18.1) were prepared from a mixture of CoSO4·7H2O and Na2WO4·2H2O salts together with 0.260 mol/l Na3C6H5O7·2H2O as a complexing agent. The mixing ratio of CoSO4·7H2O and Na2WO4·2H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.260 mol/l. The bath pH was adjusted to 8.5 with ammonium water. The plating was conducted for 30 min in 70 8C baths. Various compositions of Co – W alloy films were obtained by changing the molar ratio of the two salts or by changing the current density. An analysis of the alloy composition was done using an energy dispersive X-ray spectrometer attached to an SEM. A platinum plate was used as an anode. Rolled polycrystalline sheets were first electropolished and then employed as the substrates.


Figure 7.4.2.18.1 is a graph of the W content of plated Co – W films plotted against the W concentration in the baths. The plating was conducted at a current density of 400 A/m2. The W content of the films generally increased with increasing W concentration in the baths. Figure 7.4.2.18.2 depicts the relationship between the W content in plated Co – W films and the current density. The W content increased with increasing current density up to 300 A/m2 but did not increase above 300 A/m2. Although the W content could be changed by varying the current density, the maximum attainable W content in the films was 26.0 at.% W.

Table 7.4.2.18.1. The composition of a Co–W alloy plating bath and its plating condition. Composition

Plating condition

CoSO4·7H2O þ Na2WO4·2H2O ¼ 0.26 mol/l (NH4)2C4H4O6 ¼ 0.26 mol/l

Bath temperature: 70 8C Current density: 100–1200 A/m2 Plating time: 30 min Agitation: yes Substrate: polycrystalline Cu sheet Anode: Pt sheet

501

502

Nano-Plating

Figure 7.4.2.18.1. A graph for the W content of plated Co –W films plotted against the W concentration in the baths. The current density was 400 A/m2.

Figure 7.4.2.18.2. Relationship between the W content of plated Co –W films and the current density.

Electrolytic Co– W

503


Figure 7.4.2.18.3 shows X-ray diffraction patterns for plated Co –W alloy films having various W contents. Diffraction peaks from the Co appear to decrease with increasing W content. The location of the {101} Co peak was carefully measured with increasing W content, but did not show any peak shift, implying that the incorporated W atoms did not dissolved in the Co matrix. From this trend observed in Figure 7.4.2.18.3, it seems that the films containing less than 24.8 at.% W are crystalline but those with more than 25.4 at.% W are amorphous. It is not clear by this X-ray analysis alone whether or not the

Figure 7.4.2.18.3. X-ray diffraction patterns for plated Co –W alloy films containing various amounts of W.

504

Nano-Plating

Co – 25.4 at.% W alloy film is really amorphous. To further understand the amorphous structure, we conducted a heating (200 – 600 8C) experiment. Figure 7.4.2.18.4 depicts a change in X-ray diffraction patterns with increasing temperatures for a plated Co – 21.6 at.% W alloy film. No new diffraction peaks appeared, leaving the strong {101} Co peak unchanged by this heating. This particular alloy film already formed a thermally stable phase. This film has a {101} texture, which is also stable by heat treatment. Figure 7.4.2.18.5 depicts a change in X-ray diffraction patterns with increasing temperatures for a plated Co –24.8 at.% W alloy film. The as-deposited film exhibited

Figure 7.4.2.18.4. A change in X-ray diffraction patterns with increasing heat treatment temperature for a plated Co–21.6 at.% W alloy film.


505

Figure 7.4.2.18.5. A change in X-ray diffraction patterns with increasing heat treatment temperature for a plated Co –24.8 at.% W alloy film.

a broad peak with a small {101} Co peak. Heating the film from 200 to 600 8C promoted a reduction in the intensity of the broad peak and increased the intensity of the {101} Co peak, accompanied by the appearance of other Co-related diffraction peaks. It is clear that this alloy film is constructed from a mixture of an amorphous phase and {101}-oriented fine Co crystals. In this heat treatment, the {101}-oriented Co grains grew larger and new Co grains nucleated at the expense of the amorphous phase. Figure 7.4.2.18.6 depicts a change in X-ray diffraction patterns with respect to increasing temperatures for a plated Co –25.4 at.% W alloy film. This as-deposited film shows a very broad peak, which did not change by heating at temperatures up to 500 8C. At 600 8C, however, the peaks of the intermetallic compound, Co3W, suddenly appeared.

506

Nano-Plating

Figure 7.4.2.18.6. A change in X-ray diffraction patterns with increasing heat treatment temperature for a plated Co–25.4 at.% W alloy film.

In the case of the Co – 25.4 at.% W film, the as-deposited film is uniformly amorphous and when the heating temperature reached the crystallization temperature of Co3W, Co3W crystals nucleated. Figure 7.4.2.18.7 depicts a change in X-ray diffraction patterns with respect to increasing temperatures for a plated Co-26.0 at.% W alloy film. This film behaved in the same manner as the Co – 25.4 at.% W. Electrolytic Co – W alloy films, whose W content is less than 24.8 at.%, are considered to be crystalline because they consist of a mixture of an amorphous phase and fine Co crystals. The structure of those alloy films containing more than 25.4 at.% W were uniformly amorphous. The critical composition, at which the structure


507

Figure 7.4.2.18.7. A change in X-ray diffraction patterns with increasing heat treatment temperature for a plated Co-26.0 at.% W alloy film.

changes from crystalline to amorphous phases, was in the region of Co – 25.1 at.% W. The compositional range of crystalline and amorphous phases observed in the plated Co –W films are listed in Figure 7.4.2.18.8(b) and contrasted with the equilibrium phase diagram in Figure 7.4.2.18.8(a). The alloys containing 24.5 at.% W and 26.0 at.% W are close to the intermetallic compound, Co3W (25.0 at.% W). The structure of the films having a W content above 26.0 at.% are not well understood at the moment. Figure 7.4.2.18.9 is an SEM micrograph showing the fracture surface and surface morphology of an amorphous Co –W film obtained by a plating method. The surface

508

Nano-Plating

Figure 7.4.2.18.8. (a) The equilibrium phase diagram of Co –W alloys and (b) the phase diagram of plated Co –W films.


509

Figure. 7.4.2.18.9. An SEM micrograph showing the fracture surface and surface morphology of a plated amorphous Co–W film.

exhibits characteristic nodules, whereas the crack surface indicates the brittle fracture mode unique to amorphous materials.

REFERENCE Ito, K., Wang, F., Watanabe, T. (2003) J. Jpn Inst. Metals, 67, 499.

7.4.2.19.

Electrolytic Cr –H (Furuya, 1981 –1983)


The compositions of the Cr –H alloy plating bath used in this study, and its plating conditions, are listed in Table 7.4.2.19.1. The concentration of Cr3þ listed in the table was determined using a light absorption method. For this measurement, 1 cm3 of solution was extracted from the bath prior to the plating. A mixture of 0.2 kg CrO3 and 0.0053 kg (NH4)2SO4 (ammonium sulfate) was dissolved in 700 ml distilled water, which was subsequently heated to boiling. Ammonium citrate dibasic ((NH4)2HC6H5O7) (0.3 kg) was slowly added to the solution to reduce a proportion of Cr6þ ions to Cr3þ. After the addition of ammonium citrate dibasic, an additional 20-min heating was applied to complete the reduction process. After cooling, distilled water was added to make a 1000 ml solution, which became the E bath as shown in Table 7.4.2.19.1. The E bath, containing 50% Cr3þ, was mixed with the original mixture (200 kg/m3 CrO3 and 5.3 kg/m3 (NH4)2SO4, with no Cr3þ) in an appropriate ratio. The mixed solutions were heated for 10 min at 70 8C, and then cooled. By doing this, four baths (A, B, C, and D) having different Cr3þ concentration were prepared. The pH of the E bath was adjusted to 8.0 with the addition of ammonium water, but no pH adjustment was done for the other four baths. The plating was conducted for 4 h using a platinum plate ð25 £ 20 mm2 Þ as an anode and a polycrystalline Cu sheet ð25 £ 20 mm2 Þ as a cathode in a 300 ml beaker containing 260 ml of solution. The quality of plated Cr– H alloy films depended on the concentration of Cr3þ. Consequently, it was necessary to adjust the current density and the bath temperature to obtain high-quality Cr– H films. A current density of 1000 A/m3 and a bath temperature of 20 8C were chosen for the A, B, and C baths. For the bath D, current densities of 5000 – 21,000 A/m3 and the bath temperature of 60 8C were selected. Finally, the E bath was operated at 10,000 –25,000 A/m3 and 60 8C.

Table 7.4.2.19.1. The composition, plating condition, and current efficiency of Cr–H alloy plating baths. Bath name Cr3þ (%) CrO3 (kg/m3) AS (kg/m3) ACD (kg/m3) CD ( £ 104 A/m2) pH BT (8C) CE (%) A B C D E

5 15 18 40 50

200 200 200 200 200

5.3 5.3 5.3 5.3 5.3

4.6 15 30 150 300

0.1 0.1 0.1 0.5 –2.1 1.0 –2.5

0.5 0.5 0.6 3.0 8.0

20 20 20 60 60

22.4 25.1 5.4 0.09 0.03

AS ¼ ammonium sulfate ((NH4)2SO4); ACD ¼ ammonium citrate dibasic ((NH4)2HC6H5O7); CD ¼ current density; BT ¼ bath temperature; CE ¼ current efficiency.

510

Electrolytic Cr– H

511

It is important to note that these baths can be modified to plate Cr– W – H and Cr– Mo – H alloy films. For example, a bath for plating Cr – W – H alloy films was prepared in the following manner. During the construction of the E bath, at the reduction stage by the addition of ammonium citrate dibasic, a solution containing 0:025 , 0:1 kg WO3 and 0:025 , 0:1 kg NaCO3 was added, making 1000 ml solution. In the case of preparing a Cr –Mo alloy plating bath, again at the reduction stage by the addition of ammonium citrate dibasic, 5 , 40 kg=m3 MoO3 and 5 , 40 kg=m3 NaCO3 was added. The plating condition for Cr –W –H and Cr– W – H alloy films was the same as the E bath. The plating time for Cr –W – H films was 2 h, whereas that for Cr– Mo – H films was 4 h.


The current efficiency varied with the type of the bath, as seen in Table 7.4.2.19.1. With increasing concentrations of Cr3þ, the current efficiency increased initially but then dropped to an extremely small value. Figure 7.4.2.19.1 shows the structure of plated Cr –H alloy films obtained from the five baths listed in Table 7.4.2.19.1. The structure of Cr –H films is depicted by means of (a) bright-field (BF) and (b) dark-field (DF) images together with their electron diffraction pattern. Electron diffraction patterns from the films obtained from the A and B baths displayed single-crystal spots, indicating the formation of largegrained Cr. The film from the C bath showed spotty rings, which indicate a polycrystalline deposit. The diffraction pattern of the film from the D bath consists of a mixture of spotty and continuous broad rings. Both BF and DF images revealed the presence of 10-nm cuboidal crystals, which were identified to be Cr metal. The diffraction pattern of the film from the E bath was a broad ring characteristic of an amorphous phase. No particular image contrast was found from the BF and DF images. Figure 7.4.2.19.2 is a SEM micrograph showing the surface morphology of plated Cr– H alloy films, which were examined by TEM as seen in Figure 7.4.2.19.1. The films obtained from (A) the A and (B) B baths displayed rough surfaces, whereas those from (C) the C bath exhibited needle-shaped crystals. From the D and E baths, the films contained surface pits but their surfaces were generally smooth. Figure 7.4.2.19.3 is a TEM micrograph and electron diffraction patterns showing the structure of Cr – H films plated at current densities of (A) 10,000, (B) 18,000, and (C) 21,000 A/m2. Images marked with (a) –(c) in Figure 7.4.2.19.3 represent the bright-field image, dark-field image, and electron diffraction pattern for each film, respectively. At low current density (10,000 A/m2), the film grew as an amorphous phase. However, with increasing current density, cubic crystals appear to nucleate inside the amorphous matrix. Similar experiments were conducted in the E bath. For all the current densities, all the films became amorphous.

512

Nano-Plating

Figure 7.4.2.19.1. The structure of plated Cr– H alloy films obtained from five baths listed in Table 7.4.2.19.1. (a) The BF image, (b) the DF image, and (c) the electron diffraction pattern.


513

Figure 7.4.2.19.2. SEM micrographs showing the surface morphology of Cr– H alloy films plated at different current densities from five baths (see Table 7.4.2.19.1).

514 Nano-Plating Figure 7.4.2.19.3. TEM micrographs ((a) BF and (b) DF), and (c) electron diffraction patterns showing the structure of Cr–H films plated at current densities of (A) 10,000, (B) 18,000, and (C) 21,000 A/m2 from the D bath (40% Cr3þ).

Electrolytic Cr– H Figure 7.4.2.19.4. The weight gain, current efficiency, and structure of Cr–H films plated from (a) the D (40% Cr3þ) and (b) E (50% Cr3þ) baths plotted as a function of current density.

515

516 Nano-Plating

Figure 7.4.2.19.5. The surface morphology of Cr–H deposits plated at various current densities from the D (40% Cr3þ) and E (50% Cr3þ) baths.

517


Table 7.4.2.19.2. Impurity content in plated Cr and Cr-related alloy deposits (symbol X denotes W, Mo, or Fe). Plated films Crystalline Cr Non-crystalline Non-crystalline Non-crystalline Non-crystalline Non-crystalline

Cr (40% Cr3þ bath) Cr (50% Cr3þ bath) Cr–W Cr–Mo Cr–Fe

H (at.%)

N (at.%)

O (at.%)

Cr (at.%)

X (at.%)

2.81 14.1 20.1 25.7 22.2 17.5

0.05 0.16 0.17 0.17 0.24 0.19

0.89 0.79 0.76 0.87 0.80 0.81

96.2 85.0 78.6 67.2 74.8 79.7

– – – 6.06 1.92 1.82

The weight gain and current efficiency of Cr –H films plated from (a) the D and (b) E baths were plotted in Figure 7.4.2.19.4 as a function of current density. At low current densities, the weight gain was small in both baths, making the corresponding current efficiency low, and the films became amorphous, i.e. when the current efficiency was less than 0.009%, the films were amorphous. In comparison with the general alloy plating systems, these current efficiency values are extremely small. Consequently, the hydrogen evolution is violent and the heat generation is high during the plating. The high heat generation made it necessary to cool the plating bath externally. It is clear that the violent hydrogen evolution and subsequent incorporation caused the formation of Cr –H alloys. In Table 7.4.2.19.2, the concentrations of H, N, and O in plated Cr alloy deposits are listed. Amorphous Cr –W, Cr –Mo, and Cr –Fe deposits are listed for comparison. The hydrogen content was determined using an EMGA-521 hydrogen analyzer (Horiba, Ltd.). The amorphous Cr– H films are seen to contain much more hydrogen than the crystalline deposits. The surface morphology of these Cr – H deposits grown from the D and E baths can be observed in Figure 7.4.2.19.5, in which a number of surface pits are present in all the films. These surface pits are most likely vent holes for hydrogen gas to escape. Nevertheless, all the surfaces appear bright to the naked eye. The crystallization temperature of amorphous Cr– H deposits was greater than 370 8C and became higher if the deposits were further alloyed with Mo or W. The corrosion resistance of amorphous Cr –H deposits was superior to that from the crystalline phase.

REFERENCES Furuya, H (1983) A study on the formation, microstructure, crystallization behavior, and corrosionresistance property of electrolytic chromium and amorphous chromium binary alloys, PhD. dissertation, Tokyo Metropolitan University. Furuya, H. & Tanabe, Y. (1982) J. Jpn Inst. Metals, 46, 1042. Furuya, H., Misaki, Y. & Tanabe, Y. (1981) J. Metal Finish. Soc. Jpn, 32, 631. Furuya, H., Hasegawa, N., Misaki, Y. & Tanabe, Y. (1981) J. Metal Finish. Soc. Jpn, 32, 637. Furuya, H., Hasegawa, N., Shinohara, S. & Tanabe, Y. (1983) J. Mater. Sci. Soc. Jpn, 19, 341.

7.4.2.20. Electrolytic Cu –Ni (Mizushima, 1996) 7.4.2.20.1. PLATING METHOD

A Cu –Ni alloy plating bath is a mixed solution of nickel sulfate (NiSO4·6H2O) and copper sulfate (CoSO4·7H2O). The mixing ratio of CoSO4·7H2O and NiSO4·6H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.26 mol/l. Glycine was added to the bath at a concentration of 0.8 mol/l as a buffer. The bath pH was adjusted using NaOH. Polycrystalline Cu or splatquenched amorphous Fe – Ni– Si – B alloy foils were used as substrates. The plating was carried out at 100 A/m2 under a constant current condition in the stirred 60 8C bath. The plating was stopped when 20 C was consumed. The film thickness was , 2.5 mm.


Figure 7.4.2.20.1 is a plot of the Ni content of plated Cu –Ni alloy films against the Ni concentration in the bath. No Cu – Ni alloys containing Ni were produced from a bath having a Ni concentration of less than 60 at.%. From the baths containing more than 60 at.% Ni, Cu –Ni alloy films, whose Ni content changed with the Ni concentration of the bath, were obtained. Figure 7.4.2.20.2 illustrates a change in the Ni content of plated Cu – Ni films with a change in the bath pH. Three baths having three concentration ratios of Cu and Ni (Cu:Ni ¼ 0.05:0.95, 1:9, and 2:8) were studied.


X-ray diffraction patterns for various compositions of plated Cu –Ni alloys are shown in Figure 7.4.2.20.3. The bath pH was 6. All the patterns indicate the fcc structure. Lattice constants were calculated from these diffraction patterns and plotted in Figure 7.4.2.20.4(b). The equilibrium phase diagram of Cu –Ni alloys is shown in Figure 7.4.2.20.4(a) for comparison. According to the equilibrium phase diagram, the alloy is mutually soluble over the entire composition range at high temperatures but indicates phase separation at temperatures below 400 8C. In Figure 7.4.2.20.4(b), the composition of the Cu –Ni alloy films was varied by changing the bath pH. The lattice constant changed almost linearly from that of Ni (0.3523 nm (JCPDS File #4-0850)) to that of Cu (0.3616 nm (JCPDS File #4-0836)), following Vegards law. These alloys are considered to be solid solutions without any phase separation (Massalski, 1990). 518

Electrolytic Cu– Ni

519

Figure 7.4.2.20.1. A graph of the Ni content of plated Cu –Ni alloy films plotted against the Ni concentration in the bath.

Figure 7.4.2.20.2. A change in the Ni content of plated Cu –Ni films with a change in the bath pH.

520

Nano-Plating

Figure 7.4.2.20.3. X-ray diffraction patterns for various compositions of plated Cu –Ni alloys. The bath pH was 6.

X-ray diffraction patterns were taken from a plated Cu –75 at.% Ni alloy film, which was annealed at 100 8C for 120 and 360 min. Although grain coarsening took place during this annealing, as indicated by the sharpening of the diffraction lines, no phase separation was observed (see Figure 7.4.2.20.5). Similar annealing treatments at 200 and 300 8C were given, but these did not induce any phase separation. In Figure 7.4.2.20.6, we present a cross-section TEM micrograph of the plated alloy film (Cu – 75 at.% Ni), which does not show any layer structure, indicating no compositional fluctuation along the film thickness direction. The absence of the compositional fluctuation was also confirmed using an energy dispersive X-ray spectrometer attached to the TEM. The cross-section of the same alloy film (Cu –75 at.% Ni), which was annealed at 100 8C, was examined by TEM as shown in Figure 7.4.2.20.7. This film did not contain a layer structure and its electron diffraction pattern became spottier, indicating a grain coarsening. Furthermore, the diffraction rings did not split and thus the film was a single-phase solid solution alloy. Kurachi and Sumiyama (1973) found that plated Cu – Ni alloys form a solid solution over the whole composition range and their lattice constants follow Vegards law. Figure 7.4.2.20.8 depicts the surface morphology of plated Cu –Ni alloy films. The bath pH was 6. The surface was generally smooth for all the alloy compositions.


521

Figure 7.4.2.20.4. (a) The equilibrium phase diagram of Cu–Ni alloys and (b) the lattice constants of plated Cu – Ni alloys.

522

Nano-Plating

Figure 7.4.2.20.5. X-ray diffraction patterns showing a structural change in a plated Cu–74 at.% Ni alloy film upon heating at 100 8C for different times.

Figure 7.4.2.20.6. A cross-section TEM micrograph of a plated Cu –75 at.% Ni alloy film. Three electron diffraction patterns were taken at three locations, (a)–(c), indicated in the TEM micrograph.


523

Figure 7.4.2.20.7. A cross-section TEM micrograph of a plated Cu –75 at.% Ni alloy film, which was annealed at 100 8C for 360 min. Three electron diffraction patterns were taken at three locations, (a)–(c), indicated in the TEM micrograph.

524

Nano-Plating

Figure 7.4.2.20.8. Surface morphologies for plated Cu –Ni films having various alloy compositions.

REFERENCES JCPDS File #4-0850. JCPDS File #4-0836. Kurachi, M. & Sumiyama, B. (1973) Denki Kagaku, 41, 26. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, Vol. 2, ASM International, p. 1442. Mizushima, I., Chikazawa, M. & Watanabe, T. (1996) J. Electrochem. Soc., 143, 1978.

FURTHER READING Ishikawa, M. & Enomoto, H. (1980) J. Metal Finish. Soc. Jpn, 31, 545. Ishikawa M. & Enomoto H (1986) Extended Abstract for the 74th Meeting of Metal Finish Soc. of Japan. p. 6. Sanada, Y.N. & Venkatachalam, T.K. (1979) Metal Finish., 77, 15. Stout, L., Burch, O.G. & Langdorf, A.S. (1930) Trans. Electrochem. Soc., 57, 113. Vu Quang, K., Chassaing, E., Viet, B. Le. Celis, J.P. & Roos, J.R. (1985) Metal Finish., 83(October), 25. Ying, R.Y. (1988) J. Electrochem. Soc., 135, 2957.

7.4.2.21. Electrolytic Cu– Pb (Shimizu, 1978) 7.4.2.21.1. PLATING METHOD

The compositions of the Cu – Pb alloy plating baths used in this study are listed in Table 7.4.2.21.1. The bath was prepared based on a composition used by Piontelli et al. (1967), who used a bath containing tartaric acid, sulfamic acid, and citric acid. The mixing ratio of Cu – to-Pb was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.2 mol/l. Polycrystalline Cu sheets were used as substrates. Electron-transparent Au specimens for TEM observations were made by the method illustrated in Figure 5.5, Chapter 5. The use of electron-transparent Au substrates allowed us to observe both the deposit and the substrate simultaneously in a plan-view mode. For a structural analysis of Cu – Pb alloy films, both X-ray diffraction and high-resolution reflection electron diffraction techniques were applied. After the plated films were dissolved in nitric acid, a composition analysis was conducted using an atomic absorption method.


Figure 7.4.2.21.1 shows a graph for the Pb content of plated Cu –Pb alloy films plotted against the Pb concentration in the baths. The current density was 10 A/m2 and the plating duration was (a) 10 s, (b) 60 s, and (c) 10 minutes. Film (a) was used for TEM examinations, whereas films (b) and (c) were used for surface morphology studies. Figure 7.4.2.21.1 indicates that the plating time affects the alloy composition significantly. This phenomenon is common in alloy plating systems that contain two metal ions with markedly different deposition potentials. In such systems, the nobler metal deposits preferentially at the start of plating, followed by the formation of a uniform bulk alloy layer. The layer, whose composition is richer in the noble metal, is called an initial layer (Watanabe, 1990). In the Cu – Pb system, Cu is nobler than Pb and thus a Cu –rich film is expected to form as the initial layer. As seen in Figure 7.4.2.21.1, the thinnest film plated for 10 s contained the highest concentration of Cu among the films deposited for 60 s and 10 min.


Figure 7.4.2.21.2 shows X-ray diffraction patterns for Cu –Pb alloy films plated for 10 min. Using the X-ray diffraction patterns together with the high-resolution reflection electron diffraction patterns taken from the films plated for 60 s, we determined the lattice constants 525

526

Nano-Plating

Table 7.4.2.21.1. The composition of Cu–Pb alloy plating baths. Composition

PbO CuO HS3NH2 K3C6H5O7·H2O K2C4H4O6·1/2H2O KOH

Bath name Cu

A

B

C

D

E

F

G

Pb

0 24.0 70 100 70 70

22.3 15.9 70 100 70 70

33.5 11.9 70 100 70 70

44.6 9.7 70 100 70 70

50.3 6.0 70 100 70 70

55.8 4.0 70 100 70 70

60.9 2.1 70 100 70 70

63.8 1.2 70 100 70 70

67.0 0 70 100 70 70

Figure 7.4.2.21.1. Relationship between the Pb content of plated Cu –Pb alloy films and the Pb concentration in the baths.

Electrolytic Cu – Pb

527

Figure 7.4.2.21.2. X-ray diffraction patterns for plated Cu –Pb alloy films.

over the whole composition range (see Figure 7.4.2.21.3(b)). The equilibrium phase diagram of Cu – Pb alloys (Piontelli et al., 1967) is shown for comparison in Figure 7.4.2.21.3(a). This alloy system is a peritectic type, in which two phases (a and b phases) are present at ambient temperatures. In addition, the composition of these a and b phases is very close to the composition of pure Cu and pure Pb, respectively. From pure Cu to 6 at.% Pb, only the a phase appeared, whereas from pure Pb to 18 at.% Cu, only the

528

Nano-Plating

Figure 7.4.2.21.3. (a) The equilibrium phase diagram of Cu –Pb alloys and (b) the lattice constants of plated Cu –Pb alloy films over the whole composition range.


529

b phase appeared (see Figure 7.4.2.21.3). On the pure Cu side, the lattice constant of the a phase initially increased with increasing Pb content. On the pure Pb side, by contrast the lattice constant of the b phase decreased with increasing Cu content. This trend strongly indicates that on both the Cu and Pb sides, solid-solution two-phase alloys were formed. The equilibrium phase diagram, however, does not show any mutual solubility at ambient temperatures. Thus, the solid solution observed here is supersaturated and can be in a metastable state. Using a different Cu – Pb alloy plating bath, Raub and Muller (1967) found that a supersaturated solid-solution alloy in the Cu side formed up to 34 at.% Pb and its lattice constant obeyed Vegards law up to 3 , 4 at:% Pb. In the composition range of 6 , 82 at:% Pb, both the a and b phases appeared, which were consistent with the equilibrium phase diagram. The lattice constant of the a phase was very close to that of pure Cu but that of the b phase became smaller than that of pure Pb, indicating that Cu dissolved into pure Pb. Figure 7.4.2.21.4 contains a plan-view TEM micrograph showing the microstructures and their electron diffraction patterns of electrolytic Cu – 30 at.% Pb alloy films grown for 1 and 3 s, respectively, on (a) the {110} and (b) the {100} surfaces of Au. In all cases, the films were an assembly of fine ð10 , 20 nmÞ crystals. In addition to strong diffraction spots from the Au substrate, diffraction spots from both Cu and Pb are seen in these diffraction patterns, indicating lattice matching to the Au substrate. Figure 7.4.2.21.5 displays SEM micrographs showing the surface morphology of various Ni –Pb alloy films. It is clear that the films containing high Pb content show rough surfaces.

Figure 7.4.2.21.4. Plan-view TEM micrographs showing the microstructure and its electron diffraction pattern of two composite deposit-substrate films, in which electrolytic Cu–30 at.% Pb alloy films were deposited on (a) the {110} and (b) the {100} surfaces of Au.

530 Nano-Plating

Figure 7.4.2.21.5. SEM micrographs showing the surface morphology of various Cu –Pb alloy films. The plating time for all the films was 10 min.


531

REFERENCES Piontelli, P., Cavallotti, P. & Guilianti, L. (1967) Electrochim. Metal, 2, 222. Raub, E. & Muller, K. (1967) Fundamentals of Metal Deposition, Elsevier Pub. Co., New York, p. 136. Shimizu, Y. & Tanabe, Y. (1978) J. Metal Finish. Soc. Jpn, 29, 21. Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 652.

7.4.2.22. Electrolytic Cu –Sb 7.4.2.22.1. FILM STRUCTURE

Regarding the structure of electrolytic Cu – Sn alloy films, Brenner (1963) cites results by Raub and Sautter (1957) as seen in Figure 7.4.2.22.1(b). Since no other reports are available for this alloy system, we list the observed phases in the electrodeposited films together with its equilibrium phase diagram.

Figure 7.4.2.22.1. (a) The equilibrium phase diagram of Cu –Sb alloys and (b) the phase diagram of plated Cu – Sb alloys (Raub and Sautter, 1957; Brenner, 1963).

REFERENCES Brenner, A. (1963) Electrodeposition of Alloys, vol. I, Academic Press, New York, p. 194. Raub, E. & Sautter, F. (1957) Metalloberfla¨che, 11, 249.

532

7.4.2.23. Electrolytic Cu– Sn (Liang, 1999) 7.4.2.23.1. PLATING METHOD

The compositions of the Cu – Sn alloy plating bath used in this study are listed in Table 7.4.2.23.1. A Cu –Sn alloy plating bath is a mixed solution of tin sulfate (SnSO4) and copper sulfate (CuSO4·4H2O). The mixing ratio of SnSO4 and CuSO4·7H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.5 mol/l. To prevent hydrolysis of Cu2þ and Sn2þ, 1 mol/l H2SO4 was added to the bath. As a leveling agent, cresol sulfonic acid was added. It was confirmed that the addition of cresol sulfonic acid did not cause the incorporation of any atomic elements into Cu – Sn deposits and the amount of Sn content in the deposit did not change. Polycrystalline Fe sheets having dimensions of 40 £ 12 £ 0:2 mm3 were used as substrates after electropolishing to a mirror finish. The plating was carried out at 300 A/m2 for 30 min in the 30 8C unstirred bath. As described later, the current efficiency was almost 100% except for a few cases of 85%. Consequently, most of the films grew to a thickness of 60 mm after 30-min plating.


The Sn content of plated Cu –Sn alloy films and current efficiency are plotted in Figure 7.4.2.23.1 as a function of the Sn concentration in the bath. If the Sn concentration in the bath is less than 26 at.%, no Cu – Sn alloy deposits are obtained. Above 26 at.% Sn concentration, however, we can plate Cu – Sn alloy films. Based on this plot, we prepared various compositions of plated Cu – Sn alloy films, which were subsequently subjected to structural analysis.

Table 7.4.2.23.1. The composition of a Cu –Sn alloy plating bath. Composition


SnSO4 þ CuSO4·5H2O H2SO4 Cresol sulfonic acid

0.5 1.0 0.25

533

534

Nano-Plating

Figure 7.4.2.23.1. The Sn content of plated Cu –Sn alloy films and the current efficiency (CE) plotted as a function of the Sn concentration in the bath.


7.4.2.23.3.1 Structure Table 7.4.2.23.2 is a summary of previously reported stable and meta-stable phases in Cu – Sn alloys as well as new phases found in plated Cu –Sn films (Bernal, 1928; Carlsson and Ha¨gg, 1932; Kno¨dler, 1956; de Bondt and Deruyttere, 1967; Gangulee et al., 1973; Vandermeulen and Deruyttere, 1973; Brandon et al., 1975; Saunders and Miodownik, 1987; Massalski, 1990a). Figure 7.4.2.23.2(a) shows the equilibrium phase diagram of Cu – Sn alloys (Massalski, 1990b). Figure 7.4.2.23.2(b) and (c) will be described later. Figure 7.4.2.23.3 displays X-ray diffraction patterns from plated Cu – Sn films obtained in this study. X-ray diffraction patterns for Cu – Sn alloy films containing 3.4, 4.8, 6.6, 10.6, and 12.6 at.% Sn are shown in Figure 7.4.2.23.3(a)– (e), respectively. These five patterns are very similar except for the height of the peaks. For the Cu, the peak positions shift toward the low-angle side with increasing alloying of Sn. This peak shift suggests that the Cu is forming a solid solution alloy with Sn. According to the equilibrium phase diagram in Figure 7.4.2.23.2(a), although Sn can dissolve in Cu at high temperatures, there is no Sn solubility in Cu at ambient temperatures. Consequently, these Cu phases are meta-stable supersaturated solid-solution alloys existing at ambient temperatures. These alloys are

535

Electrolytic Cu – Sn Table 7.4.2.23.2. A list of stable and meta-stable phases observed in Cu –Sn alloys. (a) Phases listed in the JCPDS file Composition Name [1] 2 3 4 [5] 6 7 8 9 10 [11] 12 13 14 15

Cu a-(Cu, Sn) CuSn (15 at.% Sn) Cu5.6Sn b-(Cu, Sn) Cu41Sn11 Cu327.92Sn88.08 Cu81Sn22 Cu39Sn11 Cu10Sn3 1-Cu3Sn h-Cu6.25Sn5 h-Cu6Sn5 Cu6Sn5 b-Sn

a phase S.S.S.S. (b phase ?) (b phase ?) b phase (dphase ?) (dphase ?) (dphase ?) (dphase ?) (z phase ?) 1 phase (h phase ?) h phase h phase Sn

(b) Phases not listed in the JCPDS file Composition Name 1 2 3 4 5 6 7 8 9 10 [11] 12 (13) (14) (15)

8 at.% Sn 11 at.% Sn

25 at.% Sn

Crystal system

a (nm)

b (nm)

Cubic Cubic Tetragonal Tetragonal Orthorhombic Cubic Cubic Cubic Cubic Cubic Orthorhombic Hexagonal ? ? Tetragonal

0.36150 0.3655 0.3726 0.3726 0.456 1.7980 1.7964 1.7980 1.8011 0.733 0.5621 0.42062 ? ? 0.5821

Crystal system

a (nm)

cph, A3

0.2625

0.426

DO3 Ordered Hexagonal

0.899 1.799 0.7330

7.864

Tetragonal Tetragonal Orthorhombic

0.4192 2.085 2.095 2.0870 0.5985 0.4744 0.8612

0.5037 2.51 2.543 2.5081 1.1028 1.0062 0.4033

0.456

3.325 ? ?

b (nm)

c (nm)

JCPDS #

0.4328 0.5094 ? ? 0.3182

4-0836 44-1477 17-865 31-487 6-0621 30-510 30-511 31-486 31-485 26-564 1-1240 47-1575 2-713 45-1488 4-0673

c (nm)

Remark

0.3642 0.3642 4.32

0

a phase a0 phase b00 phase v phase g phase g0 phase z phase z0 phase h phase h0 phase h0 phase h0 phase T(I) T(II) Or

0.4190

New phase New phase New phase

A phase marked with “?” was discovered in this study, although it is not mentioned in the literature nor in the JCPDS file. [1], [5], [11], and [12] were determined in this study, whereas (13), (14), (15) are newly discovered meta-stable phases.

designated as Cu(Sn) in Figure 7.4.2.23.3. From these five patterns, we also discovered a new phase called T(I), which has not been previously reported. Assuming that these peaks originated from one unknown crystal, we found that the crystal was tetragonal with the lattice constants of a ¼ b ¼ 0:985 nm and c ¼ 1:1028 nm: This new crystal was further analyzed using TEM. Figure 7.4.2.23.4(a) is a bright-field image showing the microstructure of the Cu – 3.4 at.% Sn alloy film. Figure 7.4.2.23.4(b) and (c) are electron

536

Nano-Plating

Figure 7.4.2.23.2. (a) The equilibrium phase diagram of Cu –Sn alloys, (b) and (c) are the phase diagrams of plated Cu –Sn alloy films.

Electrolytic Cu – Sn

Figure 7.4.2.23.3. X-ray diffraction patterns from various compositions of plated Cu–Sn alloy films.

537

538

Nano-Plating

diffraction patterns from two new tetragonal crystals marked as (b) and (c) in Figure 7.4.2.23.4(a). Since this tetragonal crystal had not been reported in the literature, it is assumed to be a thermodynamically meta-stable phase, which will be called T(I). In summary, all the plated Cu – Sn alloy films, whose Sn composition ranged from 3.4 to 12.6 at.%, consist of Cu(Sn) and T(I) phases. Figure 7.4.2.23.3(f) –(h) shows X-ray diffraction patterns from plated Cu –Sn alloy deposits containing 13.4, 15.2, and 16.2 at.% Sn. The diffraction peaks of the 13.4 at.% Sn alloy film (Figure 7.4.2.23.3(f)) were found to correspond to that of a meta-stable b0 phase (JCPDS #6-0621). Broad diffraction peaks from a supersaturated solid-solution alloy of Cu overlap with those of the b0 phase. It was, therefore, impossible from X-ray diffraction patterns alone to judge the presence of the Cu solid-solution alloy. TEM was applied to resolve this peak overlap problem. An example of the TEM analysis is shown in Figure 7.4.2.23.5, where the bright-field image of the 13.4 at.% Sn alloy film, together with its

Figure 7.4.2.23.4. (a) A plan-view TEM micrograph of the Cu –3.4 at.% Sn alloy film and two electron diffraction patterns taken at two sites, (b) and (c).


539

electron diffraction patterns, is displayed. The bright-field image seen in Figure 7.4.2.23.5(a) consists of isolated white features enclosed by a dark background. Electron diffraction patterns from the white feature and the dark background are shown in Figure 7.4.2.23.5(c) and (b), respectively. The white feature and the dark background were identified as being the b0 phase and the Cu solid solution, respectively. Since the diffraction pattern of the Cu solid solution consists of rings, the crystal of the Cu solid solution is very small and is randomly oriented. From the above results, it can be summarized that the 13.4 at.% alloy film is made up of coarse-grained b0 phase and fine-grained Cu solidsolution alloy. The X-ray diffraction pattern of the 15.2 at.% Sn alloy film (cf. Figure 7.4.2.23.3(g)) resembles the pattern of the 13.4 at.% Sn alloy deposit (see Figure 7.4.2.23.3(f )). A TEM analysis of this film, however, revealed that the electron diffraction patterns contained two unknown crystals, described as (c) an orthorhombic crystal and (d) a tetragonal crystal, in addition to (b) broad rings from the Cu solid solution (see Figure 7.4.2.23.6). The lattice


540

Nano-Plating

constants of the orthorhombic crystal were found to be a ¼ 0:861 nm; b ¼ 0:419 nm; and c ¼ 0:403 nm: This orthorhombic crystal will be called Or. The lattice constants of the tetragonal crystal, which is different from the previously mentioned crystal, T(I), were a ¼ b ¼ 0:474 nm and c ¼ 1:006 nm: This crystal will be called T(II). Thus, the 15.2 at.% Sn alloy film was an assembly of three types of crystals, the Cu solid solution, Or, and T(II). The diffraction rings from the Cu solid-solution alloy were broader than the rings of the 13.4 at.% Sn alloy (see Figure 7.4.2.23.6(b)), indicating that the grain size of this solid solution alloy was even smaller than that alloy. An electron diffraction analysis of the 16.2 at.% Sn alloy film revealed that the film consisted of the Cu solid-solution alloy and the previously mentioned T(II) phase (see Figure 7.4.2.23.7). X-ray diffraction patterns shown in Figure 7.4.2.23.3(i) and (j), which were taken from the 20.5 and 23.1 at.% Sn alloy deposits, are very simple. According to an electron diffraction analysis, we found the T(II) phase and three equilibrium phases including h (JCPDS #47-1575), h0 , and 1 phases (JCPDS #1-1240). The h phase undergoes



541

an order – disorder transformation from the low-temperature phase of h0 to the hightemperature phase of h. The structure of both the h and h0 phases is hexagonal. The lattice constants of the h phase are a ¼ 0:42062 nm and c ¼ 0:50974 nm: According to Carlsson and Ha¨gg (1932), the lattice constants of the h phase are a ¼ 2:095 nm and c ¼ 2:543 nm; which are about five times larger than those of the h phase. Thus, the h0 phase is regarded as the superlattice of the h phase. The transformed phase of h was designated as h(h0 ) in Figure 7.4.2.23.3. It cannot be confirmed whether the 1 phase was formed during plating or during the ion-milling process used for TEM specimen preparation. The 20.5 and 23.1 at.% Sn alloy deposits were very unstable and low-temperature heating at 373 K (100 8C) transformed the entire film into the 1 phase (see Figure 7.4.2.23.9). Thus, it is conceivable that the film could have been heated during the ion-milling process and consequently part of the film may have been transformed into the 1 phase. An X-ray diffraction pattern for the 24.8 at.% Sn alloy deposit is shown in Figure 7.4.2.23.3(k). In this film, four phases including the h, h0 , T(II), and Sn phases were identified. Tin metal has two allotropic forms, i.e. gray tin (a-Sn) with a cubic structure and white tin (b-Sn) with a tetragonal structure. The Sn metal found in the 24.8 at.% Sn alloy film was b-Sn. The diffraction angle of the b-Sn shifted toward the higher angle side as the b-Sn alloyed with Cu. Cu cannot dissolve into the Sn metal at ambient temperature. b-Sn is a supersaturated solid solution and is thus meta-stable. The supersaturated solid solution of Sn will be designated as Sn(Cu), which is marked in the diffraction patterns in Figure 7.4.2.23.3 (k) – (m). X-ray diffraction patterns for the 45.7 and 69.3 at.% Sn alloy deposits are shown in Figure 7.4.2.23.3 (l) and (m). Both films were found to contain the h, h0 , and Sn(Cu) phases. Finally, an X-ray diffraction analysis of the film from the bath containing only the Sn salt revealed only the tetragonal b-Sn crystal in the film (see Figure 7.4.2.23.3(n)). In summary, we discovered the following four new meta-stable alloy phases in plated Cu – Sn alloy films, in addition to the known phases in the literature. (1) Tetragonal crystal (T(I)): the lattice constants are a ¼ b ¼ 0:985 nm and c ¼ 1:1028 nm: (2) Tetragonal crystal (T(II)): a ¼ b ¼ 0:474 nm and c ¼ 1:006 nm: (3) Orthorhombic crystal (Or): a ¼ 0:861 nm; b ¼ 0:419 nm; and c ¼ 0:403 nm: (4) A solid-solution alloy of Sn supersaturated with Cu (Sn(Cu)). We summarize the above results in Figure 7.4.2.23.2(b). This phase diagram appears to be quite different from the equilibrium phase diagram shown in Figure 7.4.2.23.2(a). To further clarify the origin of this marked difference, we shifted the phase diagram of plated Cu – Sn films (cf. Figure 7.4.2.23.2(b)) along the temperature axis on the equilibrium phase diagram (cf. Figure 7.4.2.23.2(a)) and discovered that these phases and their phase

542

Nano-Plating

Figure 7.4.2.23.7. Electron diffraction patterns from the meta-stable phase T(II), which appeared in a plated Cu –16.2 at.% Sn film.

boundaries of Figure 7.4.2.23.2(b) correspond to those of Figure 7.4.2.23.2(a) at 650 8C. Based on this, we made a pseudo phase diagram for the plated films as shown in Figure 7.4.2.23.2(c). In conclusion, the structure of plated Cu – Sn alloy deposits is considered to be formed during quenching from 650 8C.

Figure 7.4.2.23.8. Electron diffraction patterns from the (a) T(II), (b)1, (c) h, and (d) h0 phases of a plated Cu –20.5 at.% Sn alloy film.


543

Figure 7.4.2.23.9. X-ray diffraction patterns showing the structural change of a plated Cu –23.1 at.% Sn alloy film upon heat treatment at 100 8C.

REFERENCES Bernal, J.D. (1928) Nature, 122, 54. de Bondt, M. & Deruyttere, A. (1967) Acta Met., 15, 993. Brandon, J.K., Pearson, W.B. & Tozer, D.J.N. (1975) Acta Cryst., B31, 774. Carlsson, O. & Ha¨gg, G. (1932) Z. Krist., 83, 515. Gangulee, A., Das, G.C. & Bever, M.B. (1973) Met. Trans., 4, 2063. Kno¨dler, H. (1956) Acta Cryst., 9, 1036. Liang, H.-Y., Chikazawa, M. & Watanabe, T. (1999) J. Jpn Inst. Metals, 63, 454. Massalski, T.B. (1990a) Binary Alloy Phase Diagrams, 2nd Edition, vol. 3, ASM International, p. 2863. Massalski, T.B. (1990b) Binary Alloy Phase Diagrams, 2nd Edition, vol. 2, ASM International, p. 1481. Saunders, N. & Miodownik, A.P. (1987) J. Mater. Sci., 22, 629. Vandermeulen, W. & Deruyttere, A. (1973) Met. Trans., 4, 1659.

7.4.2.24. Electrolytic Cu –Zn 7.4.2.24.1. PLATING METHOD

The mixing ratio of CuCN and ZnSO4 salts was changed to obtain various Cu – Zn alloy compositions while keeping the total metal ion concentration in the 200 ml bath constant at 30 g/l. In addition, 10 g/l NaOH was added to the bath as a supporting electrolyte. The plating, which used a Zn sheet as an anode and a stainless steel sheet as a cathode, was done at a current density of 10 A/m2 for 3 h in the room-temperature bath. After plating, the film was peeled off the substrate. The Cu content of the alloy film was determined using a sodium thiosulfate method. A structural analysis of the alloy films was performed using a Debye camera.


Unfortunately precise data were not provided in Aotani (1951).


Phases and their lattice constants obtained from electrolytic Cu – Zn alloy films are listed in Figure 7.4.2.24.1(b) and compared with the corresponding equilibrium phase diagram in Figure 7.4.2.24.1(a). In the alloy films of Cu-(73 , 88 at.%) Zn, diffraction lines from the 1 or g phase (hcp) appeared, whereas for the Cu-(0 , 49) at.% Zn alloy films, an fcc phase corresponding to the a phase emerged. The Cu-49 at.% Zn alloy film was composed of only the a phase and did not show the b phase. In the alloy films having 88 at.% Zn, an hcp phase was dominant with only a small degree of bcc phase present. For the alloy films containing less than 71 at.% Zn, only hcp crystals were observed. Note that the b and g phases are bcc crystals of CsCl type and a-Mn type, respectively.

REFERENCE Aotani, K. (1951) J. Jpn Inst. Metals, B-15, 52.

544

Electrolytic Cu –Zn

545

Figure 7.4.2.24.1. (a) The equilibrium phase diagram of Cu– Zn alloys and (b) the lattice constants of phases observed in plated Cu– Zn alloy films.

7.4.2.25. Electrolytic Fe – Mo (Watanabe et al., 1989) 7.4.2.25.1. PLATING METHOD

The compositions of the acid and alkaline Fe – Mo alloy plating bath used in this study are listed in Table 7.4.2.25.1(a) and (b), respectively. For the acid bath, the mixing ratio of iron sulfate (FeSO4·7H2O) and sodium molybdate (Na2MoO4·2H2O) salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.26 mol/l. Sodium citrate (C3H4(OH)COONa3) was added as a complexing agent, and the bath pH was adjusted using dilute sulfuric acid or ammonium water. A detailed analysis of plated Fe – Mo alloy films was performed for the cases of bath pH ¼ 3 and 5. The bath temperature was 33 8C. A compositional analysis of the films was performed using an atomic absorption method. A structural analysis of the films was performed using an X-ray diffractometer, which used a cobalt Ka line as the irradiation source. The composition of the alkaline bath is seen in Table 7.4.2.25.1(b). The bath was an ammonium molybdate ((NH4)2Mo2O7) solution. The pH of the as-prepared bath was 14 and thus the bath was strongly alkaline.


Figure 7.4.2.25.1(A) and (B) are graphs for the Mo content in Fe –Mo alloy films obtained from the acid bath with pH ¼ 3 and 5 plotted as a function of the Mo concentration in the bath. Three current densities of 400, 600, and 800 A/m2 were applied. The dotted line in the graph represents the ratio of the Mo content of the film to the Mo concentration

Table 7.4.2.25.1. The composition of Fe–Mo alloy plating baths. Bath name

Composition (g/l)

Mo/(Mo þ Fe) (%) 31.6

(a)

FeCl2 Fe2þ (NH4)2Mo2O7

75 21 18

(b)

(NH4)2C5H5O7 NaOH C6H15O3N

40 140 50

546

53

62

66

76.3

81.5

30 8.4 18

40 11.2 34

75 21 75

45 12.6 75

35 9.8 80

Electrolytic Fe – Mo

547

Figure 7.4.2.25.1. (A) and (B) are graphs for the Mo content in Fe–Mo alloy films obtained from the acid bath (Bath (a) in Table 7.4.2.25.1) with pH ¼ 3 and 5 plotted as a function of the Mo concentration in the bath. Three current density values (400, 600, and 800 A/m2) were used.

of the bath when it was equal to 1. The Mo content appears to increase with increasing Mo concentration in the bath regardless of the magnitude of the current density. Figure 7.4.2.25.2 is a graph of the Mo content in the film plotted against the bath pH. These baths contained 60 mol/l sodium molybdate, and were operated at a current density of 400 A/m2. The Mo content is seen to display a maximum at pH ¼ 2:5: The maximum attainable Mo content was 55 at.%, which was obtained at high current densities from the pH ¼ 2 bath. Figure 7.4.2.25.3 is the Mo content of Fe – Mo alloy films from the alkaline bath plotted as a function of the Mo concentration in the bath. The Mo content of the film from this bath also increased with increasing Mo concentration. A graph for the Mo content in the film versus the bath temperature is shown in Figure 7.4.2.25.4. An increase in the bath temperature appears to lead to a reduction in the Mo content. Figure 7.4.2.25.5 illustrates how the Mo content in the film changes with current density. The Mo content decreased with increasing current density.


X-ray diffraction patterns were taken from Fe– Mo alloy films plated in the acid bath as seen in Figure 7.4.2.25.6. The bath pH was 3 and the current density was 400 A/m2. Although the presence of diffraction peaks from Cu substrates complicates

548

Nano-Plating

Figure 7.4.2.25.2. The Mo content of Fe–Mo alloy films plated from the acid bath, Bath (a), is plotted against the bath pH.

Figure 7.4.2.25.3. The Mo content of Fe –Mo alloy films obtained from the alkaline bath, Bath (b), is plotted as a function of the Mo concentration in the bath.


549

Figure 7.4.2.25.4. The Mo content in Fe–Mo alloy films plated from the alkaline bath, Bath (b), versus the bath temperature.

the overall patterns, diffraction peaks from the film clearly broadened with increasing Mo content in the film. The crystal size was reducing and the film became amorphous. Although the exact composition at which a crystalline-to-amorphous phase transition took place is not known, it is estimated to be about 30 at.% Mo. The structure (crystalline, amorphous, or mixed) of Fe –Mo alloy films plated from the baths having pH ¼ (A) 3 and (B) 5 are plotted in Figure 7.4.2.25.7 as a function of the current density and the Mo concentration in the bath. It is difficult to define a boundary

Figure 7.4.2.25.5. The Mo content of plated Fe–Mo alloy films obtained from the alkaline bath (cf. Table 7.4.2.25.1(b)) versus the current density.

550

Nano-Plating

Figure 7.4.2.25.6. X-ray diffraction patterns from various compositions of plated Fe– Mo alloy films.

Figure 7.4.2.25.7. The structure (crystalline, amorphous, or mixed) of Fe– Mo alloy films plated from the acid baths (cf. Table 7.4.2.25.1(a)) having pH ¼ (A) 3 and (B) 5.


551

Figure 7.4.2.25.8. X-ray diffraction patterns from Fe–Mo alloy films plated in the alkaline bath (cf. Table7.4.2.25. 1 (b)).

between the crystalline and amorphous phases from these plots. It is, however, clear that the film becomes amorphous when the Mo concentration in the bath increased. As shown in Figure 7.4.2.25.1, the Mo content in the film does not depend on the current density but depends instead on the Mo concentration in the bath. The increased Mo content in the film was responsible for the formation of the amorphous phase. The results shown in Figure 7.4.2.25.7 demonstrates that the film became amorphous at about 30 at.% Mo. Figure 7.4.2.25.7 illustrates the fact that the Mo content was a determining factor for producing the amorphous phase. All the films containing 30 , 55 at.% Mo were amorphous. X-ray diffraction patterns were also taken for

552

Nano-Plating

Figure 7.4.2.25.9. (a) The equilibrium phase diagram of Fe–Mo alloys and (b) the composition range of forming an amorphous phase in plated Fe–Mo alloy films.

Fe –Mo alloy films plated in the alkaline bath (see Figure 7.4.2.25.8). As the Mo content increased, the film also became fine-grained, finally leading to the formation of an amorphous phase. The composition range for forming an amorphous phase in plated Fe– Mo alloy films is seen in Figure 7.4.2.25.9(b) together with (a) the equilibrium phase diagram of Fe –Mo alloys. Figure 7.4.2.25.10 contains SEM micrographs showing the surface morphology and fracture surface of various compositions of plated Fe– Mo alloy films. It is interesting to note that the fracture surfaces of the amorphous films, i.e. (c) –(e), exhibit a fibrous structure along the film thickness direction. Such a fracture morphology was not expected for amorphous films, because they generally exhibit a smooth, featureless morphology similar to the surface of a fractured glass.

Electrolytic Fe – Mo Figure 7.4.2.25.10. SEM micrographs showing the surface morphology and fracture surface of various compositions of plated Fe–Mo alloy films obtained from the acid bath (cf. Table 7.4.2.25.1(a)).

553

554

Nano-Plating

REFERENCES Watanabe, T., Naoe, T., Mio, J. & Katsumata, A. (1989) J. Surf. Finish. Soc. Jpn, 40, 458. Zhu, L. & Watanabe, T. (1992) J. Jpn Inst. Metals, 56, 664.

7.4.2.26. Electrolytic Fe – Ni (Fukumuro, 1996, 1998) 7.4.2.26.1. PLATING METHOD

Various Fe –Ni plating baths were prepared by changing the mixing ratio of iron sulfate (FeSO4·7H2O) and nickel sulfate (NiSO4·7H2O) salts was while keeping the total metal ion concentration constant at 1 mol/l. These two metal salts were mixed together with a buffer, 0.5 mol/l boric acid, in deoxidized/distilled water. The bath pH was adjusted to 2.0 with dilute sulfuric acid. A saturated calomel electrode (SCE) was used as the reference electrode. Iron – nickel films were electroplated by a controlled potential electrolysis method. The constant potentials ranging from 2 0.8 to 2 1.3 V versus SCE were applied to plate the films. The bath temperature was 50 8C and the bath was agitated with a stirrer.


Figure 7.4.2.26.1 shows a graph for the Ni content of plated Fe –Ni alloy films plotted against the Ni concentration in the bath. It is interesting to note from the figure that the Ni

Figure 7.4.2.26.1. The Ni content of plated Fe–Ni alloy films plotted against the Ni concentration in the bath.

555

556

Nano-Plating

content in the film depended primarily on the Ni concentration in the bath but not on the deposition potential. Another characteristic of this plating system is that the formation of the 50 at.% Ni films requires the concentration of NiSO4 in the bath to be 10 times greater than FeSO4. Because of such an unbalanced mixture, this Fe –Ni alloy plating system is regarded as an anomalous co-deposition type.


7.4.2.26.3.1 Structure Figure 7.4.2.26.2 displays X-ray diffraction patterns for plated Fe– Ni alloy films with five different compositions. The diffraction pattern of the Fe-87.1 at.% Ni alloy film plated at 2 0.7 V versus SCE was compared with the JCPDS file and was found to be that of a fcc g solid solution. The Fe– 76.0 at.% Ni alloy film plated at 2 1.3 V versus SCE, had the composition of the intermetallic compound, FeNi3, which is listed in the equilibrium phase diagram (cf. Figure 7.4.2.26.3(a)), but did not exhibit the characteristic diffraction lines associated with the ordered structure. For the Fe – 76.0 at.% Ni alloy film, therefore, only the disordered g phase was formed instead of the ordered FeNi3 intermetallic compound. As the Ni content decreased sequentially from 58.6, to 48.0 and down to 38.3 at.%, the {111} peak of the g phase continuously shifted toward the lower-angle side, which corresponded to the expansion of the g phase lattice, and, at the same time, the intensity of the {200} peak became progressively weaker. For the Fe – 27.8 at.% Ni alloy, a new smooth peak appeared near the 528 diffraction angle, while the {200} peak almost vanished, and the overall pattern changed to a bcc structure. With a further decrease in the Ni content to 19.3 at.%, an a solid solution having a bcc structure emerged, followed by the appearance of the sharp {110} peak of the a phase when the Ni content reached 11.6 at.%. Figure 7.4.2.26.2 also supports the fact that the structure of plated Fe – Ni films can be solely determined by the alloy composition, independent of the deposition potential value used. Figure 7.4.2.26.3 shows (a) the equilibrium phase diagram of Fe –Ni alloys, (b) a phase diagram determined from plated Fe –Ni alloys, and (c) the lattice constants of the observed phases. The phase diagram of plated Fe – Ni alloy films can be established by the film composition alone, without need for considering the deposition potential. The bcc a phase was formed in the composition range of 0 –15 at.% Ni, whereas the fcc g phase was generated in the range of 35 –100 at.% Ni. The a/g mixed phases were formed between 15 and 35 at.% Ni. The lattice constant of the a phase tended to decrease slightly from the constant of pure Fe (0.2867 nm) with increasing Ni content (see Figure 7.4.2.26.3(c)). The lattice constant of the g phase increased linearly from pure Ni (0.3519 nm) up to 55 at.% Ni with increasing Fe content. The lattice constant of the g phase apparently obeys

Electrolytic Fe –Ni

557

Figure 7.4.2.26.2. X-ray diffraction patterns for plated Fe–Ni alloy films with eight different compositions.

558

Nano-Plating

Figure 7.4.2.26.3. (a) The equilibrium phase diagram of Fe– Ni alloys; (b) and (c) are the phase diagram of plated Fe–Ni alloys and the lattice constants of the observed phases, respectively.


559

Vegards law because it follows the chained line, which was drawn by connecting the lattice constant of pure fcc Ni to the lattice constant (0.36394 nm) of fcc g-Fe. It is important to note that in order to be consistent with the same structural type (fcc) as Ni, the lattice constant of the fcc g-Fe at the allotropic-transformation temperature (912 8C) (Basinski et al., 1955) was adopted instead of the bcc a-Fe. The lattice constant of the fcc g-Fe took a maximum value of 0.3575 nm at the alloy composition of about 38 at.% Ni, but decreased with a further increase in the Fe content, accompanied by the formation of the a phase. Although the lattice constant change in the a and g phases with the alloy composition differed slightly from those reported previously (Brenner, 1963; Jartych et al., 1992; Grimmett et al., 1993), it generally agreed (Owen and Yates, 1937) with the metallurgically processed Fe – Ni alloys. The lattice constant of plated films could, of course, be affected by the presence of inclusions, such as hydrogen and organic/inorganic additives, or internal stresses. If plated films are free of these effects, they should form a structure similar to metallurgically processed alloys. 7.4.2.26.3.2 Crystal size and surface morphology Figure 7.4.2.26.4 shows a relationship between the grain size of the a/g phases in plated Fe –Ni alloy films and their Ni content. The grain size (or crystal size) was directly measured from TEM micrographs. An X-ray diffraction technique was also used to determine the grain size from the full width at half maximum (FWHM) of a diffraction

Figure 7.4.2.26.4. Relationship between the grain size of the a/g phases in plated Fe–Ni alloy films and their Ni content.

560

Nano-Plating

peak using Scherrer’s equation. As noted in Figure 7.4.2.26.4, both X-ray and TEM measurements gave comparable values for the grain size of the a and g phases. Other investigators also reported (Jartych et al., 1992; Grimmett et al., 1993; Cheung et al., 1995) a similar grain-size dependence on the composition but did not clarify how such a dependence originated. Comparing Figure 7.4.2.26.3 with Figure 7.4.2.26.4, we find that the grain size of the a phase decreased with increasing Ni content from the 30 nm of pure Fe, to less than 10 nm at a composition of 28 at.% Ni, and finally diminished to zero. The grain size of the g phase was also about 10 nm at 28 at.% Ni but increased with increasing

Figure 7.4.2.26.5. SEM micrographs revealing the surface morphology and fracture surface of plated Fe –Ni alloy films containing various amounts of Ni, i.e. (A) 82, (B) 60, (C) 48, (D) 28, (E) 19, and (F) 11 at.%. The corresponding cross-section TEM micrographs and electron diffraction patterns are shown in (b) and (c).


561

Ni content up to 40 nm, which is the grain size of pure Ni. The alloy composition of 15– 35 at.% Ni is considered to be a transitional range, where the structure of plated Fe– Ni alloy films gradually changes from the a phase to the g phase. From the above results, it is possible to envisage the following mechanism. The relative volume ratio of deposited Fe and Ni atoms in plated Fe – Ni alloy films determines the ratio of the a and g phases, which in turn establish their grain (crystal) size distribution. Figure 7.4.2.26.5(a) displays SEM micrographs revealing the surface morphology and fracture surface of plated Fe – Ni alloy films having various amounts of Ni, i.e. (A) 82,


562

Nano-Plating

(B) 60, (C) 48, (D) 28, (E) 19, and (F) 11 at.%. The corresponding cross-section TEM micrographs and electron diffraction patterns are shown in Figure 7.4.2.26.5(b) and (c), respectively. As will be discussed later, the deposition potential is known (Fukumuro et al., 1996) to affect the surface morphology markedly. All the films were grown at the deposition potential of 21:1 , 1:3 V versus SCE, which yielded a relatively smooth surface as seen in Figure 7.4.2.26.5(a). The TEM cross-sections of all the films appear to exhibit a uniform structure along the film thickness direction. According to Figures 7.4.2.26.3 and 7.4.2.26.4,



563

Figure 7.4.2.26.6. Cross-section TEM micrographs showing the structure of plated Fe–Ni alloy films, which were obtained at different deposition potentials.

564

Nano-Plating

the Fe – 82 at.% Ni film consists of only the g phase with a grain size of 30 nm. These crosssection TEM micrographs indicate that the grain size of the g phase decreased continuously as the Ni content decreased sequentially from (B) 60, to (C) 48, and to (D) 28 at.%. The Fe – 28 at.% Ni film containing a mixture of the a and g phases is very fine-grained (, 10 nm). For the Fe– 19 at.% Ni film, the grain size of the a phase increased, while the g phase almost vanished. Finally, only the a phase with a grain size of 15 nm remained in the Fe –11 at.% Ni film, which had a columnar structure as seen in Figure 7.4.2.26.5(F) (b). Figure 7.4.2.26.6 illustrates the cross-sectional view of Fe– Ni alloy films plated at different deposition potentials. The films shown in Figure 7.4.2.26.6(a) and (b) were obtained at 2 1.3 and 2 0.7 V versus SCE from a bath having the metal salt ratio of NiSO4 =FeSO4 ¼ 0:98=0:02: Both the films had the same lattice constant of the g phase, the same Ni content (80 at.%), and the same grain size (, 30 nm). The surface of the 2 1.3 V film was smooth but the surface of the 2 0.7 V was rough. Similarly, the films shown in Figure 7.4.2.26.6(c) and (d) were obtained at 2 1.3 and 2 0.8 V versus SCE from a bath having the metal salt ratio of NiSO4 =FeSO4 ¼ 0:90=0:10: These two depositions yielded two films with the same lattice constant of the g phase, the same Ni content (48 at.%), and the same grain size (, 20 nm). For the surface morphology, which can be estimated from the cross-sections in Figure 7.4.2.26.6, the film obtained at the less noble potential (2 1.3 V versus SCE) showed a smoother surface. The film obtained with the more noble potential (2 1.0 V), displayed a rough surface although the grain size decreased. To establish the effect of the deposition potential, the current density, and the deposition rate on the surface morphology of plated Fe– Ni alloy films, we conducted additional TEM studies on the surface morphology and cross-section structure. The following trend was observed. For the films plated under the plating conditions of . 2 1.0 V versus SCE, , 500 A/m, and , 15 nm/s, the surface was very rough regardless of the grain size. Conversely, for the films plated under the plating conditions of , 2 1.0 V versus SCE, . 500 A/m, and . 15 nm/s, the surface was smooth.

REFERENCES Fukumuro, N., Chikazawa, M. & Watanabe, T. (1996) J. Surf. Finish. Soc. Jpn, 47, 461. Fukumuro, N., Imai, M., Chikazawa, M. & Watanabe, T. (1998) J. Magn. Soc. Jpn, 22, 1268.

FURTHER READING Aotani, K. (1950) J. Jpn Inst. Metals, B-14, 55. Aotani, K. (1952) Denki Kagaku, 20, 1. Basinski, Z.S., Hume-Rothery, W. & Sutton, A.L. (1955) Proc. R. Soc., A229, 459. Brenner, A. (1963) Electrodeposition of Alloys, vol. II, Academic Press, New York, p 301.

Electrolytic Fe –Ni Cheung, C., Djuanda, F., Erb, U. & Palumbo, G. (1995) Nanostruct. Mater., 5, 5. Dahms, H. & Croll, I.M. (1965) J. Electrochem. Soc., 112, 771. Grande, W.C. & Talbot, J.B. (1993) J. Electrochem. Soc., 140, 669. Grimmett, D.L., Schwartz, M. & Nobe, K. (1993) J. Electrochem. Soc., 140, 973. Hessami, S. & Tobias, C.W. (1989) J. Electrochem. Soc., 136, 3611. Jartych, E., Budzynski, M. & Zurawicz, J.K. (1992) Hyperfine Interact., 73, 255. Kamo, Y. (1987) J. Magn. Soc. Jpn, 11, 291. Matlosz, M. (1993) J. Electrochem. Soc., 140, 2272. Ohno, I. & Haruyama, S. (1991) J. Jpn Inst. Metals, 30, 735. Omata, Y., Asai, H. & Shinozaki, T. (1994) J. Magn. Soc. Jpn, 18, 285. Owen, E.A. & Yates, E.L. (1937) Proc. Phys. Soc., 49, 17. Owen, E.A. & Yates, E.L. (1937) Proc. Phys. Soc., 49, 178. Owen, E.A. & Yates, E.L. (1937) Proc. Phys. Soc., 49, 307. Popov, B.N., Yin, K.-M. & White, R.E. (1993) J. Electrochem. Soc., 140, 1321. Romankiw, L. (1970) IEEE Trans. Magn., MAG-6, 597. Takai, M. & Osaka, T. (1997) J. Magn. Soc. Jpn, 21, 51. Venkatasetty, H.V. (1970) J. Electrochem. Soc., 117, 403. Winand, R. (1994) Electrochim. Acta, 39, 1091. Yasuda, S. & Koura, N. (1982) J. Metal Finish. Soc. Jpn, 33, 427.

565

7.4.2.27. Electrolytic Fe – W (Wang, 2000) 7.4.2.27.1. PLATING METHOD

The compositions of the Fe –W alloy plating bath used in this study, and deposition conditions, are listed in Table 7.4.2.27.1. The mixing ratio of FeSO4·7H2O and Na2WO4·2H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.26 mol/l. Different film compositions could also be obtained by changing the current density. The 80 8C bath temperature was found to yield an excellent film. The plating, however, could be performed at 35 8C, which was used in the present experiment. The bath pH was adjusted to 8 with dilute sulfuric acid or ammonium water. Each plating was terminated when 200 C was consumed. These plating conditions produced films about 6 mm thick.


In Figure 7.4.2.27.1, the W content of Fe– W films plated at (a) 500 and (b) 1000 A/m2 and the current efficiency are plotted against the W concentration in the bath. For the two current densities used, the W content in the film increased slowly with the W concentration in the bath. Figure 7.4.2.27.2 is the W content of Fe – W alloy films plated in a 40 at.% W bath plotted against the current density. The W content increased with increasing current density, while the current efficiency decreased. Figure 7.4.2.27.3 shows the effect of the bath pH on the W content of various Fe –W films, which were plated at 500 A/m2. The Fe-to-W mixing ratio was also changed. The maximum W content was obtained at pH ¼ 4: Based on this experiment, we found that Fe –W alloy films containing 11.5– 31 at.% W can be obtained from the present plating bath. Table 7.4.2.27.1. The composition of Fe –W alloy plating baths and the plating condition. Composition

Plating condition

FeSO4·7H2O þ Na2WO4·2H2O 0.26 mol/l (NH4)2C4H4O6 0.26 mol/l

Bath temperature: 35 8C Current density: 100– 1200 A/m2 Consumed electricity: 200 C Agitation: yes Substrate: polycrystalline Cu Anode: Pt plate

566

Electrolytic Fe – W Figure 7.4.2.27.1. The W content of Fe–W films plated at (a) 500 and (b) 1000 A/m2 and the current efficiency are plotted against the W concentration in the bath.

567

568

Nano-Plating

Figure 7.4.2.27.2. The W content of Fe– W alloy films plated in a 40 at.% W bath plotted against the current density.


Figure 7.4.2.27.4 contains X-ray diffraction patterns from various compositions of plated Fe –W alloy films. Sharp diffraction peaks from a-Fe crystals appeared for the 11.6 and 16.9 at.% W alloy films. These peaks broadened considerably for the films containing

Figure 7.4.2.27.3. The effect of the bath pH on the W content of various Fe– W films plated at 500 A/m2.

Electrolytic Fe – W

569

Figure 7.4.2.27.4. X-ray diffraction patterns from various compositions of plated Fe–W alloy films.

more than 17.5 at.% W. The peak broadening persisted up to the composition of 31 at.% W. This continuous peak broadening demonstrates that plated Fe– W alloy films were initially crystalline for the low W content and their grain size was steadily refined with increasing W content. The films finally transformed to an amorphous phase. It is not clear at the moment at which composition the Fe –W film became amorphous. From the X-ray diffraction study, we identified the presence of bcc a-Fe crystals in the low W-content films. The peak location was shifted toward the lower-angle side, which is equivalent to the expansion of the bcc lattice. A change in the lattice constant with the W content in the film is plotted in Figure 7.4.2.27.5. As seen in the plot, the lattice constant increased linearly with increasing W content and the extrapolated value coincided with the lattice constant of pure bcc W. This trend obeys

570

Nano-Plating

Figure 7.4.2.27.5. A change in the lattice constant of the a-Fe of plated Fe –W alloy films as a function of the W content in the films.

Vegards law, proving that W atoms are alloyed with Fe substitutionally. According to the equilibrium phase diagram of Fe –W alloys shown in Figure 7.4.2.27.6(a), Fe metal can dissolve up to 14.3 at.% W at high temperatures but no W solubility is present at ambient temperatures. Consequently, these Fe –W solid-solution alloy films obtained by a plating method are in a state of supersaturation and thus a thermodynamically meta-stable phase. It was shown that the structure of plated Fe –W alloy films changed from crystalline to amorphous phases with increasing W content. To determine the transitional composition for the crystalline-to-amorphous phase, we conducted a heating experiment to understand their thermal behavior. We chose three alloy films; (1) crystalline film with the 16.9 at.% W, (2) crystalline or amorphous films with 17.5 at.% W, and (3) completely amorphous deposit (31.0 at.% W). Figure 7.4.2.27.7 shows X-ray diffraction patterns for the 16.9 at.% W film, which was heat-treated at various temperatures. The as-deposited film exhibited a {110} diffraction peak of a supersaturated solid-solution alloy at an angle of 44.68. The intensity of the {110} peak increased with increasing heat-treatment temperature and the peak shifted suddenly at 700 8C toward the high-angle side, coinciding with the {110} peak of the a-Fe phase. This change is due to the out-diffusion of W atoms into the grain boundaries of the supersaturated Fe– W solid-solution alloy. A further increase in the


571

Figure 7.4.2.27.6. (a) The equilibrium phase diagram of Fe–W alloys and (b) the phase diagram of plated Fe–W alloy films.

heat-treatment temperature to 800 and 1000 8C led to the formation of Fe2W crystals, which are probably formed between W and Fe atoms segregated at the grain boundaries. The final structure consists of a-Fe and Fe2W phases, which is in agreement with the equilibrium phase diagram (see Figure 7.4.2.27.6(a)). From the above result, we can conclude that the as-deposited 16.9 at.% W alloy film initially contained fine a-Fe crystals (supersaturated with W), which transformed to more stable a-Fe crystals and decomposed into the a-Fe and Fe2W phases.

572

Nano-Plating

Figure 7.4.2.27.7. X-ray diffraction patterns showing a structural change in a plated Fe– 16.9 at.% W film upon heat treatment at various temperatures.

The effect of heating on the structure of the 17.5 at.% W alloy film is shown in Figure 7.4.2.27.8. The diffraction pattern did not change up to 400 8C, but showed the formation of the a-Fe suddenly at 600 8C. At 700 8C, new peaks marked with a symbol, M.S., appeared and the corresponding crystal to these peaks has never been reported previously. Since these peaks disappeared at higher temperatures, this phase is believed to be a metastable phase having an unknown composition of FexWy. The disappearance of the peaks was accompanied by the appearance of the peaks from the Fe2W phase. The final structure


573

Figure 7.4.2.27.8. X-ray diffraction patterns showing a structural change in a plated Fe–17.5 at.% W film upon heat treatment at various temperatures.

was made up of the a-Fe and Fe2W phases, consistent with the equilibrium phase diagram (cf. Figure 7.4.2.27.6(a)). Heating of the 31.0 at.% W alloy film yielded X-ray diffraction patterns (cf. Figure 7.4.2.27.9) similar to those seen for the 17.5 at.% W alloy film (see Figure 7.4.2.27.8). The amount of a-Fe formed at 600 8C was extremely small. At 1000 8C, the a-Fe became almost non-existent but the intensity of the Fe2W peaks became stronger. Since the amount of the a-Fe and Fe2W phases is proportional to their peak intensity, it is in agreement with

574

Nano-Plating

Figure 7.4.2.27.9. X-ray diffraction patterns showing a structural change in a plated Fe– 31.0 at.% W film upon heat treatment at various temperatures.

the equilibrium phase diagram (see Figure 7.4.2.27.6(a)). From the present heating experiment, the following conclusion can be made. As-plated Fe –W films having a W content of 17.5– 31 at.% were amorphous. At 600 8C, a-Fe and meta-stable phases were nucleated. The meta-stable phase and a-Fe crystals reacted to form Fe2W crystals. Based on the heating experiment, we consider that there is a amorphous/crystalline boundary between 16.9 and 17.2 at.% W for the plated Fe– W alloy films. This result is illustrated in Figure 7.4.2.27.6(b). It is important to remember that the present experiment


575

Figure 7.4.2.27.10. SEM micrographs showing the surface morphology and fracture surface of plated Fe–W films containing various W content.

576

Nano-Plating

was conducted up to an alloy composition of 31 at.% W and thus does not cover the whole composition range. Figure 7.4.2.27.10 displays SEM micrographs showing the surface morphology and fracture surface of plated Fe –W films with varying W content. Crystalline films display the smooth surface for both surface morphology and fracture surface, whereas the amorphous deposits exhibits a smooth surface and the fracture surface is like that of glass, indicating structural uniformity.

REFERENCE Wang, F., Ito, K. & Watanabe, T. (2000) J. Jpn Inst. Metals, 64, 1335.

FURTHER READING Barnes, C. (1985) Trans. Inst. Metal Finish., 2, 47. Fukushima, H., Akiyama, T., Akagi, S. & Azuma, T. (1978) J. Jpn Inst. Metals, 42, 980. Omi, T. & Yamamoto, H. (1973a) J. Metal Finish. Soc. Jpn, 24, 428. Omi, T. & Yamamoto, H. (1973b) J. Metal Finish. Soc. Jpn, 24, 612.

7.4.2.28.

Electrolytic Fe – Zn (Hu and Watanabe, 2000)


An Fe – Zn alloy plating was prepared by dissolving two metal salts, FeCl2·4H2O and ZnCl2, in pure water, with 20 g/l NaCl as a supporting electrolyte. The mixing ratio of FeCl2·4H2O and ZnCl2 was changed to obtain various alloy compositions at different molar ratios. The bath pH was adjusted to 1.5 with hydrochloric acid. The plating was carried out at a current density of 750 A/m2 in the 24 8C bath until 1000 kC/m2 was reached. In order to avoid the effect of Fe oxidation in the bath, the plating was generally done immediately after the bath was prepared.


Figure 7.4.2.28.1 shows the Zn content of plated Fe– Zn alloy films plotted against the Zn concentration in the bath. The Zn content of the film increased with increasing Zn

Figure 7.4.2.28.1. Relationship between the Zn content of plated Fe–Zn alloy films and the Zn concentration in the bath.

577

578

Nano-Plating

Figure 7.4.2.28.2. Relationship between the current efficiency and the Zn concentration in the Fe –Zn alloy plating bath.

concentration in the bath. When the Zn concentration in the bath was increased to 70 at.%, the Zn content of the film reached 99 at.%. The Zn content of all the plated Fe –Zn films was higher than the Zn concentration of the bath. This trend is also indicated by the upward concave shaped curve. Although Zn metal is less noble than Fe metal, it deposits preferentially in the Fe –Zn alloy plating system, which is thus designated as an anomalous codeposition type. Figure 7.4.2.28.2 plots the current efficiency as a function of the Zn concentration in the bath. The current efficiency was about 95% over the entire Zn concentration. From the current density and the current efficiency, the deposition rate can be estimated to be about 0.10 nm/s.


7.4.2.28.3.1 Structure An X-ray diffraction analysis was conducted on a wide compositional range of Fe– Zn alloy films, as seen in Figure 7.4.2.28.3. In the films containing 0 , 25 at.% Zn, only the bcc a phase (a solid-solution alloy of Fe) was present, but in the films having more than

Electrolytic Fe– Zn

Figure 7.4.2.28.3. X-ray diffraction patterns from various compositions of plated Fe–Zn alloy films.

579

580

Nano-Plating

30.5 at.% Zn, the cubic G phase with the stoichiometry of Fe4Zn9 appeared. When the Zn content reached 38.7 at.%, a small amount of cubic G1 phase (FeZn4) started to form. At a Zn content of 71.9 at.%, the a phase disappeared completely, and when the Zn content increased beyond 74.8 at.%, the hcp h phase started emerging. Beyond 94.9 at.% Zn, only the h phase was present. Based on the above X-ray results, we constructed a phase diagram for plated Fe– Zn alloy films in Figure 7.4.2.28.4(b) and (c). For comparative purposes, the equilibrium phase diagram of Fe– Zn binary alloys is illustrated in Figure 7.4.2.28.4(a). It is interesting to note that the d and z phases present in the equilibrium phase diagram were not observed in the plated films and that only the a, G, G1, and h phases were found. Adaniya et al. plated Fe –Zn alloy films at a current density of 5000 A/m2 in a 40 8C sulfate-type bath ðpH ¼ 3:0Þ and obtained films containing the a, G, d, and h phases. Kondo (1988, 1989) used a sulfate-type bath to plate Fe – Zn films and found a, G (bcc), G1 (fcc), and h (hcp) phases similar to those identified in this study. 7.4.2.28.3.2 Texture and lattice constant Based on the X-ray results in Figure 7.4.2.28.3, the lattice constants and texture of the a phase were plotted in Figure 7.4.2.28.5(a) and (b) as a function of the Zn content in the film. The texture index was calculated using Willson’s equation (see Section 6.5.2). In Figure 7.4.2.28.3(a), the (110) texture index is seen to be about 1.5 for all compositions. Thus, the texture for these Fe –Zn alloy films is 110, although it is not a strongly preferred orientation. The lattice constant of the a phase, which is a solid-solution alloy of Fe, was calculated from the strongest peak, i.e. 110. The lattice constant of a plated pure Fe film, containing no Zn, agreed well with the lattice constant (0.2866 nm—a dotted horizontal line) of pure Fe. The lattice constant of the a phase increased with increasing Zn content, accompanied by dissolution of Zn atoms in Fe metal. According to the equilibrium phase diagram, Zn cannot dissolve in Fe at ambient temperatures. The a phase is thus a solid solution of Fe supersaturated with Zn, which is in a meta-stable state. In Figure 7.4.2.28.6(a) and (b), we plotted the lattice constants of the G (Fe4Zn9) and G1 (FeZn4) phases as a function of the Zn content in their phases. To calculate the lattice constant of the G phase, we used the {321} peak located at 37.48. The published lattice constant (0.895 nm) of the G phase from the JCPDS file is drawn as a horizontal dotted line in Figure 7.4.2.28.6(a). It is clear that the lattice constant of the G phase generally increases with increasing Zn content in the film and intersects with the dotted line at a composition of , 75 at.% Zn, which is not far from the stoichiometric value of 69.2 at.% Zn for the intermetallic compound, Fe4Zn9. To calculate the lattice constant of the G1 phase, we used the (733) peak located at 41.18, the (880) 58.08, and (1171) 68.08. The published lattice constant value (1.798 nm) of the G1


581

Figure 7.4.2.28.4. (a) The equilibrium phase diagram of Fe–Zn alloys is listed in (a), whereas the phase diagrams of plated Fe –Zn alloy films are shown in (b) and (c).

582 Nano-Plating Figure 7.4.2.28.5. A change in the (a) lattice constant and (b) texture of the a-Fe phase in plated Fe–Zn alloy films is plotted as a function of the Zn content in the films.


Figure 7.4.2.28.6. A change in the lattice constants of the (a) G and (b) G1 phases in plated Fe–Zn films with increasing Zn content.

583

584

Nano-Plating

phase in the JCPDS file is drawn as a horizontal dotted line in Figure 7.4.2.28.6(b). The lattice constants determined from the (733) and (1171) peaks appear to be in good agreement with the published value and increased slowly with increasing Zn content. The lattice constant values obtained from the (880) peak are slightly larger. It is not clear at moment why the (880) peak gave higher values. Figure 7.4.2.28.7(a)– (d) displays, respectively, the lattice constant, a; lattice constant, c; the c=a ratio, and the texture index of the h phase plotted as a function of the Zn content of the film. The lattice constant, a, is shifted by , 0.0015 nm (0.56%) upward from the standard value (0.2665 nm) of Zn listed in the JCPDS. The lattice constant, c,

Figure 7.4.2.28.7. A change in (a) the lattice constant,a, (b) lattice constant, c, (c) the c/a ratio, and (d) the texture index of the h phase in plated Fe –Zn alloy films versus a change in the Zn content in the films.


585

changed in the form of an upward concave curve around the standard c value (0.4947 nm) of Zn (drawn as a horizontal dotted line). From the above results, we also plotted a change in the c=a value with Zn content together with the c=a value for ideal packing (drawn as a horizontal dotted line) (see Figure 7.4.2.28.7(c)). Here, the ideal packing value of 1.856 for the c=a ratio can be obtained if each atom is assumed to be spherical. All the data points lie below the ideal value, indicating that the h phase is a hexagonal close-packed crystal, whose c axis is more compressed than the axis of the Zn crystal. Zinc metal cannot dissolve Fe atoms at ambient temperatures; therefore there is a possibility that the Zn crystal is supersaturated with Fe atoms. According to Figure 7.4.2.28.7(a), however, the lattice constant, a, of the plated pure Zn film is already expanded by as much as 0.0007 nm (0.26%) without alloying. In addition, the upward concave curves in the lattice constants, a and c (cf. Figure 7.4.2.28.7(a) and (b)), cannot be explained by the alloying phenomenon. An alternate explanation of the lattice constant change could be due to hydrogen inclusion. The texture index value of the h phase appears to display a scatter but the 102 or 103 texture can be observed at the Zn-rich side. 7.4.2.28.3.3 TEM observations To further confirm the X-ray results, we conducted TEM studies on the structure of plated Fe –Zn films. Figure 7.4.2.28.8 shows plan-view bright-field TEM micrographs and their associated electron diffraction patterns for five different Fe– Zn alloy films. Figure 7.4.2.28.8(a) is a TEM micrograph and an electron diffraction pattern for the 35.3 at.% Zn alloy film, which contained the a phase and a small amount of G phase according to the X-ray analysis. The electron diffraction pattern revealed the presence of the bcc a phase, but did not show the G phase. Although the G phase was not detected by TEM, it could have been distributed too sparsely, for the TEM to detect. The 35.3 at.% Zn alloy film is therefore most likely to contain both the a and G phases. The crystal (or grain) size of the a phase can be determined from the full width at the half maximum (FWHM) of an X-ray diffraction peak using Scherrer’s equation. Since the {111} peak of the Cu substrate overlapped with a peak from the a phase, we could not determine the crystal size precisely by the X-ray method. Consequently, the crystal size of the a phase was determined by TEM, which yielded 15 , 30 nm: Figure 7.4.2.28.9(b) shows the structure of the 43.9 at.% Zn alloy film. The associated electron diffraction pattern indicates the presence of the bcc G phase. Both the G and G1 phases appeared in the 61.2 at.% Zn film (see Figure 7.4.2.28.8(c)). An increase in the Zn content to 74.8 at.% led to the formation of three phases, G, G1, and h (see Figure 7.4.2.28.8(d)). A further increase to 94.9 at.% Zn resulted in the production of the same three phases, G, G1, and h as shown in Figure 7.4.2.28.9(a). The lattice matching relationship among the three phases is shown schematically in Figure 7.4.2.28.9(b). The following relationship was established.

586

Nano-Plating

Figure 7.4.2.28.8. Plan-view bright-field TEM micrographs and their associated electron diffraction patterns for five different electrolytic Fe–Zn alloy films.

(121) G//(111) G1//(0111) h , 121 . G//, 011 . G1//, 7253 . h Therefore, these three phases (G, G1, and h) are formed by matching crystallographically. The pure Zn film in Figure 7.4.2.28.8(e) had the expected hexagonal structure.


587

Figure 7.4.2.28.9. (a) An electron diffraction pattern from a plated Fe-94.9 at.% Zn film containing the h, G, and G1 phases, and (b) its index pattern.

The crystal size of the a phase was 15 , 30 nm; whereas the G and G1 phases were 50 , 80 nm: The G1 phase disappeared when the Zn content in the film became high. Finally, the h phase has a large crystal size with a hcp structure.

7.4.2.28.3.4

Surface morphology

Figure 7.4.2.28.10(a) – (f) comprises SEM micrographs showing the surface morphologies of various compositions of plated Fe– Zn alloy films. Phases in these films are indicated in the figure and the surface roughness and cracks can be seen. The surface of the pure Fe is relatively smooth but contains cracks (cf. Figure 7.4.2.28.10(a)). Numerous microcracks were generated in the Fe films plated at a current density of more than 500 A/m2 from the room-temperature bath. Cracks were observed in the films containing up to 43.9 at.% Zn. Alloying of Fe with Zn was accompanied by the development of large surface irregularities in the form of nodules. These surface nodules consist of many fine grains as noted in TEM micrographs in Figure 7.4.2.28.8. For the 74.8 at.% Zn film, small (, 1 mm) nodules are seen to form larger ð3 , 7 mmÞ ones, which are actually aggregates of 50 , 80 nm grains. Finally, when the Zn content reached 94.9 at.%, the film exhibited a characteristic morphology of hexagonal plates.

588

Nano-Plating

Figure 7.4.2.28.10. SEM micrographs showing the surface morphologies of various compositions of plated Fe –Zn alloy films. The composition and phases are indicated in each micrograph.

REFERENCES Hu, X.-T. & Watanabe, T. (2000) J. Jpn Inst. Metals, 64, 234. Kondo, K. (1988) Tetsu to-Hagane, 74, 2300. Kondo, K. (1989) ISIJ Inc., 29, 517. FURTHER READING Adaniya, T., Hara, T., Sigiyama, M., Homma, T. & Watanabe, T. (1985) Plat. Surf. Finish., 72(August), 52. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 2, ASM International, p. 1795– 1797.

7.4.2.29. Electrolytic In– Sn (Watanabe et al., 1999)


The compositions of the In –Sn alloy plating bath used in this study are listed in Table 7.4.2.29.1. In these baths, several reagents were added for the following purposes: (1) complexing agents (citric acid and hydrazine chloride) were added to stabilize Sn2þ ions and (2) addition agents (gelatin and dextrin) were put into the bath to obtain a uniform film by preventing the formation of Sn whiskers. Prior to the use of these reagents, we made sure that no atomic elements other than In and Sn were incorporated into the deposits. The bath pH was adjusted to 2.5 using ammonium chloride as a buffer. Different compositions of In –Sn alloy films were obtained by changing the bath chemistry and the current density.


The In content of plated In – Sn alloy films obtained from seven baths (cf. Table 7.4.2.29.2) is plotted against the current density in Figure 7.4.2.29.1. In this graph, the crystal structure of the deposit is also indicated with different symbols, which will be discussed later. For all the baths, the In content in the film increased with increasing current density and increasing In concentration in the bath.

Table 7.4.2.29.1. The composition of an In–Sn alloy plating bath. Composition

Concentration

InCl3·4H2O SnCl2·2H2O C2H8O7 N2H4·2H2O NH4Cl Gelatin Dextrin

X mol/l (0.15 2 X) mol/l 0.5 mol/l 0.7 mol/l 1.5 mol/l 2 g/l 10 g/l

589

590

Nano-Plating

Figure 7.4.2.29.1. The In content of plated In-Sn alloy films obtained from 7 baths is plotted against the current density (see text for details).

Table 7.4.2.29.2. Indium–tin alloy plating baths containing various mixing ratios of In and Sn metal salts. Metal salts (mol/l)

InCl3·4H2O SnCl2·2H2O

Bath name A

B

C

D

E

F

G

0.138 0.012

0.136 0.014

0.128 0.022

0.119 0.031

0.112 0.038

0.099 0.051

0.059 0.091

Electrolytic In– Sn

591


7.4.2.29.3.1 Structure X-ray diffraction patterns from plated In – Sn films with different alloy compositions are displayed in Figure 7.4.2.29.2. In these plated films, all the phases listed in the equilibrium phase diagram (cf. Figure 7.4.2.29.3(a)) (Massalski, 1990) appeared. These phases include

Figure 7.4.2.29.2. X-ray diffraction patterns from plated In–Sn films with different alloy compositions.

592

Nano-Plating

Figure 7.4.2.29.3. (a) The equilibrium phase diagram of In-Sn alloys and (b) the phase diagram of plated In– Sn alloy films.

pure Sn, tin solid solution, b phase, g phase, In, and indium solid solution. No meta-stable phases were found in these alloy deposits. Dotted horizontal lines in Figure 7.4.2.29.1 represent constant alloy compositions. The same phase or phases appeared between any two neighboring lines, regardless of the type of baths used. Although the current density changed the alloy composition, it did not determine the structure of the plated In – Sn films. Since the current density changed the alloy composition, we might erroneously assume that the current density is a controlling


593

factor in determining the resulting film structure. Like the construction of the equilibrium phase diagram, the phases are determined only by the composition and temperature. When the temperature was fixed at an ambient temperature, the only remaining factor is the alloy composition. Therefore, the current density has nothing to do with the structure of plated alloy films. Similarly, we can conclude that other electrochemical parameters like the overpotential or the deposition rate do not change the structure of plated films. The structure is uniquely determined by the alloy composition. In Figure 7.4.2.29.3(b), we summarized the phases observed in plated In –Sn alloy films over the whole composition range. Although a determination of the exact phase boundaries requires more experiments, the phases (cf. Figure 7.4.2.29.3(b)) observed in the plated films are in good agreement with those listed in the equilibrium phase diagram (cf. Figure 7.4.2.29.3(a)). From the locations of diffraction peaks shown in Figure 7.4.2.29.2, we studied a change in the diffraction planes of the observed phases as a function of the alloy composition. Indium metal has a tetragonal structure, and according to the JCPDS file (JCPDS File #5-0642), its lattice constants are a ¼ b ¼ 0:32517 nm and c ¼ 0:49459 nm: The lattice constants of the indium solid solution increased by about 1% after alloying with Sn. Tin metal, which appeared as white tin (b-Sn), has a tetragonal structure, and its lattice constants (JCPDS File #4-0673) are a ¼ b ¼ 0:5821 nm and c ¼ 0:3182 nm: The lattice constants of the tin solid solution decreased by 1% after alloying with In. Although these lattice-constant changes are very small, they demonstrate the formation of a solid solution alloy in both In and Sn. The b phase known as In3Sn also has a tetragonal structure with lattice constants a ¼ b ¼ 0:489 nm and c ¼ 0:445 nm (JCPDS File #7-0345). According to the equilibrium phase diagram in Figure 7.4.2.29.3(a), the b phase has a wide solubility range, suggesting that the lattice constants may change to accommodate the compositional change. Figure 7.4.2.29.4 illustrates how the interplanar spacing of six planes (110, 200, 210, 220, 300, and 311) in the b phase changes with the composition. Horizontal dotted lines represent the values of the lattice constants quoted in the JCPDS file (JCPDS File #7-0345). The interplanar spacing is seen to increase, decrease, or remain unchanged with the composition, depending upon the type of plane. No systematic trend can be discerned from this graph. For example, the 111 spacing increased by 0.2% regardless of the composition change, whereas the (220) plane spacing expanded by only 0.4%. The (220) interplanar spacing decreased with increasing In content but the (300) spacing increased. A sound understanding of this behavior is impossible. The dependence of the lattice constants on the composition is not well understood. The g phase designated as InSn4 also has a tetragonal structure. According to the JCPDS File #7-0396, its lattice constants are a ¼ b ¼ 0:3211 nm and c ¼ 0:2992 nm: The g phase has a wide solubility range and thus the lattice constants are expected to vary with alloy composition. Contrary to expectation, only the (001) lattice spacing expanded

594

Nano-Plating

Figure 7.4.2.29.4. A lattice spacing change in the 6 planes (110, 200, 210, 220, 300, and 311) of the b phase in plated In–Sn alloy films.

by 1% and no other planes exhibited any changes within the accuracy of the measurement. Differential scanning calorimetry (DSC) experiments were conducted on plated In – Sn alloy films and the results are shown in Figure 7.4.2.29.5. It is important to remember that in DSC curves, an exotherm is represented as an upward deviation from the baseline, whereas an endotherm appears as a downward deviation from the baseline. In Figure 7.4.2.29.5, an endothermic peak appeared at 103 8C. The appearance of this endothermic


595

Figure 7.4.2.29.5. Differential scanning calorimetry (DSC) curves for various compositions of plated In–Sn alloy films.

596

Nano-Plating

peak is due to the initial temperature imbalance associated with the DSC equipment. Each specimen appears to display one or two exothermic peaks. The starting and finishing temperatures of these peaks were superimposed on the equilibrium phase diagram as seen in Figure 7.4.2.29.6. In Figure 7.4.2.29.5, we note that a sharp exothermic peak appeared between 120 and 140 8C in In – Sn alloy films containing more than 51.9 at.% In. This peak corresponds to the melting temperature of the b phase. Plated In – Sn alloy films containing 7:1 , 18:2 at:% In showed a sharp exothermic peak at 180 , 230 8C, which corresponds to the melting temperature of the g phase. A broad exothermic peak extending from 135 to 195 8C in an In – Sn alloy film with 27 at.% In indicates a melting reaction in the g phase. The alloy films having the compositions of 18.2 and 27.9 at.% In show a small exothermic peak at 120 8C. Although there are no apparent reactions at this temperature according to Figure 7.4.2.29.6, this temperature corresponds to the eutectic reaction temperature

Figure 7.4.2.29.6. The starting and finishing temperatures of exothermic peaks were plotted on top of the equilibrium phase diagram.


597

Figure 7.4.2.29.7. SEM micrographs showing the surface morphologies of plated In–Sn films having various alloy compositions.

598

Nano-Plating

between the b and g phases. Apparently a small amount of b phase was present in the film and then transformed to the g phase, giving rise to the exothermic peak. This example confirms the fact that a solid formed by a plating method is not necessarily thermodynamically in equilibrium and thus could be in a meta-stable state. This phenomenon is similar to the formation mechanism of meta-stable solids obtained by quenching a metal from a liquid state. The alloy film with 85 at.% In exhibited a broad endothermic peak ranging from 165 to 190 8C. Since the film is in a liquid state at these temperatures, it is not clear what kind of reactions took place. As demonstrated in Figure 7.4.2.29.6, the final temperatures of all the exothermic reactions are slightly higher than the liquidus temperature of the equilibrium phase diagram. This was caused by the heating rate during the DSC experiments being too fast. A reduction in the heating rate should bring the final temperature down to the liquidus temperature. From the above results, the structure of plated In –Sn alloy films is considered to be in a supercooled state. The structure is thus very close to the equilibrium phase diagram. As noted in Table 7.4.2.29.1, several reagents were added to the bath, and it is possible that the films incorporated foreign elements from these reagents. These incorporated foreign elements might have modified the phase diagram shown in Figure 7.4.2.29.3(b). However, the use of an energy dispersive X-ray spectrometer attached to our SEM did not reveal any foreign elements except their main elements of In and Sn. 7.4.2.29.3.2 Surface morphology Figure 7.4.2.29.7 displays SEM micrographs showing the surface morphologies of plated In – Sn films having various alloy compositions. The composition and observed phases are denoted in the figure. The surface is generally rough for all the films. The surface roughness is high for the high Sn content films, or conversely, the surface is smooth for the high In content films. Particles lying on these surfaces were not identified here.

REFERENCES JCPDS File #5-0642. JCPDS File #4-0673. JCPDS File #7-0345. JCPDS File #7-0396. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 2, ASM International, p. 2295– 2296. Watanabe, T., Sato, I. & Chikazawa, M. (1999) J. Jpn Inst. Metals, 63, 483.

7.4.2.30. Electrolytic Ni –B (Onoda et al., 1990, 1992, 1998) 7.4.2.30.1. PLATING METHOD

The compositions of the Ni– B alloy plating bath, and the deposition conditions used in this study are listed in Table 7.4.2.30.1. The basic ingredients of a Ni– B plating bath were derived from a Watts-type bath (240 g/l NiSO4·6H2O and 45 g/l NiCl2·6H2O). In addition, 30 g/l (CH3)3NBH3 was added to the bath as a complexing agent and H3BO3 as a buffer. Trimethylamine borane (H3BO2) known as TMAB was also applied as a source of B. The bath pH was adjusted to 1:5 , 5:0 with dilute sulfuric acid or nickel carbonate. The bath temperature was 55 8C. A schematic view of the plating setup is illustrated in Figure 7.4.2.30.1, where a rotating disc electrode was used as a cathode. The plating was conducted in a still bath without rotating the disc electrode. The substrate, which was in the form of a drum, had an area of 20 £ 50 mm2 : A Ni ball was kept inside a titanium basket as an anode, which was enclosed in an anode bag. All the film thicknesses were 20 mm.


In Figure 7.4.2.30.2, the B content in plated Ni– B alloy films was plotted as a function of the current density. The B content appears to decrease with increasing current density. The maximum B content obtainable in this bath was 24 at.% B. Figure 7.4.2.30.3 depicts a graph for the B content versus the bath pH and current density. The B content increased with increasing pH but reached a plateau for all the current densities (10, 30, and 500 A/m2). Figure 7.4.2.30.4 shows the effect of the addition of TMAB on the B content in a plated Ni –B alloy film. The B content increased with increasing amount of TMAB.

Table 7.4.2.30.1. The composition of a Ni–B alloy plating bath and its plating condition. Composition

Concentration (g/l)

Plating condition

NiSO4·6H2O NiCl2·6H2O (CH3)3NBH3

240 45 30

pH Current density Bath temperature

599

5.0–10 10–500 A/m2 55 ^ 1 8C

600

Nano-Plating

Figure 7.4.2.30.1. A schematic diagram illustrating an experimental setup for plating Ni–B alloy films using a rotating electrode.

Figure 7.4.2.30.2. The B content of plated Ni– B alloy films is plotted as a function of the current density.

Electrolytic Ni – B

Figure 7.4.2.30.3. The B content of plated Ni– B alloy films versus the bath pH and current density.

Figure 7.4.2.30.4. The effect of the addition of TMAB on the B content of plated Ni– B alloy films.

601

602

Nano-Plating


7.4.2.30.3.1

Structure

X-ray diffraction patterns from Ni – B alloy films plated under various plating conditions are shown in Figure 7.4.2.30.5. The film became amorphous when the B content was high, but if the B content was low, the film became an assembly of fine Ni crystals, which exhibited the strong {111} diffraction peak. From these diffraction patterns it is difficult to determine at which composition the film became amorphous. The composition range together with the equilibrium phase diagram of Ni– B alloys are presented in Figure 7.4.2.30.6. The film having the composition near the intermetallic compound Ni3B appears to become amorphous and becomes an assembly of fine Ni crystals as the B content decreases. From the presence of the strong {111} Ni peak, it can be predicted that the fine Ni crystals are needle-shaped and their long direction is aligned along the film thickness direction.

7.4.2.30.3.2

Surface morphology

Figure 7.4.2.30.7(a) is the fracture surface of a crystalline Ni –B alloy film containing low B content, exhibiting a ductile fracture morphology. The fracture surfaces of the films with high B content are smooth like that of a glass, as a result of undergoing a brittle fracture mode.

Figure 7.4.2.30.5. X-ray diffraction patterns from plated Ni –B alloy films containing various B contents.

Electrolytic Ni – B

603

Figure 7.4.2.30.6. (a) The equilibrium phase diagram of Ni–B alloys and (b) the composition range for forming an amorphous phase in plated Ni–B alloy films.

Figure 7.4.2.30.7. The fracture surfaces of (a) crystalline and (b) non-crystalline Ni– B alloy films.

604

Nano-Plating

REFERENCES Onod, M., Tsuchiya, T., Ogawa, K. & Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 388. Onod, M., Tsuchiya, T., Shimizu, S. & Watanabe, T. (1992) J. Surf. Finish. Soc. Jpn, 43, 138. Onod, M., Shimizu, K., Tateishi, Y. & Watanabe, T. (1998) Trans. IMF, 76, 41.

FURTHER READING Onod, Onod, Onod, Onod,

M., Shimizu, M., Shimizu, M., Shimizu, M., Shimizu,

K., Tsuchiya, K., Tsuchiya, K., Tsuchiya, K., Tsuchiya,

T. & T. & T. & T. &

Watanabe, Watanabe, Watanabe, Watanabe,

T. (1992) J. Surf. Finish. Soc. Jpn, 43, 862. T. (1993) J. Magn. Magn. Mater., 126, 595. T. (1994) J. Surf. Finish. Soc. Jpn, 45, 714. T. (1995) Mater. Trans., JIM, 36, 1104.

7.4.2.31 Electrolytic Ni– Mo (Watanabe et al., 1989) 7.4.2.31.1. PLATING METHOD

The main constituents of a Ni– Mo alloy plating bath are two metal salts, nickel sulfate (NiSO4·2H2O) and sodium molybdate (Na2MO4·4H2O). The mixing ratio of NiSO4·2H2O and Na2MO4·4H2O salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.26 mol/l. Sodium citrate (C6H8O7Na3·2H2O) was used as a complexing agent. The bath pH was adjusted with dilute sulfuric acid or ammonium water. A detailed structural study was performed on Ni– Mo alloy films plated from the baths whose pH was 3 and 5. The bath temperature was 33 8C. The composition of the alloy films was determined using an atomic absorption method. For a structural analysis of these alloy films, an X-ray diffraction technique was applied using the cobalt Ka-line as an X-ray source.


Figure 7.4.2.31.1 presents the Mo content in plated Ni– Mo alloy films as a function of the Mo concentration in the bath. Two bath pH values of (A) 3 and (B) 5 and three current

Figure 7.4.2.31.1. The Mo content of plated Ni –Mo alloy films plotted as a function of the Mo concentration in the bath.

605

606

Nano-Plating

densities (400, 600, and 800 A/m2) were used. The dotted line in the graph represents the ratio of the Mo content of the film to the Mo concentration of the bath being equal to 1. The Mo content appears to increase with increasing Mo concentration in the bath but does not appear to be affected by the magnitude of the current density significantly. Figure 7.4.2.31.2 is a graph of the Mo content in the film plotted against the bath pH. These baths contained 60 mol./l sodium molybdate and the plating was carried out at 400 A/m2. The Mo content changed markedly with the pH and displayed the maximum at pH ¼ 5:


X-ray diffraction patterns from various compositions of Ni– Mo alloy films plated at 400 A/m2 in the pH ¼ 3 bath are shown in Figure 7.4.2.31.3. In addition to the peaks from the alloy films, strong diffraction peaks appeared from the Cu substrate. Ignoring these substrate peaks, we note that the diffraction peaks of the film increasingly broadened with

Figure 7.4.2.31.2. Relationship between the Mo content in Ni–Mo alloy films and the bath pH. The bath having the ratio of Ni : Mo ¼ 6 : 4 was used.

Electrolytic Ni – Mo

607

Figure 7.4.2.31.3. X-ray diffraction patterns from various compositions of plated Ni–Mo alloy films.

increasing Mo content, indicating that the crystals in the film were being continuously refined. It is difficult to ascertain at which composition the film became amorphous. From

608

Nano-Plating

the disappearance of crystalline peaks in the X-ray diffraction patterns, we assumed that the film was already amorphous at the composition of 35 at.% Mo. The structure (crystalline or amorphous) of Ni –Mo alloy films plated from the baths having pH ¼ (A) 3 and (B) 5 are plotted in Figure 7.4.2.31.4 as a function of the current density and the Mo concentration in the bath. Although an unambiguous determination of the film crystallinity is not an easy matter, it is safe to state that the film becomes amorphous when the Mo concentration in the bath is high. As explained in Figure 7.4.2.31.1, the Mo content increases proportionally with increasing Mo concentration in the bath regardless of the current density. The results in Figure 7.4.2.31.4, therefore, confirm that the Mo content in the film controls the amorphous structure of the film. We summarize the structure data in Figure 7.4.2.31.5 together with the equilibrium phase diagram of Ni – Mo alloys. It is important to note that the highest obtainable Mo content in this experiment was 80 at.% Mo. The film appears to be amorphous between 30 and 80 at.% Mo content. The film becomes amorphous near the composition at which their intermetallic compounds are present. Figure 7.4.2.31.6 contains SEM micrographs showing the surface morphology and fracture surface of three different Ni – Mo alloy deposits, which were determined to be amorphous. Both the surface morphology and fracture surface are seen to change with the composition.

Figure 7.4.2.31.4. The structure (crystalline or amorphous) of Ni–Mo alloy films plated from various baths having pH ¼ (A) 3 and (B) 5.

Electrolytic Ni – Mo

609

Figure 7.4.2.31.5. (a) The equilibrium phase diagram of Ni –Mo alloys and the composition range of forming an amorphous phase in plated Ni–Mo alloy films.

REFERENCE Watanabe, T., Naoe, T., Mio, J. & Katsumata, A. (1989) J. Surf. Finish. Soc., Jpn, 40, 458.

610

Nano-Plating

Figure 7.4.2.31.6. SEM micrographs showing the surface morphology and fracture surface of three different Ni– Mo alloy deposits.

7.4.2.32. Electrolytic Ni– P (Ito et al., 2001; Watanabe et al., 1989, 1994) 7.4.2.32.1. PLATING METHOD

The compositions of the Ni –P alloy plating bath used in this study are listed in Table 7.4.2.32.1. The bath, which is similar to the Watts bath, contains 0.9 mol/l NiSO4·6H2O and 0.1 mol/l NiCl2·H2O. In addition, 0.5 mol/l H2BO3 was added as a buffer and 0:04 , 0:5 mol=l phosphite (H3PO3) as a source of P. The concentration of P in the Ni– P films was varied by changing the quantity of phosphite, the current density, or the bath pH. In some cases, hypophosphite can be added instead of phosphite,or both chemicals can be used together. The bath temperature was either at an ambient temperature or 30 8C.


A relationship between the P content in plated Ni – P alloy films and the current density is plotted in Figure 7.4.2.32.1 as a function of the concentration of phosphite in the bath. At low current densities, the P content is high when the amount of phosphite is high in the bath. In addition to the trend seen in Figure 7.4.2.32.1, the P content was high if the film was grown in the low-pH (acidic) bath.


Figure 7.4.2.32.2 consists of X-ray diffraction patterns from various compositions of plated Ni –P alloy films, which were obtained by changing the current density and the concentration of phosphite in the bath. The diffraction peaks became broader with increasing P, i.e. Ni grains were continuously refined as the P content was increased and

Table 7.4.2.32.1. The composition of a Ni–P alloy plating bath. Composition


NiSO4·6H2O NiCl2·H2O H3BO3 H3PO3

0.95 0.17 0.32 0.04–0.5

611

612

Nano-Plating

Figure 7.4.2.32.1. The P content of plated Ni– P alloy films plotted as a function of the current density and the phosphite concentration in the bath.

finally became amorphous. Cross-section TEM micrographs for these films are shown in Figure 7.4.2.32.3– 7.4.2.32.10, together with their electron diffraction patterns. The grain size of each film can be easily estimated from the dark-field images. Figure 7.4.2.32.3 is a bright-field (BF)/dark-field (DF) images and their electron diffraction pattern showing the cross-section of a plated Ni film containing no P. Since the specimen was prepared by a microtome method, it is compressed laterally and consequently the grains are stretched along the vertical direction. Despite the damage, the presence/absence of compositional or structural modulations along the thickness direction can be easily detected from these microtomed specimens. Figure 7.4.2.32.3(a) is the cross-section TEM micrograph of the Ni film, taken at low magnification. The BF and DF images of the near-surface region taken at a higher magnification are seen in Figure 7.4.2.32.3(b) and (b0 ), respectively. The film thickness was about 19 mm and its structure was very uniform along the film thickness direction. From the DF image in Figure 7.4.2.32.3(b0 ), the average grain size was determined to be about 100 nm. Figure 7.4.2.32.4 is the BF image of a plated Ni –P alloy film containing few at.% P and four electron diffraction patterns taken from four sites (marked with symbols, (a)– (d)) along the film thickness direction. The structure of this film appears to be uniform along

Electrolytic Ni – P

613

Figure 7.4.2.32.2. X-ray diffraction patterns from plated Ni– P alloy films containing various amounts of P.

the thickness direction. The near-surface and near-interface regions in Figure 7.4.2.32.4 were further observed in the BF and DF imaging modes at higher magnification in Figure 7.4.2.32.5. The grain size was less than 10 nm in both regions. Figure 7.4.2.32.6 comprises the BF and DF images of a plated Ni – P alloy film containing 8 at.% P together with their electron diffraction patterns taken from five locations along the thickness direction. All the patterns except for those from

614

Nano-Plating

Figure 7.4.2.32.3. TEM micrographs showing the cross-section of a plated pure Ni film. The low-magnification image shown in (a) was magnified in (b) and (b0 ), which are the bright-field (BF)/dark-field (DF) images of the near-surface region.


615

Figure 7.4.2.32.4. The BF image of a plated Ni–P alloy film containing few at.% P and four electron diffraction patterns taken from four sites (marked with symbols, (a)–(d)) along the film thickness direction.

the substrate region display uniform diffraction rings, which are an indication that the grains were very small. Both BF and DF images were taken at higher magnification from both the near-surface and near-interface regions of the cross-section shown in Figure 7.4.2.32.6 (see Figure 7.4.2.32.7). Consistent with the X-ray diffraction results, the grain size was indeed small. Figure 7.4.2.32.8 is a cross-section TEM micrograph for a plated Ni –P alloy film containing the nominal composition of 24 at.% P. There is a 50-nm thick layer of finegrained Ni present immediately above the deposit –substrate interface. This crystalline

616

Nano-Plating

Figure 7.4.2.32.5. The near-surface and near-interface regions of the film shown in Figure 7.4.2.32.4 were further taken at a higher magnification in the BF and DF imaging modes. (a) and (a0 ) are a BF and DF pair for the nearsurface region, whereas (b) and (b0 ) are a BF and DF pair for the interface region.


617

Figure 7.4.2.32.6. The (a) BF and (a0 ) DF images of a plated Ni–P alloy film containing 8 at.% P together with their electron diffraction patterns taken from five locations ((b)–(f)) along the thickness direction.

618

Nano-Plating

Figure 7.4.2.32.7. The near-surface and near-interface regions of the film shown in Figure 7.4.2.32.6 were further taken at a higher magnification in the BF and DF imaging modes. (a) and (a0 ) are a BF and DF pair for the nearsurface region, whereas (b) and (b0 ) are a BF and DF pair for the interface region.


619

Figure 7.4.2.32.8. A cross-section TEM micrograph for a plated Ni –P alloy film containing the nominal composition of 24 at.% P and electron diffraction patterns taken from three locations, i.e. (a) the amorphous region (24 at.% P), (b) the fine crystalline region (9 at.% P), and (c) the Cu substrate.

620 Nano-Plating Figure 7.4.2.32. 9. A cross-section TEM micrograph for a plated Ni– P alloy film containing 25 at.% P. In contrast to the film seen in Figure 7.4.2.32.8, the film did not contain the initial layer of fine crystalline region. The TEM specimen was prepared by an ion-milling method.


621

Figure 7.4.2.32. 10. A cross-section TEM micrograph for a plated Ni-31.6 at.% P alloy film, electron diffraction patterns taken from three locations ((a)–(c)), and its X-ray diffraction pattern. The TEM specimen was prepared by an ion-milling method.

622

Nano-Plating

Figure 7.4.2.32. 11. (a) The equilibrium phase diagram of Ni –P alloys and (b) the phase diagram of plated Ni –P alloy films.

layer was followed by a thick amorphous layer, which is believed to represent the bulk of the film. The composition of the crystalline layer was analyzed using an energy dispersive X-ray spectrometer and was found to be 8.71 at.% P, which was much lower than the nominal composition (24 at.% P). The composition of the amorphous layer, was 24.4 at.% P. The 50-nm thick crystalline layer is considered to be the initial layer explained in Section 1.4.7(e). Figure 7.4.2.32.9 is a cross-section TEM micrograph for a plated Ni –P alloy film containing 25 at.% P. This film shows a uniform amorphous structure along the thickness direction, without any trace of the initial layer. Here, the formation of an initial layer was successfully avoided by agitating the bath violently during the plating.


623

Electron diffraction patterns were taken inside the amorphous region and of the deposit – substrate interface as seen in Figure 7.4.2.32.9(a) and (b), respectively. Finally, an electron diffraction analysis of the 31.6 at.% P Ni – P alloy film revealed the presence of NiP, Ni5P2, Ni3P, and Ni12P5 crystals in addition to an amorphous phase (see Figure 7.4.2.32.10). According to more detailed experimental results on electrolytic Ni –P alloy films (Ito et al., 2001), the distribution of P atoms was found to be very uniform. Furthermore, the P composition, at which Ni – P alloy films became amorphous, was found to be about 20 at.%. It is interesting to note that this composition is very close to 25 at.% P, the intermetallic compound, Ni3P. The above results are summarized in Figure 7.4.2.32.11(b) together with Figure 7.4.2.32.11(a) the equilibrium phase diagram of Ni (a)P alloys. For plated Ni –P films, the P content of , 20 at.% appears to define the phase boundary between crystalline and amorphous phases. Below the 20 at.% P, the films consist of fine Ni crystals. Above the 20 at.% P, the films become amorphous, but with a further increase in P, they will take various structural forms including their intermetallic compounds (Ito et al., 2001), NiP, Ni5P2, Ni3P, and Ni12P5.

REFERENCES Ito, K., Wang, F. & Watanabe, T. (2001) J. Jpn Inst. Metals, 65, 695. Watanabe, T. & Kanayama, T. (1989) J. Surf. Finish. Soc. Jpn, 40, 425. Watanabe, T. & Kanayama, T. (1994) J. Jpn Inst. Metals, 58, 138.

7.4.2.33. Electrolytic Ni – S (Narita et al., 1991) 7.4.2.33.1. PLATING METHOD

The compositions of the Ni – S alloy plating bath used in this study are listed in Table 7.4.2.33.1.While keeping the concentration of nickel sulfate constant (0.25 mol/l), we changed the concentration of sodium thiosulfate from 0.01 to 0.25 mol/l. The current density was also changed to obtain various compositions of Ni –S alloys. The bath pH was varied from 5 to 8, with sulfuric acid or ammonium water. The film thickness was controlled by keeping the total current constant (150 C). The plating was conducted under constant-potential conditions by varying the deposition potential from 2 460 to 2 640 mV. The amount of Ni in plated Ni – S alloy films was determined by an atomic absorption method. Since these alloy films contained only Ni and S, the S content was calculated by simply subtracting the Ni content from the film weight.


Figure 7.4.2.33.1 shows a graph for the S content of plated Ni –S alloy films plotted against the current density. Four baths containing different amounts ((a) 0.01, (b) 0.02, (c) 0.04, and (d) 0.25 mol/l) of sodium thiosulfate were used in this experiment. The S content generally decreased with increasing current density and with decreasing concentration of sodium thiosulfate. Figure 7.4.2.33.2 displays how the S content changes with the bath pH and the current density. For all the current densities, the S content slowly decreased with increasing bath pH. In Figure 7.4.2.33.3, we plotted the S content as a function of the concentration of sodium citrate in the bath and the current density. This bath contained 0.02 mol/l sodium thiosulfate. The S content decreased with increasing concentration of sodium citrate, which was used as a complexing agent and buffer for this bath. Table 7.4.2.33.1. The composition of a Ni– S alloy plating bath. Composition


NiSO4·6H2O (NH4)SO4 Na3C6H5O7·5H2O Na2S2O3·2H2O

0.25 0.25 0.05 0.02

624

Electrolytic Ni – S

625

Figure 7.4.2.33.1. The S content of plated Ni–S alloy films plotted as a function of the current density and the amount of Na2SO4.


Figure 7.4.2.33.4 shows X-ray diffraction patterns for plated Ni –S alloy films obtained at different current densities from the bath containing 0.02 mol/l sodium thiosulfate. These films were similar to those used for plotting Figure 7.4.2.33.1(c). The film plated at 15 A/ m2 contained 43 at.% S, and consisted of Ni and Ni3S2 crystals. The film obtained at 30 A/ m2 included 39 at.% S and its diffraction pattern exhibited a broad peak extending from 40 to 708. A further increase in the current density to 50, 100, 200, and 300 A/m2 reduced the S content to 30, 22, 15, and 10 at.%, respectively. At the same time, the intensity of the {111} Ni diffraction peak became increasingly stronger and the {200} Ni peak appeared. From these X-ray diffraction patterns, it is clear that the S content decreased with increasing current density. Except for the case of 15 A/m2, the structure of the Ni– S alloy films changed from amorphous to fine-grained phases with increasing current density and eventually became an assembly of fine Ni crystals. It has been shown above that different compositions of Ni– S alloy films can be prepared by changing the bath composition and plating conditions. The crystal structure can be

626

Nano-Plating

Figure 7.4.2.33.2. The S content of plated Ni–S alloy films plotted as a function of the bath pH and the current density.

Figure 7.4.2.33.3. The S content plotted as a function of the amount of sodium citrate (Na3C6H5O7) added to the bath and the current density.

Electrolytic Ni – S

627

Figure 7.4.2.33.4. X-ray diffraction patterns for plated Ni–S alloy films obtained at different current densities from the bath containing 0.02 mol/l sodium thiosulfate.

determined directly by the S content in the film but not by the plating conditions. When the S content reached a critical value of 40 at.%, the film changed from a crystalline to an amorphous phase. The composition range, in which plated Ni– S films had an amorphous structure, is plotted in Figure 7.4.2.33.5(b). The equilibrium phase diagram of Ni– S alloys is shown for comparison in Figure 7.4.2.33.5(a). It is interesting to note that the critical S content of 40 at.% corresponds to the stoichiometry of its intermetallic compound, Ni3S2. A further increase in the S content to 43 at.% indeed produced alloy films containing Ni3S2 crystals. Figure 7.4.2.33.6 displays SEM micrographs showing the surface morphology and fracture surface of an amorphous Ni –S alloy film. The surface exhibits a characteristic nodule structure, whereas the fracture face shows a smooth surface, which was formed by brittle fracture.

628

Nano-Plating

Figure 7.4.2.33.5. (a) The equilibrium phase diagram of Ni– S alloys and (b) the composition range of forming an amorphous phase in plated Ni –S alloy films.

Figure 7.4.2.33.6. SEM micrographs showing the surface morphology and fracture surface of an amorphous Ni –S alloy film.

Electrolytic Ni – S REFERENCE Narita, A. & Watanabe, T. (1991) J. Surf. Finish. Soc. Jpn, 42, 559.

629

7.4.2.34. Electrolytic Ni – Sn (Tanabe and Shimizu, 1995; Watanabe and Kaneyamn, 1994; Watanabe et al., 1999a,b, 2000) 7.4.2.34.1. PLATING METHOD

A Ni– Sn alloy plating bath developed by Enomoto et al. (Tanabe and Shimizu, 1975; Watanabe et al., 1999a,b, 2000) was used here (see Table 7.4.2.34.1). The mixing ratio of nickel chloride (NiCl2·2H2O) and tin chloride (SnCl2·2H2O) salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.15 mol/l. Pyrophosphate and glycine were employed as a complexing agent. The bath pH was adjusted to 8.1 with KOH or HCl. Annealed polycrystalline Cu sheets were used as the substrate. The plating was conducted at 50 A/m2 under the constant current condition.


In Figure 7.4.2.34.1, the Sn content of plated Ni – Sn alloy films and the current efficiency are plotted against the Sn concentration of the bath. Although the Sn content increased with increasing Sn concentration in the bath, it did not follow the dotted linear line, which represents a 1:1 composition relationship. In the Ni –rich bath, Sn deposited preferentially, whereas in the Sn-rich bath, Ni grew preferentially. From the bath containing 40– 80 at.% Sn, the Sn content of the film did not change significantly, which is consistent with the results reported by Tanabe and Shimizu (Enomoto and Fujiwara, 1981) and Enomoto et al. (Watanabe et al., 1999a,b, 2000). The current efficiency was 85% from the Ni –rich bath but was almost 100% for the rest of the concentration range. Based on this result, we plated various compositions of Ni– Sn alloy films.

Table 7.4.2.34.1. The composition of Ni–Sn alloy plating baths. Composition


NiCl2·2H2O SnCl2·2H2O K4P2O7 Glycine

(0.15 2 X) X 0.45 0.15

630

Electrolytic Ni– Sn

631

Figure 7.4.2.34.1. The Sn content of plated Ni–Sn alloy films and their current efficiency plotted against the Sn concentration of the bath.


7.4.2.34.3.1 Structure Figure 7.4.2.34.2 displays X-ray diffraction patterns from different compositions of plated Ni –Sn alloy films. Figure 7.4.2.34.2(a) shows diffraction patterns for the films having large compositional ranges (10:6 , 97:1 at:% Sn) plotted over large diffraction angles ð2u ¼ 20 , 908Þ; whereas Figure 7.4.2.34.2(b) presents detailed diffraction patterns plotted over narrower diffraction angles ð2u ¼ 25 , 508Þ for the Sn-rich films. All the films were plated for 30 min, which corresponds to a film thickness of , 66 mm. If the current efficiency was assumed to be 100%, the deposition rate then becomes 37 nm/s ð¼ 66 £ 1000=ð30 £ 60ÞÞ: These X-ray results are summarized in Figure 7.4.2.34.3. The phase diagrams of plated Ni– Sn alloy films produced by our group, Tanabe and Shimizu (Enomoto and Nakagawa, 1976), and Enomoto et al. (Bennet and Tompkins, 1976) are shown in Figure 7.4.2.34.3(b) – (d), respectively, whereas the corresponding equilibrium phase diagram [34,35] is listed in Figure 7.4.2.34.3(a) for comparison. The difference between the phase diagram of our plated Ni –Sn alloy films and the standard equilibrium phase diagram will be considered from the Ni –rich side. Very few Sn atoms can dissolve in pure Ni at ambient temperatures [34,35]. Plated Ni films, however, were found to dissolve up to 30 at.% Sn as a Ni solid-solution alloy. The exact solubility limit of the Sn content, could not be easily determined because the diffraction peaks of the Ni solid solution began to overlap with those of a meta-stable phase, which will be described below.

632

Nano-Plating

The lattice constant of the plated Ni solid-solution alloy having the Sn content of 12.2 at.% was 0.361 nm, which was larger than that (0.352 nm) of a pure Ni deposit. This increased lattice constant agrees with the lattice constant value (0.36084 nm (Nash and Nash, 1985)) extrapolated from that of a supersaturated Ni solid-solution alloy (Ni – 10.80 at.% Sn), which was obtained by a splat-quenching method. For this reason,

Figure 7.4.2.34.2. X-ray diffraction patterns from different compositions of plated Ni–Sn alloy films.


633

the plated Ni– 12.2 at.% Sn alloy film can be regarded as a super-saturated solid solution alloy of Ni (Jette and Fets, 1935; Mikulas and Thomassen, 1934; Pearson and Thompson, 1957; Polesya and Slipchenko, 1972; Klement, 1962) and thus is in a meta-stable state. This meta-stable phase was marked with a symbol, Ni(M), in Figure 7.4.2.34.3. The use of a liquid-quenching technique has been reported (Polesya and Slipchenko, 1972) to extend the Sn solubility up to 16.8 at.%. In the X-ray diffraction patterns of the 10 , 60 at:% Sn alloy films, there is a broad diffraction peak present between the diffraction angles of 20 and 408 (see Figure 7.4.2.34.2). This broad peak suggests the presence of an amorphous phase. Electroplated binary alloy films are known to form an amorphous phase near the composition of their intermetallic compound. Under normal equilibrium condition, the intermetallic compound of the b phase is supposed to form in this composition range ð10 , 60 at:%SnÞ; but instead the pattern contained only the broad peaks from the Ni(M) phase, and peaks from a metastable phase (M1) (which will be described below). Amorphous Ni(M) and meta-stable M1 phases may coexist in this alloy deposit. Diffraction peaks labeled as M1 in Figure 7.4.2.34.2(a) originate from a meta-stable phase reported previously (Mikulas and Thomassen, 1934; Pearson and Thompson, 1957; Dutta and Clarke, 1968; Tanabe and Shimizu, 1975; Bennet and Tompkins, 1976; Enomoto et al., 1981, 1982; Enomoto and Fujiwara, 1981; Tamura and Yasuda, 1986; Izaki et al., 1987a,b; Watanabe and Kaneyama, 1994; Watanabe et al., 1999a,b, 2000). This phase was described in detail by Dutta and Clarke (Watanabe et al., 1999) who reported that the structure of this phase is of hcp NiAs type having the composition of 50 at.%Ni –50 at.% Sn. It was also stated (Dutta and Clarke, 1968) that the composition ratio


634

Nano-Plating

Figure 7.4.2.34.3. (a) The equilibrium phase diagram of Ni–Sn alloys (Massalski, 1990); (b), (c), and (d) are the phase diagrams of plated Ni–Sn alloy films obtained by Watanabe et al., Shimizu et al., Enomoto and Nakagawa (1976), and Enomoto et al., Watanabe et al., (2000), respectively.


635

of Ni to Sn can be changed in this phase. A chemical shift for the M1 phase was also measured using Mo¨ssbauer spectroscopy (Silver et al., 1976). The M1 phase was reported to appear in Ni –Sn alloy films obtained by a sputtering method (Augis, 1977). It is interesting to note that the M1 phase appears in the composition range of 18 , 60 at.% Sn, where the intermetallic compounds, b (Ni3Sn) and g (Ni3Sn2) phases form according to the equilibrium phase diagram (Nash and Nash 1985; Thaddeus and Massalski, 1990) (see Figure 7.4.2.34.3(a)). In the composition range of 18 , 30 at:% Sn, meta-stable phases of Ni(M) and M1 coexist. Note that the composition range of the M1 phase determined in the present study is different from the range found by Tanabe and Shimizu (Enomoto and Fujiwara, 1981) as well as by Enomoto et al. (Bennet and Tompkins, 1976) (see Figure 7.4.2.34.3(b) – (d)). Two unknown peaks marked with a symbol,?, appeared in the 55.7 at.% Sn film (see Figure 7.4.2.34.2(a)). These peaks were strong when the peaks of the M1 phase were strong. Conversely, the unknown peaks were weak if the M1 peaks were weak. The phase associated with these unknown peaks, therefore, must be related to the M1 phase. It is possible that these unknown peaks could originate from the ordered structure of the M1 phase. The d phase (Ni3Sn4) listed in the equilibrium phase diagram was found in plated Ni– Sn alloy films in the composition range of 56 , 89 at:% Sn. The d phase forms at a composition of 57 at.% Sn under the equilibrium condition. The d phase of plated Ni – Sn alloys, however, exhibited not only the compositional shift, but also extended solubility. Consequently, the M1 and d phases coexist in the composition range of 56 , 60 at:% Sn. According to Tanabe and Shimizu (Enomoto and Fujiwara, 1981) and Enomoto et al. (Bennet and Tompkins, 1976), there is no composition range where the d phase exists singly. The present study, however, does indicate that the d phase could exist singly in the composition range of 60 , 75 at:% Sn. In the composition range of 75 , 100 at:%; Sn crystals appeared. There is no Ni solubility in Sn at ambient temperatures (Nash and Nash, 1985; Thaddeus and Massalski, 1990) (see Figure 7.4.2.34.3(a)). The lattice constant of the Sn crystal in this composition range was found to be the same as the one listed in the JCPDS file. Consequently, we concluded that the Sn crystal did not contain Ni, consistent with the equilibrium phase diagram. Diffraction peaks marked with a symbol, M2, shown in Figures 7.4.2.34.2(b) and 7.4.2.34.3(b) have not been previously reported. Figure 7.4.2.34.4 is the effect of annealing (100 8C) time on the X-ray diffraction pattern of a plated Ni– 87.1 at.% Sn film. Diffraction peaks from the M2 phase became weaker with the annealing time and disappeared after about 360 min. It can be inferred from this annealing behavior that those diffraction peaks do indeed belong to one crystal, i.e. the M2 phase. The M2 phase, which does not appear under the equilibrium condition, is a meta-stable phase because it disappears with annealing. The M2 phase appeared in the composition range of 75 , 90 at:% Sn. Thus, the d phase, M2 phase, and Sn crystal coexist in the composition range of 75 , 90 at:% Sn,

636

Nano-Plating

Figure 7.4.2.34.4. A structural change in the meta-stable M phase of a plated Ni–87.1 at.% Sn film upon heating at 100 8C.

beyond which only Sn crystal was present. As described above, Ni does not dissolve in Sn and consequently it segregates at the Sn grain boundaries. 7.4.2.34.3.2 Surface morphology Figure 7.4.2.34.5 shows the surface morphology and fracture surfaces of plated Ni– Sn alloy films with various Sn contents. The cracks seen in each film were introduced deliberately in order to study the surface and fracture surface morphologies simultaneously by SEM. Only the 98.8 at.% Sn film exhibited a ductile fracture behavior, but the rest displayed brittle fracture morphologies. Up to 43.6 at.% Sn, the surface was smooth but beyond 56.3 at.%, the surface became very rough. 7.4.2.34.3.3 Cross-section TEM Figure 7.4.2.34.6 shows TEM micrographs showing the cross-section specimens of various electrolytic Ni – Sn alloy films, which were prepared by microtoming. Although some specimens were cracked due to their high internal stress during microtoming, all the specimens exhibit a uniform structure along the film thickness direction. The composition was measured at three locations along the film thickness direction using a 2-mm-diameter beam of an energy dispersive X-ray spectrometer attached to the TEM. This composition data is shown in Figure 7.4.2.34.7, where the Sn content of plated Ni –Sn alloy films was plotted as a function of the Sn concentration of the bath and the measured locations. The composition near the surface, and at the middle point between the top surface and the

Electrolytic Ni– Sn Figure 7.4.2.34.5. Relationship between the surface morphology and fracture surface of plated Ni–Sn alloy films containing various amounts of Sn and the phase diagram.

637

638

Nano-Plating

Figure 7.4.2.34.6. TEM micrographs showing the cross-section specimens of various Ni–Sn alloy films.


639

Figure 7.4.2.34.7. The Sn composition of Ni– Sn alloy films plotted as a function of the Sn content of the films. The Sn composition was measured at three locations of the cross-section along the film thickness direction (see Figure 7.4.2.34.6).

deposit – substrate interface, is the same as the composition shown in Figure 7.4.2.34.1. For the Ni – and Sn-rich alloy films, the composition near the deposit-substrate interface is different from the one presented in Figure 7.4.2.34.1, i.e. the initial layer of the films is different. The Ni content is higher in the Sn-rich film than the bulk value, and, conversely, the Sn content is higher in the Ni– rich film. As described in Figure 7.4.2.34.1, these composition deviations are connected to the fact that Sn is diffusion-controlled in the Ni– rich bath and Ni is diffusion-controlled in the Sn-rich bath. 7.4.2.34.3.4 Crystal size The crystal (grain) sizes of various plated Ni– Sn alloy films were measured directly from the dark-field images of TEM micrographs seen in Figure 7.4.2.34.6. An X-ray diffraction technique was also employed to determine the grain size from the full width at the half maximum (FWHM) of a diffraction peak using Scherrer’s equation. These results are shown in Figure 7.4.2.34.8, where the type of a crystal used for the X-ray measurement is also indicated. Since the grains of the films were severely deformed during specimen

640

Nano-Plating

Figure 7.4.2.34.8. The grain size of plated Ni–Sn alloy films plotted as a function of the Sn content of the films. Both TEM and X-ray diffraction techniques were used to determine the grain size.

preparation by microtoming, they could not provide absolute sizes. A relative comparison, however, can be made from those measured values. It is interesting to note that the grain size obtained by the X-ray measurements shows a similar trend to the one by TEM. The grain size is seen to increase with increasing Sn content in the film and takes the maximum value at the composition of 55 at.% Sn. Interestingly, this composition is very close to the composition (56.3 at.% Sn) of the film which exhibited very rough surface (see Figure 7.4.2.34.5), as well as to the composition at which the M1 peaks appeared most strongly (see Figure 7.4.2.34.2). From these results, we can conclude that the meta-stable phase was most stable at the composition of 55 at.% Sn and thus could grow large. The crystal is smallest at the composition of 20 , 30 at:% Sn, where the Ni(M) and M phases coexist. Although the existence of an amorphous phase is possible as discussed previously, it was not possible to detect such a phase from the TEM micrographs presented in Figure 7.4.2.34.6(A) – (G).

REFERENCES Augis, J.A. & Bennet, J.E. (1977) J. Electrochem. Soc., 124, 1455. Bennet, J.E. & Tompkins, H.G. (1976) J. Electrochem. Soc., 123, 999.


641

Dutta, P.K. & Clarke, M. (1968) Trans. Inst. Metal Finish., 46, 20. Enomoto, H. & Fujiwara, Y. (1981) J. Metal Finish. Soc. Jpn, 32, 555. Enomoto, H. & Nakagawa, A. (1976) J. Metal Finish. Soc. Jpn, 27, 569. Enomoto, H., Fujiwara, Y. & Ishikawa, M. (1981) J. Metal Finish. Soc. Jpn, 32, 23. Enomoto, H., Ishikawa, M. & Fujiwara, Y. (1982) J. Metal Finish. Soc. Jpn, 33, 332. Izaki, M., Enomoto, H. & Omi, T. (1987a) J. Metal Finish. Soc. Jpn, 38, 189. Izaki, M., Enomoto, H. & Omi, T. (1987b) Plat. Surf. Finish., 74(June), 84. Jette, E.R. & Fets, E. (1935) Metallwirk. Wiss. Tech., 19, 165. Klement, W. (1962) Can. J. Phys., 40, 1397. Massalski, T.B. (1990) Binary Alloy Phase Diagrams, 2nd Edition, vol. 3, p. 2863. Mikulas, W. & Thomassen, L. (1934) Trans. AIME, 124, 111. Nash, P. & Nash, A. (1985) Bull. Alloy Phase Diagrams, 6, 350. Pearson, W.B. & Thompson, L. (1957) Can. J. Phys., 35, 349. Polesya, A.F. & Slipchenko, L.S. (1972) Izv. V. U. Z. Tsvetn. Metall., 15, 128. Schubert, R. (1973) J. Electrochem. Soc., 125, 1215. Silver, J., Mackary, C.A. & Donaldson, J.D. (1976) J. Mater. Sci., 11, 836. Tamura, T. & Yasuda, Y. (1986) J. Metal Finish. Soc. Jpn, 37, 406. Tanabe, Y. & Shimizu, Y. (1975) J. Metal Finish. Soc. Jpn, 26, 406. Thaddeus, G. & Massalski, B. (1990) ASM International, Binary Alloy Phase Diagrams, 3, 2863. Watanabe, T. (1990) J. Surf. Finish. Soc. Jpn, 41, 652. Watanabe, T. & Kaneyama, T. (1994) J. Jpn Inst. Metals, 58, 138. Watanabe, T., Arai, K., Hirose, T. & Chikazawa, M. (1999a) J. Jpn Inst. Metals, 63, 489. Watanabe, T., Hirose, T., Arai, K. & Chikazawa, M. (1999b) J Jpn Inst. Metals, 63, 496. Watanabe, T., Arai, K. & Hirose, T. (2000) J. Jpn Inst. Metals, 64, 242.

FURTHER READING Antler, M., Feder, M., Horing, C.F. & Bohland, J. (1976) Plat. Surf. Finish., 63, 30. Augis, J.A. & Bennet, J.E. (1978) J. Electrochem. Soc., 125, 330. Clarke, M. & Dutta, P.K. (1971) J. Phys. D: Appl. Phys., 4, 1652. Beltowska-Lehman, E. (1982) Surf. Technol., 15, 191. Bennet, J.E. & Tompkins, H.G. (1976) J. Electrochem. Soc., 123, 1003. Culbertson, J.W., Parkinson, N. & Rooksby, H.P. (1951) J. Electrochem. Soc., 17, 107. Enomoto, H. & Ishikawa, M. (1979) J. Metal Finish. Soc. Jpn, 30, 284. Enomoto, H., Fujiwara, Y., Isaki, M. & Ono, H. (1982) J. Metal Finish. Soc. Jpn, 33, 369. Lay, D.E. (1988) Plat. Surf. Finish., 75, 26. Lo, C.C. (1980) J. Appl. Phys., 51, 2007. Nelson, G.C. (1978) J. Electrochem. Soc., 125, 403. Tamura, T. & Hosokawa, K. (1977) J. Metal Finish. Soc. Jpn, 28, 564. Tamura, T. & Yasuda, S. (1984) Metal Finish., 82, 8. Tamura, T. & Yasuda, Y. (1987) J. Metal Finish. Soc. Jpn, 38, 518. Wilson, G.C. (1972) Trans. Inst. Metal Finish., 50, 109. Wynne, B.E., Edington, J.W. & Rothwell, G.P. (1972) Met. Trans., 3, 301. Yamashita, H., Yamamura, T. & Yoshimoto, K. (1994) Denki Kagaku, 62, 48.

7.4.2.35. Electrolytic Ni – W (Ito et al., 2001) 7.4.2.35.1. PLATING METHOD

The compositions of the Ni– W alloy plating bath used in this study are listed in Table 7.4.2.35.1.The mixing ratio of nickel sulfate (NiSO4·6H2O) and sodium tungstate (Na2WO4·6H2O) salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant at 0.32 mol/l. Varying the current density also yielded a variety of film compositions. Citric acid was added as a complexing agent and the bath pH was adjusted to 6. The plating was conducted on electropolished polycrystalline Cu substrates for 30 min in the 70 8C bath, which was agitated with a magnetic stirrer. A Pt plate was chosen as an anode. The film composition was determined using an energy dispersive X-ray spectrometer attached to the SEM. 7.4.2.35.2. PLATING CONDITIONS AND FILM COMPOSITION

The W content of Ni– W films plated at 1500 A/m2 is plotted in Figure 7.4.2.35.1 against the W concentration of the bath. The W content in the films generally increased with increasing concentration of W in the bath. The W content of the films obtained from a bath with the W concentration of 0.29 mol/l is plotted in Figure 7.4.2.35.2 as a function of current density. The W content shows a maximum at 800 A/m2. By changing the mixing ratio of the two salts or the current density, we were able to change the composition of Ni – W deposits and obtained the maximum W composition of 35.1 at.%. 7.4.2.35.3. FILM MICROSTRUCTURE

Figure 7.4.2.35.3 shows X-ray diffraction patterns from various compositions of plated Ni– W films grown on polycrystalline Cu substrates. These patterns are somewhat complicated Table 7.4.35.1. The composition of Ni– W alloy plating baths and its plating condition. Composition

Plating condition

NiSO4·6H2O þ Na2WO4·2H2O 0.32 mol/l (NH4)2C4H4O6 0.32 mol/l

Bath temperature: 70 8C Current density: 250– 1,500 A/m2 pH: 6 Plating time: 30 min Agitation: yes Substrate: polycrystalline Cu Anode: Pt plate

642

Electrolytic Ni –W

643

Figure 7.4.2.35.1. The W content of Ni–W films plated at 1500 A/m2 plotted against the W concentration of the bath.

Figure 7.4.2.35.2. The W content of Ni–W films plated from a bath containing the W concentration of 0.29 mol/l is plotted as a function of current density.

644

Nano-Plating

Figure 7.4.2.35.3. X-ray diffraction patterns from various compositions of plated Ni–W films.


645

Figure 7.4.2.35.4. The shortest atomic distance of plated Ni –W alloy films plotted as a function of the W content in the films.

by the appearance of diffraction peaks from the Cu substrate. The lattice constant of Cu ða ¼ 0:3153 nmÞ is very close (only 2.6% difference) to Ni ða ¼ 0:3238 nmÞ: The films with low W content displayed a sharp {111} Ni peak and were crystalline. With increasing W content, however, diffraction peaks from Ni continuously decreased and became a broad peak, i.e. the film finally became an amorphous phase. The {111} Ni diffraction angle in the low W content side is seen to shift toward the lower angle (or the direction toward the lattice expansion) with increasing W content. This angle shift was plotted in terms of the lattice constant in Figure 7.4.2.35.4. Since the crystal structure of Ni (fcc) is different from W (bcc), we cannot compare the value of the lattice constant directly. Consequently, we compared the closest atomic distance. The plot in Figure 7.4.2.35.4 indicates a linear relationship, indicating that W is substitutionally alloyed in Ni. According to the equilibrium phase diagram in Figure 7.4.2.35.5(a), Ni can dissolve about 17.5 at.% W at 1495 8C but 13.5 at.% W at ambient temperatures. Therefore, any alloy deposits containing 13.5 – 20.0 at.% W are

646

Nano-Plating

Figure 7.4.2.35.5. (a) The equilibrium phase diagram of Ni–W alloys and (b) the phase diagram of plated Ni–W alloy films.

considered to be a supersaturated solid solution and thus a thermodynamically meta-stable phase. To further understand the structure of as-deposited Ni– W films, we conducted a heating experiment. Figure 7.4.2.35.6 consists of X-ray diffraction patterns from a plated Ni –W alloy deposit containing 19.7 at.% W, which was heat-treated at different temperatures. The height of the {111} Ni peak appears to increase with increasing temperature and at the same time other peaks from the Ni appeared, suggesting that Ni grains were growing. Furthermore, the diffraction angle of the {111} Ni did not shift after the 700 8C heating and thus the film remained in the state of a supersaturated solid solution. Similar heating experiments were conducted for Ni – W alloy deposits containing 24.4 at.% W. Figure 7.4.2.35.7 shows this result. The as-deposited film exhibited a broad diffraction


647

Figure 7.4.2.35.6. A change in the X-ray diffraction patterns of a plated Ni–W alloy deposit containing 19.7 at.% W upon heating at various temperatures.

648

Nano-Plating


peak, which did not change up to the 700 8C heating. After the 700 8C heat treatment, many diffraction peaks suddenly appeared. These peaks are not from Ni or Ni4W crystals and are not even listed in the JCPDS file. For this reason, these unknown crystals cannot be identified at the moment. Nevertheless, these crystals are undoubtedly the intermetallic


649


compounds, NixWy. Figure 7.4.2.35.8 consists of X-ray diffraction patterns from a Ni– W alloy film containing 27.6 at.% W, which was given the same heat treatment as the 19.7 and 24.4 at.% W alloy films. The as-deposited film exhibits a broad diffraction peak. Similar to the case of the 24.4 at.% W film, the broad peak did not change up to 700 8C, but many diffraction peaks suddenly appeared after the 700 8C heating. These peaks are

650

Nano-Plating

Figure 7.4.2.35.9. The surface morphology and fracture surface of plated Ni–W alloy films containing various amounts of W.

identical to those observed in the 24.4 at.% W film (see Figure 7.4.2.35.7). This behavior is strong evidence that the as-deposited 27.6 at.% W film was amorphous prior to the 700 8C heating. Based on these heating experiments, we can conclude that the as-deposited 19.7 at.% W film was crystalline, whereas the as-deposited 24.4 and 27.6 at.% W films were amorphous. The critical composition, at which the as-deposited Ni –W alloy films


651

transform from crystalline to amorphous phases, is considered to lie between 19.7 and 24.4 at.% W and is about 22.2 at.% W if the middle value is taken. This conclusion is included in the phase diagram of plated Ni– W alloy films in Figure 7.4.2.35.5(b). Although the Ni4W intermetallic compound is present in the equilibrium phase diagram (cf. Figure 7.4.2.35.5(a)), no amorphous phase was observed at this composition. An amorphous deposit was, however, obtained for the alloy composition of more than 22.2 at.% W. The surface morphology and fracture surface of plated Ni –W alloy films containing various amounts of W are shown in Figure 7.4.2.35.9. These films correspond to those used in Figure 7.4.2.35.3 for the X-ray study. The 12.5, 14.9, and 20.0 at.% W films were crystalline and displayed a ductile fracture mode, whereas the 27.8 and 35.1 at.% W films were amorphous and thus showed a typical brittle fracture.

REFERENCE Ito, K., Wang, F. & Watanabe, T. (2001) J. Jpn Inst. Metals, 65, 1028.

7.4.2.36. Electrolytic Ni – Zn 7.4.2.36.1. PLATING METHOD

A Ni – Zn alloy plating bath is a mixture of nickel sulfate and zinc sulfate. The first set of baths was made by changing the concentration of zinc sulfate while keeping the concentration of nickel sulfate constant (0.5 mol/l). Similarly, the second set of baths was prepared by varying the concentration of nickel sulfate while keeping the concentration of zinc sulfate constant (0.5 mol/l). Boric acid (20 g/l) and ammonium chloride (14 g/l) were added to the above solutions. The bath pH was adjusted to 4 by adding ammonium solution or dilute sulfuric acid. The cathode was a pure copper sheet, in which the planar section of 10 £ 60 mm2 was exposed for plating. A platinum plate ð10 £ 60 £ 0:25 mm3 Þ was used as an anode. The plating was conducted in the 30 8C unstirred bath until 60 C were consumed. A structural analysis of the alloy films was performed using an X-ray diffraction method. The alloy composition was determined as follows. Plated Ni– Zn alloy films were first dissolved in a mixed solution of sulfuric acid and nitric acid, followed by dilution in water. The amount of metals dissolved in the solution was then analyzed using an atomic absorption spectrophotometer.

7.4.2.36.2. PLATING CONDITIONS AND FILM COMPOSITION 7.4.2.36.3. FILM COMPOSITION AND STRUCTURE

Figure 7.4.2.36.1 consists of (a) the equilibrium phase diagram of Ni – Zn alloys and (b) a plot for a cathodic potential versus the alloy composition in terms of current densities (20, 50, 100, and 500 A/m2) and the type of phases (Ni solid solution, g phase, Zn solid solution, and amorphous phase). Alloy phases obtained in electroplated Ni– Zn alloy films are indeed present in the equilibrium phase diagram. It is interesting to note that neither the b or d phases appeared in the plated films and, instead, the g phase appeared over a wide range of compositions. At a current density of 20 A/m2, a phase marked with a symbol, X, emerged in the Ni-rich side. This phase did not show appreciable diffraction peaks, and thus was determined to be an amorphous phase. A phase diagram for electroplated Ni– Zn alloys is drawn over the whole composition range in Figure 7.4.2.36.1(c).

652

Electrolytic Ni– Zn

653

Fig. 7.4.2.36.1. (a) The equilibrium phase diagram of Ni–Zn alloys. (b) A plot for a cathodic potential versus the alloy composition in terms of current densities (20, 50, 100, and 500 A/m2) and the type of phases (Ni solid solution, g phase, Zn solid solution, and amorphous phase). (c) The phase diagram of plated Ni–Zn alloys.

654

Nano-Plating

REFERENCES Imai, Y. & Kurachi, M. (1977) Denki Kagaku, 45, 728. Imai, Y. & Kurachi, M. (1979) Denki Kagaku, 47, 89. Imai, Y., Watanabe, T. & Kurachi, M. (1978) Denki Kagaku, 46, 202. Kurachi, M. & Fujiwara, K. (1970) Denki Kagaku, 38, 600. Kurachi, M., Fujiwara, K. & Tanaka, T. (1972) Proceedings of the 8th International Union Electrodeposition & Surface Finishing, Interfinish Basel CH, Forster-Verlag AG, Zurich, p 152– 157.

7.4.2.37. Electrolytic Sn– Zn 7.4.2.37.1. PLATING METHOD

A Sn – Zn alloy plating bath consists of sodium stannate (Na2Sn) and zinc cyanide (Zn(CN)2). The mixing ratio of Na2Sn and Zn(CN)2 salts was changed to obtain various alloy compositions while keeping the total metal ion concentration constant. The substrates were iron plates with dimensions of 2 £ 1 cm2 : The plating was conducted at a current

Fig. 7.4.2.37.1. (a) The equilibrium phase diagram of Sn –Zn binary alloys and (b) the phase/lattice constants of plated Sn– Zn alloys.

655

656

Nano-Plating

density of 200 A/m2 from the 70 8C bath without any agitation. The structure of the deposits was studied using a Debye camera.

7.4.2.37.2. PLATING CONDITION AND FILM COMPOSITION

Unfortunately precise data were not provided in Aotani (1952).

7.4.2.37.3. FILM STRUCTURE

Figure 7.4.2.37.1 shows (a) the equilibrium phase diagram of Sn –Zn binary alloys and (b) the phase/lattice constant of Sn– Zn alloy deposits obtained by a plating method. According to the equilibrium phase diagram, the Sn – Zn binary system is a eutectic type and separates into two phases without any significant mutual solubility. In the plated alloys, bct Sn and hcp Zn crystals appeared and these two phases coexisted in the Zn composition range of 11 , 15 at.%. In addition, the lattice constants of these crystals did not change over the entire composition range, suggesting that they did not form solidsolution alloys. In other words, these alloys were thermally in an equilibrium condition and thus agreed well with the equilibrium phase diagram.

REFERENCE Aotani, K. (1952) Denki Kagaku, 20, 611.

7.4.3 Electroless Plating 7.4.3.1. Electroless Ni– B 7.4.3.1.1. PLATING METHOD

To obtain electroless Ni– B alloy films, either sodium borohydride (NaBH4) or n-dimethyl borazane (DMAB) ((CH3)2HNBH3) can be used as a reducing agent. We will describe both reducing agents here. Table 7.4.3.1.1 lists the compositions of electroless Ni – B baths using (a) NaBH4 (Tanabe and Watanabe, 1970, 1972a,b; Watanabe and Tanabe, 1974, 1975a,b, 1978) and (b) DMAB (Suda et al., 1981a,b) as a reducing agent. The baths using NaBH4 are commercial solutions, which were operated at 91 8C, and thus contain several additives. Nickel chloride is added as a metal salt, C2H4(NH2)2 as a complexing agent, NaOH as a buffer, arsenic acid (H3AsO4) as an accelerator, and Tl2SO4 as a stabilizer. The baths using DMAB as a reducing agent contain nickel sulfate as a metal salt, HOCH2COOH as a complexing agent, sodium acetate (CH3COONa) as a buffer, and 2-mercaptobenzothiazole (2-MBT) as a stress-reducing agent. These baths were operated at 70 8C.

Table 7.4.3.1.1. The composition of electroless Ni–B baths. Composition

Bath name

(a) Baths using NaBH4 as a reducing agent R800

R400

R200

R70

R0

NiCl2·6H2O C2H4(NH2)2 NaOH H3AsO4 Tl2SO4 NaBH4

30 60 40 200 30 400

30 60 40 200 30 200

30 60 40 200 30 70

30 60 40 200 30 0

30 60 40 200 30 800

(b) Baths using DMAB ((CH3)2HNBH3) as a reducing agent D-0.5 D-1

D-3

D-6

D-9

D-12

NiSO4·6H2O (kmol/m3) HOCH2COOH (kmol/m3) CH3COONa (kmol/m3) 2-MBT (kmol/m3) DMAB (kmol/m3) B content in Ni films (at.%)

0.114 0.342 0.122 0.0108 0.051 8.5

0.114 0.342 0.122 0.0108 0.102 12.7

0.114 0.342 0.122 0.0108 0.153 15.5

0.114 0.342 0.122 0.0108 0.204 16.2

0.114 0.342 0.122 0.0108 0.0085 4.0

0.114 0.342 0.122 0.0108 0.017 5.2

657

658

Nano-Plating

Figure 7.4.3.1.1. The weight change of Ni–B alloy deposits as a function of the deposition time, the bath temperature, and the amount of NaBH4 added to the bath.

Figure 7.4.3.1.2. The deposition rate and the B content of electroless Ni– B films plotted against the concentration of DMAB in the bath.

Electroless Ni – B

659

After electroless Ni– B films were stripped off the Cu substrates in a chromic acid solution, a quantitative analysis of the B content was made using a mannitol alkali neutralization titration method (Suda et al., 1981a). The Ni content was determined using a dimethyl glyoxime method (Suda et al., 1981b). Other alloying elements were studied using an electron probe micro-analysis (EPMA) technique.


As long as electroless Ni baths contain NaBH4 and DMAB as a reducing agent, boron atoms from the reducing agent will be incorporated into the Ni deposit, thus forming Ni – B alloy films. Conversely, pure Ni films without B cannot be obtained from these baths. 7.4.3.1.2.1 NaBH4 as a reducing agent Figure 7.4.3.1.1 plots the weight of Ni –B alloy deposits as a function of the deposition time. In this plot, the bath temperature was 80 and 91 8C, and the concentration of NaBH4

Figure 7.4.3.1.3. The deposition rate and the B content of electroless Ni–B films plotted as a function of the bath pH. The bath pH was adjusted with NaOH.

660

Nano-Plating

Figure 7.4.3.1.4. The deposition rate and the B content of electroless Ni–B films plotted against the amount of a buffer, CH3COONa, in the bath.

was 200 and 400 mg/l. The increase in the bath temperature from 80 to 91 8C resulted in a doubling of the deposition rate. 7.4.3.1.2.2 DMAB as a reducing agent Figure 7.4.3.1.2 depicts the deposition rate and B content of electroless Ni– B films plotted against the concentration of DMAB in the bath. With an increasing concentration of DMAB, both the deposition rate and B content appear to increase. Figure 7.4.3.1.3 depicts the deposition rate and B content of electroless Ni– B films plotted as a function of bath pH. With increasing bath pH, the deposition rate increased but the B content decreased. Figure 7.4.3.1.4 depicts the deposition rate and B content of electroless Ni– B films plotted against the concentration of buffer, CH3COONa, in the bath. The addition of CH3COONa slightly changed the deposition rate but did not change the B content. The deposition rate and B content of electroless Ni – B films are plotted as a function of the bath temperature in Figure 7.4.3.1.5. An increase in the bath temperature increased the deposition rate and increased the B content.


661

Figure 7.4.3.1.5. The deposition rate and the B content of electroless Ni–B films plotted as a function of the bath temperature.

The deposition rate and B content of electroless Ni –B films were plotted as a function of the concentration of a stress-reducing agent, 2-MBT, in the bath (see Figure 7.4.3.1.6). The addition of a small quantity ð0:0 , 2:0 mg=lÞ of 2-MBT appears to affect the deposition rate but a further addition of 2-MBT did not further change the deposition rate. The B content in the film was not affected by the addition of 2-MBT.

7.4.3.1.3. STRUCTURE

Figure 7.4.3.1.7 displays TEM micrographs and electron diffraction patterns showing the structure of electroless Ni – B films obtained from four baths containing NaBH4 as a reducing agent (see Table 7.4.3.1.1(a)). These four baths contained (a) 0, (b) 70, (c) 200, and (d) 400 g/l NaBH4 as a reducing agent. The B content in the film increased in the same order. For those films with a low B content, the structure appears to be grainy but for those with high B content, the structure became featureless. The electron diffraction patterns display spotty polycrystalline rings but change to a broad halo as the B content is increased. By comparison with the electron diffraction patterns, fine-grained polycrystal-

662

Nano-Plating

Figure 7.4.3.1.6. The deposition rate and the B content of electroless Ni–B films plotted as a function of the amount of a stress-reducing agent, 2-MBT, in the bath.

Figure 7.4.3.1.7. TEM micrographs and electron diffraction patterns showing the structure of electroless Ni– B films obtained from 4 baths containing NaBH4 as a reducing agent (see Table 7.4.3.1.1(a)).


663

Figure 7.4.3.1.8. A TEM micrograph for an electroless Ni–B film grown on a polycrystalline Fe substrate for 240 s from the 67 8C bath containing a small amount (200 mg/l) of NaBH4.

Figure 7.4.3.1.9. SEM micrographs showing the effect of the amount ((a) 0 and (b) 1.8 mg/l) of a stress-reducing agent, 2-MBT, on the surface morphology of electroless Ni–B films.

664

Nano-Plating

Figure 7.4.3.1.10. X-ray diffraction patterns from electroless Ni–B films containing various amounts of B obtained in the baths, which used DMAB as a reducing agent.

line films change to amorphous films with increasing B content. However, we do not know at which composition the film became amorphous. Figure 7.4.3.1.8 is a TEM micrograph for an electroless Ni – B film grown on a polycrystalline Cu substrate for 240 s from a 67 8C bath containing a small amount (200 mg/l) of NaBH4. The micrograph displays the structure of the Ni– B film grown on two differently oriented Cu grains. The grain boundary running from the upper left to the lower right appears to separate the two grains. Two different electron diffraction patterns from these grains are shown in Figure 7.4.3.1.8(a) and (b), which indicate that electroless Ni –B films tend to grow lattice-matched to the Cu substrate. Figure 7.4.3.1.9 is a SEM micrographs showing the effect of the concentration ((a) 0 and (b) 1.8 mg/l) of a stress-reducing agent, 2-MBT, on the surface morphology of


665

electroless Ni –B films. Without addition of 2-MBT, cracks are generated, but with addition of 2-MBT, no cracks appear, although nodules are present. Figure 7.4.3.1.10 is an X-ray diffraction pattern from electroless Ni– B films containing various concentrations of B. These films were obtained from the baths that used DMAB as a reducing agent. The B content in the film was varied by changing the concentration of DMAB (see Figure 7.4.3.1.2). Diffraction peaks representing Ni appeared from Ni – B films with low B content but became obscure with increasing B. The film finally became amorphous.

REFERENCES Suda, H., Watanabe, T., Misaki, Y. & Tanabe, Y. (1981a) J. Jpn Inst. Metals, 45, 5. Suda, H., Watanabe, T., Misaki, Y. & Tanabe, Y. (1981b) J. Jpn Inst. Metals, 45, 117. Tanabe, Y. & Watanabe, T. (1970) J. Metal Finish. Soc. Jpn, 21, 665. Tanabe, Y. & Watanabe, T. (1972a) J. Metal Finish. Soc. Jpn, 23, 38. Tanabe, Y. & Watanabe, T. (1972b) J. Metal Finish. Soc. Jpn, 23, 642. Watanabe, T. & Tanabe, Y. (1974) J. Metal Finish. Soc. Jpn, 25, 36. Watanabe, T. & Tanabe, Y. (1975a) J. Jpn Inst. Metals, 39, 1. Watanabe, T. & Tanabe, Y. (1975b) J. Jpn Inst. Metals, 39, 831. Watanabe, T. & Tanabe, Y. (1978) J. Jpn Inst. Metals, 42, 563.

7.4.3.2. Electroless Ni– P (Ito et al., 2001) 7.4.3.2.1. PLATING METHOD

An electroless Ni plating method uses sodium hypophosphate (NaHPO2·6H2O) as a reducing agent. Phosphorus atoms are generally incorporated into Ni films, forming Ni– P alloy deposits. The composition of the electroless Ni– P bath and the deposition conditions are listed in Table 7.4.3.2.1. The plating was carried out for 60 min in the 90 8C bath, which was agitated by magnetic stirrer. The bath pH was varied from 5.7 to 10 with the addition of NaOH. Polycrystalline Cu sheets were used as substrates. Since the electroless Ni –P deposition did not initiate spontaneously on the Cu substrate, a 5-s electrical bias was applied to start the deposition process electrolytically at a current density of 170 A/m2. The P content of the Ni– P films was determined using an energy dispersive X-ray spectrometer attached to the SEM machine.


Figure 7.4.3.2.1 depicts the P content of electroless Ni– P films plotted as a function of bath pH (which was changed with the addition of NaOH). With increasing pH, the P content decreased. Figure 7.4.3.2.2 is a graph of P content plotted against the NiSO4 concentration of the bath. With increasing NiSO4, the P content decreased. Various alloy compositions of electroless Ni –P films were prepared by changing the bath pH and concentration of NiSO4. 7.4.3.2.3 Structure Figure 7.4.3.2.3 shows X-ray diffraction patterns from different electroless Ni –P films. The films containing low P exhibited a sharp {111} diffraction peak of Ni crystals and

Table 7.4.3.2.1. The composition of electroless Ni–P plating baths and their plating conditions. Bath composition (mol/l) NiSO4·H2O NaHPO2·6H2O Citric acid (NH4)2SO4

Plating condition 0.1 0.2 0.5 0.5

Bath temperature Current density Plating time Substrates Bath pH Agitation

666

90 8C 170 A/m2 60 min Cu Plates 5.7–10 Magnetic stirrer

Electroless Ni – P

667

Figure 7.4.3.2.1. The P content of electroless Ni –P films plotted as a function of the bath pH.

thus consisted mostly of crystalline Ni. With increasing P content the {111} intensity decreased continuously. Here, as the Ni crystal size continued to be refined, the film structure finally reached an amorphous state. X-ray diffraction patterns alone, however, could not provide information on the exact alloy composition at which these Ni – P films transformed from crystalline to amorphous phases. The crystallization behavior of these Ni – P films can often offer information on the structure of the as-deposited films. For example, if the structure of the as-deposited film was uniformly amorphous, its crystallization should involve nucleation and grain growth processes. Alternatively, if the as-deposited film consisted of fine crystalline grains, these grains should grow without involving nucleation processes. Therefore, an inspection of the microstructure after heat treatment should tell us whether or not the original structure was crystalline or amorphous. Such a crystallization study can be conveniently done by an X-ray diffraction technique in combination with a heating experiment. Figure 7.4.3.2.4 shows a change in the X-ray diffraction pattern of an electroless Ni-19.8 at.% P film upon heating at various temperatures (100, 200, 300, and 400 8C).

668

Nano-Plating

Figure 7.4.3.2.2. Relationship between the P content of Ni–P films and the NiSO4 concentration of the bath.

The structure of the as-deposited film with a broad peak appears to be amorphous. With increasing temperature the intensity of the Ni {111} peak became stronger. From this observation, it can be envisioned that the as-deposited film already contained fine Ni crystals and these Ni crystals grew upon heating. Accordingly, it can be concluded that the electroless Ni-19.8 at.% P film was crystalline prior to the heat treatment. Figure 7.4.3.2.5 shows a change in the X-ray diffraction pattern of an electroless Ni21.7 at.% P film after heating at various temperatures (100, 200, 300, and 400 8C). Again, the structure of the as-deposited film appears to be amorphous. Similar to the Ni-19.8 at.% P film, the Ni {111} peak increased with increasing heating temperature. This film also contained fine Ni crystals and thus the as-deposited film was already crystalline. X-ray diffraction patterns in Figure 7.4.3.2.6 indicate structural changes in a Ni25.4 at.% P film as a result of heat treatment. A broad peak in the diffraction pattern of the as-deposited film did not change up to a temperature of 200 8C. Upon heating to 300 8C, Ni3P crystals appeared and its diffraction peaks became stronger with increasing temperature. In other words, the film maintained an amorphous state up to the 200 8C


669

Figure 7.4.3.2.3. X-ray diffraction patterns from electroless Ni– P films containing various P contents.

annealing, but Ni3P crystals nucleated upon heating at 300 8C. These Ni3P crystals subsequently grew with further heating at higher temperatures. From the above experiments, the critical P composition of as-deposited Ni– P films bounding the crystalline and amorphous phases lies between 21.7 and 24.5 at.% P and is most likely to be 23 at.% P. Figure 7.4.3.2.7(a) and (b) show the equilibrium phase diagrams of Ni – P alloys and the phases observed in electroless Ni – P films, respectively. It is interesting to note that the critical composition is close to the composition of the intermetallic compound (Ni3P).

670

Nano-Plating

Figure 7.4.3.2.4. A change in the X-ray diffraction patterns of an electroless Ni-19.8 at.% P film upon heating at different temperatures (100, 200, 300, and 400 8C).


671


672

Nano-Plating



673

Figure 7.4.3.2.7. (a) The equilibrium phase diagram of Ni– P alloys and (b) the phase diagram of electroless Ni– P films.

REFERENCE Ito, K., Wan, F. & Watanabe, T. (2001) J. Japan Inst. Metals, 65, 495.

7.4.4 Displacement Plating 7.4.4.1. Displacement Ag (Tanabe et al., 1971) 7.4.4.1.1. PLATING METHOD

The compositions of four displacement Ag plating baths are listed in Table 7.4.4.1.1. Polycrystalline Cu, single-crystal Cu and Fe sheets are used as substrates. The plating was conducted at ambient temperatures without solution agitation.


Figure 7.4.4.1.1 shows TEM micrographs and electron diffraction patterns for displacement Ag films obtained by immersing polycrystalline Cu in the (A) A and (B) B baths. After a 3-s immersion (cf. Figure 7.4.4.1.1(A) (a)), Ag particles (or islands) with a size of 5 , 30 nm appeared and their number density was 4:5 £ 1011 cm22. The number density of Ag islands obtained under all other deposition conditions was found to lie between 1 £ 1010 and 1 £ 1011 cm22. These islands were randomly oriented without lattice-matching to the substrate, as indicated by the formation of rings in the electron diffraction pattern. In other words, the islands did not develop strong chemical bonds to the substrate at this stage of growth. After a 7.5-s immersion (Figure 7.4.4.1.1(A) (b)), many islands with sizes of 5 , 30 nm are now linked together and at the same time their electron diffraction pattern contained a singlecrystal pattern superimposed on the rings. Some islands became lattice-matched to the substrate and some remained randomly oriented. After a 15-s immersion (Figure

Table 7.4.4.1.1. The composition of displacement Ag baths. Composition (g/l)

AgCN NaCN

Bath A

B

C

D

1 4

7.5 7.5

7.5 15.0

7.5 30.0

674

Displacement Ag

675

Figure 7.4.4.1.1. Plan-view TEM micrographs and their electron diffraction patterns displaying the growth processes of displacement Ag films formed on poly-crystalline Cu substrates from the A and B baths.

676

Nano-Plating

Figure 7.4.4.1.2. A (a) bright- and (b) dark-field TEM image pair showing the structure of a displacement Ag film obtained from the A bath after 120-s immersion.

7.4.4.1.1(A) (c)), the surface coverage by the Ag islands increased but did not fill the whole surface. The electron diffraction pattern was spotty, indicating an improved lattice-matching. After a 120-s immersion (Figure 7.4.4.1.1(A) (d)), the substrate surface was covered completely by the Ag film and the electron diffraction pattern became a single-crystal pattern. Since the B bath contains more Ag ions than the A bath, the film deposition rate from the B bath is much faster than that from the A bath. In fact, the substrate surface was completely covered even after a 0.3-s immersion (Figure 7.4.4.1.1(B) (a)). Figure 7.4.4.1.2 displays a (a) bright-field and (b) dark-field TEM image pair for a displacement Ag film obtained from the A bath after a 120-s immersion. The size of the sub-grains can be measured from these micrographs and was found to be 100 , 150 nm: Figure 7.4.4.1.3 is an electron diffraction pattern for displacement Ag films stripped from single-crystal (a) {001}, (b) {110}, (c) {111} Cu substrates after a 1-min immersion in the A bath. The Ag films appear to have grown epitaxially on the {001} and {110} Cu substrates, although the amount of lattice misfit between Ag and Cu is 13.0%. We found that lattice-matching with other orientations did not appear to reduce the amount of the misfit. Therefore, we consider that the following relationship prevailed the Ag-to-Cu lattice-matching.

Displacement Ag

677

Figure 7.4.4.1.3. Electron diffraction patterns of displacement Ag films grown on single-crystal (a) {001}, (b) {110}, (c) {111} Cu substrates after 1-min immersion in the A bath.

{001} Cu//{001} Ag and , 100 . Cu//, 100 . Ag {110} Cu//{110} Ag and , 110 . Cu//, 110 . Ag The Ag film grown on the {111} Cu substrate contained the , 111 . - and , 332 . oriented grains (see Figure 7.4.4.1.4). It can be concluded that a displacement Ag film generally grows on Cu in a lattice-matched mode. Figure 7.4.4.1.5(A) and (B) are SEM micrographs showing the surface morphologies of displacement Ag films obtained after a 24-h immersion in the B and C baths, respectively. The surface structure of the copper substrate appears to be replicated in the Ag film, as is

Figure 7.4.4.1.4. An indexed electron diffraction pattern showing how a ,331 . -oriented displacement Ag film lattice-matches on the {111} Cu substrate.

678 Nano-Plating

Figure 7.4.4.1.5. The surface morphologies of Ag films obtained by a displacement method after 24-h immersion in the B and C baths.

Displacement Ag

679

clear from the crystallographic patterns in (B) (a). It is interesting to note that the film grew thicker at the grain boundaries, indicating an enhanced growth rate at the grain boundaries except at the coherent twin boundaries. Another notable feature is the presence of numerous pits on the film surface (see (A)). It is possible that these pits were the sites where the copper substrate has dissolved. No such pits, however, are seen in (B). In this case, we cannot determine which part of the copper substrate dissolved.

REFERENCE Tanabe, Y., Asano, S. & Nakajima, H. (1971) J. Metal Finish. Soc. Jpn, 22, 461.

7.4.4.2. Displacement Au (Tanabe and Sudoh, 1968; Tanabe and Urao, 1968; Tanabe and Urao, 1968a,b) 7.4.4.2.1. PLATING METHOD

Various concentrations of gold chloride solutions were used to obtain displacement Au films. The effect of HCl addition was also studied. Single-crystal ({001} and {110}) and polycrystalline Fe sheets were used as the substrates after electropolishing in the Jacquet solution. Single-crystal ({001}, {110}, and {111}) Cu was also employed as the substrates after electropolishing in a solution containing 2 parts phosphoric acid and 1 part water.


7.4.4.2.2.1 Films grown on polycrystalline iron substrates Figure 7.4.4.2.1 depicts the initial stages of displacement Au films grown on Fe substrates in various concentrations of HAuCl4 baths. Each displacement Au film can be regarded as an assembly of fine granular Au particles. From a bath containing a low concentration of HAuCl4, the Au particles are sparsely distributed. With increasing HAuCl4 concentration, the surface coverage improved and the displacement rate increased. An electron diffraction pattern from the low-coverage film showed rings, suggesting that the Au particles are not lattice-matched to the substrate. With an increasing concentration of HAuCl4 in the bath, the electron diffraction pattern displayed spotty rings rather than continuous rings. Although some Au particles are still randomly oriented, some are lattice-matched to the substrate. Figure 7.4.4.2.2 depicts the initial stages of displacement Au films obtained by immersion for 1 and 5 s in a 0.05 wt% HAuCl4 bath containing various amounts of HCl. With an increasing concentration of HCl, the surface coverage improved and the displacement rate increased. Electron diffraction patterns from the low-coverage films showed rings, which indicate that the deposits are not lattice-matched to the substrate. With an increasing concentration of HCl, the deposits showed an increased tendency to lattice-match to the substrate as indicated by the formation of spotty rings. Figure 7.4.4.2.3 depicts the initial stages of displacement Au films obtained by immersion for 0.5, 1, and 30 s in a 0.1 wt% HAuCl4 bath containing 0.8% HCl. The initial morphology can be characterized as a tabular film coated with small dendrites. With increasing immersion time, these dendrite crystals grew larger until the whole surface was covered with needle-shaped crystals. 680

Displacement Au

681

Fig. 7.4.4.2.1. The structure of displacement Au films grown on Fe substrates from various concentrations of HAuCl4 baths.

682

Nano-Plating

Figure 7.4.4.2.2. The effect of HCl addition on the structure of displacement Au films grown on Fe substrates by immersing for 5 s in the 0.05 wt% HAuCl4 baths.

Displacement Au

683

Figure 7.4.4.2.3. The surface morphologies of displacement Au films grown on Fe substrates by immersing for 0.5, 1, and 30 s in the 0.1 wt% HAuCl4 baths containing 0.8 % HCl.

684

Nano-Plating

Figure 7.4.4.2.4. A time-dependent change in the growth morphology and electron diffraction pattern of a displacement Au film grown on polycrystalline Cu substrates from the 0.1 wt% HAuCl4 bath.

Displacement Au

685

Figure 7.4.4.2.5. TEM micrographs showing the growth morphology and electron diffraction patterns of a displacement Au immersed for 5 s in the bath containing 0.2 ml/l HCl. A dark thick line running from the top of the left side to the bottom of the right side is the grain boundary of the Cu substrate.

7.4.4.2.2.2 Films grown on pure copper substrates Figure 7.4.4.2.4 shows the time sequence of a displacement Au film grown on polycrystalline Cu substrates from the 0.1 wt% HAuCl4 bath. Similar to the case of the Fe substrates, granular Au crystals are distributed sparsely after a 1-s immersion and the corresponding electron diffraction pattern showed rings. After a 15-s immersion, the Au film covered the whole substrate surface. At the same time, single-crystal spots started appearing in the electron diffraction pattern, thus indicating that the Au grains started to lattice-match to the substrate. Figure 7.4.4.2.5 contains TEM micrographs showing the morphology and electron diffraction patterns of a displacement Au film immersed for 5 s in a bath containing 0.2 ml/ l HCl. A dark thick line running from the top of the left side to the bottom of the right side is the grain boundary of the Cu substrate. The growth morphology of the lower grain is clearly different from the upper grain and their diffraction patterns are indeed different from each other. The appearance of the dark thick line is an indication that more Au is deposited at the grain boundary.

REFERENCES Tanabe, Tanabe, Tanabe, Tanabe,

Y. & Sudoh, T. (1968) J. Metal Finish. Soc. Jpn, 19, 333. Y. & Urao, R. (1968) J. Metal Finish. Soc. Jpn, 19, 265. Y., Urao, R., Ogura, T. & Kurihara, T. (1968a) J. Metal Finish. Soc. Jpn, 19, 217. Y., Urao, R. & Ogura, T. (1968b) J. Metal Finish. Soc. Jpn, 19, 261.

7.4.4.3. Displacement Cd (Kanda, 1970) 7.4.4.3.1. PLATING METHOD

A displacement Cd plating bath is a 10 wt.% CdCl2 aqueous solution. In this study the solution also contained HCl. Polycrystalline Fe foils are used as the substrate.


Since it was difficult to peel displacement Cd films off the Fe substrate, we could not obtain reasonable TEM micrographs. For this reason, the deposition processes were studied using SEM. Figure 7.4.4.3.1(a) is an SEM micrograph showing a displacement Cd film grown on a polycrystalline Fe substrate after an 8-min immersion in a 10 wt.% CdCl2 aqueous solution, which was kept at 40 8C. Part of the surface image in Figure 7.4.4.3.1(a) is magnified in Figure 7.4.4.3.1(b). Here we note that the deposition rate depends on the orientation of the substrate grains. The surface after a 45-min immersion is shown in Figure 7.4.4.3.2. The surface morphology is clearly affected by the orientation of the underlying substrate. The surface morphology variation with the substrate grain orientation can be understood from the fact that the substrate dissolution rate depends on the orientation and the displacement Cd films tend to lattice-match to the substrate.

686

Displacement Cd

687

Figure 7.4.4.3.1. SEM micrographs showing a displacement Cd film grown on a polycrystalline Fe substrate after 8-min immersion in a 10 wt% CdCl2 aqueous solution.

688 Nano-Plating Figure 7.4.4.3.2. The surface morphology of a displacement Cd film on a polycrystalline Fe substrate after 45-min immersion in a 10 wt% CdCl2 aqueous solution.

Displacement Cd

689

REFERENCE Kanda, M. (1970) Chemical displacement plating of cadmium on pure iron substrates, PhD. dissertation, Tokyo Metropolitan University.

7.4.4.4. Displacement Cu (Tanabe, 1967) 7.4.4.4.1. PLATING METHOD

A displacement Cu bath was prepared by mixing 100 parts Cu(NH4)2Cl4·2H2O and 1.6 parts 37% HCl. The bath temperature during displacement reactions was kept at ambient temperature. Both single-crystal and polycrystalline Fe sheets were used as the substrates, which were electropolished in the Jacquet solution to a mirror finish.


The adhesion of a displacement Cu film to the substrate was good if the immersion time was limited to less than several seconds but deteriorated badly after a 3-min immersion. This deterioration was attributed to a dissolution of the Fe substrate during immersion processes. In other words, the interface between the displacement copper and the substrate dissolved, causing the peeling of the copper film from the substrate. A displacement copper film obtained after a 3-min immersion consisted of two layers. The layer in contact with the Fe substrate was a tabular film with a thickness of less than 1 mm, followed by a layer of

Figure 7.4.4.4.1. (a) A TEM micrograph and electron diffraction pattern for a tabular film of displacement Cu, which was obtained by immersing an Fe plate into a hydrochloric acid solution of Cu(NH4)2Cl4·2H2O for 30 s, and (b) its reflection electron diffraction pattern.

690

Displacement Cu

691

Figure 7.4.4.4.2. A TEM micrograph for a tabular film of displacement Cu, which was obtained by immersing an Fe plate into a hydrochloric acid solution of Cu(NH4)2Cl4·2H2O for 30 s, and its electron diffraction pattern.

1 –5 mm-long needle-shaped crystals. These needle-shaped crystals grew into resinoid crystals, with some long needles reaching as long as16 mm. In addition, the needle crystals were randomly oriented and did not appear to show any lattice-matching to the underlying layer. The thickness of the initial layer (free of the needle crystals) was less than 0.1 mm after a 0.1-s immersion. Figure 7.4.4.4.1 is a TEM micrograph for a tabular film of displacement Cu and its transmission/reflection electron diffraction patterns. From this study, we found that each needle-shaped crystal was a single crystal, on which fine (4 , 15 nm in diameter) Cu particles are attached (see Figure 7.4.4.4.2). Furthermore, these fine Cu particles were also found on the surface of the initial tabular layer. In both circumstances, the fine Cu particles were oriented randomly. The reflection electron diffraction pattern in Figure 7.4.4.4.1(b) indicates that the displacement Cu film has grown epitaxially on the single-crystal Fe substrate.

REFERENCE Tanabe, Y. & Ishibashi, K. (1967) J. Metal Finish. Soc. Jpn, 18, 41.

7.4.4.5. Displacement Zn (Nakata et al., 1997) 7.4.4.5.1. PLATING METHOD

The application of Zn coatings by a displacement technique is generally known as a zincating treatment. The composition of the displacement Zn bath is tabulated in Table 7.4.4.5.1. Polycrystalline and single-crystal ({001}, {110}, and {111}) Al plates were used as substrates.


Figure 7.4.4.5.1 comprises SEM micrographs showing the surfaces of single-crystal ({001}, {110}, and {111}) Al plates, which were immersed in the zincating bath for 15, 30, and 45 s. In all cases, small (, 1 mm) and large (, 10 mm) Zn particles are seen to have grown on the surfaces. The distribution of these particles is non-uniform on the {001} and {110} surfaces, but is uniform on the {111} surface. In the metal finishing industry, the zincating treatment is generally performed twice to provide a proper surface finish to Al surfaces. After the first 30-s zincating treatment, Zn particles grown on the surface are dissolved completely in nitric acid and then the second zincating is applied. Figure 7.4.4.5.2 shows the surfaces of single-crystal ({001}, {110},

Table 7.4.4.5.1. A displacement Zn plating bath (zincating bath). Composition

Concentration (g/l)

NaOH ZnO FeCl2·6H2O Additive

170 25 2 10

692

Displacement Zn

Figure 7.4.4.5.1. Zinc particles grown on the surfaces of single-crystal ({001}, {110}, and {111}) Al plates, which were immersed in the zincating bath for 15, 30, and 45 s.

693

694

Nano-Plating

Figure 7.4.4.5.2. The surfaces of single-crystal ({001}, {110}, and {111}) Al plates (a) after the 1st 30-second zincating and (b) after the 2nd zincating.

and {111}) Al plates (a) after the first zincating treatment and (b) after the second zincating treatment, given without dissolving Zn particles in nitric acid. In this instance, since the nitric acid step involving a dissolution of Zn particles was not performed, the surfaces became rougher. Figure 7.4.4.5.3 depicts the surface morphologies of twice-zincated single-crystal Al plates, which were given the nitric acid dip between the zincating treatments. Fine and coarse Zn particles are seen on the {001} and {110} surfaces, which became rough. Zinc particles formed on the {111} surface are uniformly fine, producing a very smooth surface. It is clear that the surface morphologies induced by zincating processes depend on the orientation of Al surface. Figure 7.4.4.5.4(a) depicts the surfaces of single-crystal ({001}, {110}, and {111}) Al plates after the first zincating treatment, followed by the nitric acid step. Protrusions are

Displacement Zn

695

Fig. 7.4.4.5.3. The surface morphologies of single-crystal ({001}, {110}, and {111}) Al plates (a) after the 1st 30-second zincating, which was followed by (b) the 2nd zincating treatment after Zn particles on the Al surfaces were dissolved completely in a nitric acid solution.

seen on all the surfaces and some crystallographic facets are likely to have remained after the dissolution in nitric acid. Large protrusions are present on the {001} and {110} surfaces, but small triangular protrusions were distributed uniformly over the {111} surface. The surface morphology of Zn deposits obtained by a zincating treatment was shown to depend on the orientation of Al substrates. This dependence can be attributed to a difference in the chemical property of metal surfaces. In Figure 7.4.4.5.4(b), we show the surfaces of single-crystal Al plates, which were etched in a HF solution. Square or triangular etch pits are seen on the {001} and {111} surfaces, and appear to either line up along the low-angle grain boundaries, or to decorate isolated dislocations inside the grains. These etching morphologies are clearly different from that produced when nitric acid etching is performed between the two zincatings.

696

Nano-Plating

Fig. 7.4.4.5.4. (a) The surfaces of single-crystal ({001}, {110}, and {111}) Al plates after the 1st zincating treatment, followed by the nitric acid step and (b) the surfaces of single-crystal Al plates, which were etched in an HF solution.

REFERENCE Nakata, T., Wada, I., Imai, H., Ikejima, K., Inoue, K. & Watanabe, T. (1997) J. Surf. Finish. Soc. Jpn, 48, 820.

FURTHER READING Tashiro, T., Chiba, K., Fukuda, Y., Nakao, H. & Honma, H. (1994) J. Surf. Finish. Soc. Jpn, 45, 720.

Index Adatom 7, 9, 12 – 15, 32, 97, 101, 102, 123, 134, 263 Anchoring force 129 Bridgeman technique 212 Crystallographic facet 32, 211, 376, 695

Nodule 73, 83, 84, 126, 436, 509, 587, 627, 665 Non-equilibrium phase 484

Damascene 204 Deformation texture 209 Dendrite 14, 37, 38, 101, 104, 147, 169, 229, 258, 403, 499, 680 Diffusion layer 9, 15, 97

Overpotential theory (OT) 8, 118, 161 Photolithography 203, 204 Pseudomorphism 70, 79 Replica method 203, 204, 246, 247

Electrical double layer 97 Embryo 191, 192 Epitaxy 68, 199, 200, 223, 335, 494 Equilibrium phase diagram 3, 10, 12, 14, 43, 106– 113, 212, 216, 257, 384– 391, 400– 440, 446–470, 478–486, 507–558, 571–631, 634–656, 669, 673

Stranski-Krastanov (S-K) 97, 99, 101 Stress-reducing agent 657, 661– 664 Strike Au plating 274 Substitutional type solid-solution 109 Supercooled solid 14 Superlattice 106, 199, 541 Supersaturated solid-solution 2, 13, 113, 432, 463, 529, 534, 538, 541, 570, 632, 646 Surface deformation zone (SDZ) 210 Surface diffusion 8, 12, 15, 32, 119, 120

Focused ion-beam (FIB) 221, 235 Frank-van der Merwe type 97, 99, 105, 106 Helmholz 7, 97 Initial layer 80, 82, 399, 427, 437, 438, 446, 459, 525, 620, 622, 639, 691 Initiation treatment 124 Interfacial energy 70, 99 Interstitial type solid-solution 10, 109 Laves layer 210 Leveling agent 6, 38, 533 LIGA (Lithographie Galvanoforumung Abformung) 204 Metal-ion deficient layer (MIDL) 124 Metal-ion denuded layer (MIDL) 14, 15, 38

Metallizing 128 Meta-stable phase 109– 114, 202, 216, 428– 434, 449, 463, 534– 542, 570– 574, 592, 631– 640 Misfit dislocation 70, 77, 79, 100, 223 Misfit twin crystal 61, 70, 78, 79

TTT (time-temperature-transformation) 192 Type 1 additive 38 Type 2 additive 38, 39 Uni-directional solidification 212 Vegerds law 109, 432, 438, 470, 486, 518, 520, 529, 559, 570 Vomer-Waber (V-W) type 97, 99, 104– 106 Watts bath 32, 34, 53 – 55, 74 –77, 87, 344– 356, 599, 611 Willson Equation 169, 376, 580 Zincate 132, 694

697


Nano Plating - Microstructure Formation Theory of Plated Films and a Database of Plated Films

Microstructure evaluation of thin films during growth

International Dictionary of Films and Filmmakers. - Films

Analysis and design of plated structures

Art of Watching Films

Handbook of Plastic Films

Preparation of Thin Films

International Dictionary of Films and Filmmakers. Volume 1. Films

Waveguide Spectroscopy of Thin Films, Volume 33 (Thin Films and Nanostructures) (Thin Films and Nanostructures)

Crime Films

Analysis and Design of Plated Structures: Volume 2: Dynamics

London Films

The Films of Vincent Price

The Films of Jack Nicholson

Thin films of soft matter

Ceramic Films and Coatings

The Films of Federico Fellini

The Art of Watching Films

Adhesion Aspects Of Thin Films

Materials Science of Thin Films

Polymolecular and Monomolecular Films

The Art of Watching Films

Adhesion Aspects of Thin Films

The Films of Robert Wise

Savory Sweets : From Ingredients to Plated Desserts

foam and foam films. Theory, Ecxperiment, Application

Handbook of Industrial Diamonds and Diamond Films

International Dictionary of Films and Filmmakers

International Dictionary of Films and Filmmakers

Principles of Vapor Deposition of Thin Films

Handbook of Infrared Spectroscopy of Ultrathin Films

Nano Plating - Microstructure Formation Theory of Plated Films and a Database of Plated Films

Microstructure evaluation of thin films during growth

International Dictionary of Films and Filmmakers. - Films

Analysis and design of plated structures

Art of Watching Films

Handbook of Plastic Films

Preparation of Thin Films

International Dictionary of Films and Filmmakers. Volume 1. Films

Waveguide Spectroscopy of Thin Films, Volume 33 (Thin Films and Nanostructures) (Thin Films and Nanostructures)

Crime Films

Analysis and Design of Plated Structures: Volume 2: Dynamics

London Films

The Films of Vincent Price

The Films of Jack Nicholson

Thin films of soft matter

Ceramic Films and Coatings

The Films of Federico Fellini

The Art of Watching Films

Adhesion Aspects Of Thin Films

Materials Science of Thin Films

Polymolecular and Monomolecular Films

The Art of Watching Films

Adhesion Aspects of Thin Films

The Films of Robert Wise

Savory Sweets : From Ingredients to Plated Desserts

foam and foam films. Theory, Ecxperiment, Application

Handbook of Industrial Diamonds and Diamond Films

International Dictionary of Films and Filmmakers

International Dictionary of Films and Filmmakers

Principles of Vapor Deposition of Thin Films

Handbook of Infrared Spectroscopy of Ultrathin Films

Recommend Documents