DATE DUE

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SEP 1 ( 1986

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4W nil 4 r JUL 1 5 1992

C AVLOKO rRINTEOINU.S A.


UNITED STATES DEPARTMENT OF COMMERCE • Peter G. Peterson, Secretary

iq Science of Ceramic Machining

c and Surface Finishing

Proceedings of a Symposium Sponsored by

the American Ceramic Society, the Office

of Naval Research, and the National

Bureau of Standards, November 2-4, 1970,

held at NBS, Gaithersburg, Maryland

Edited by

S. J. Schneider, Jr.

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

and

R. W. Rice

Naval Research Laboratory

Washington, D. C. 20390

Associate Editors

C. F. Bersch, Naval Air Systems Command,

Washington, D. C. 20360; A. M. Diness,

Office of Naval Research, Arlington, Virginia

and J. B. Wachtman, Jr., Institute for Materials Research, National Bureau of Standards,

Washington, D. C. 20234

US. National Bureau of Standards Special Publication 348

4 I <

Nat. Bur. Stand (U.S.), Spec. Publ. 348, 431 pages (May 1972)

CODEN: XNBSAV

Issued May 1972 -

For sale by the Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402—Price S5.25 (Order by SD Catalog No. C13.10:348).
Abstract

This volume presents the proceedings of the Symposium on the Science of Ceramic

Machining and Surface Finishing held at the National Bureau of Standards in Gaithersburg, Maryland on November 2-4, 1970. The symposium was jointly sponsored by

the Baltimore-Washington Section of the American Ceramic Society, the Office of

Naval Research, and the National Bureau of Standards. The purpose of the conference was to survey the developing science of ceramic machining and to stimulate further

progress by discussion of current problems and research programs. In addition to a

panel discussion, 37 review and original research papers were presented with attention

focused on four main areas: (1) Techniques and Mechanisms of Removal and Shaping

(2) Surface Treatment (3) Analysis and Characterization of Machining Effects and

(4) Mechanical and Other Effects of Finishing. An edited version of the floor discussion following each paper is given.

Key words: Ceramics; ceramic machining; mechanical effects of machining; removal

and shaping of ceramics; surface treatment.

Library of Congress Catalog Card Number: 79-611323
Foreword

One way in which the National Bureau of Standards attempts to achieve its goal—^to strengthen and advance the nation's science and technology and to facilitate their effective application for public benefit—is by insuring the availability of basic data on the

properties of materials. In this effort the Institute for Materials Research holds symposia that provide a forum for discussion of the latest research results, and that focus

attention on areas of material science in which basic knowledge is lacking or needs improvement.

The latest of these symposia was on the Science of Ceramic Machining and Surface

Finishing. Ceramics have long been used in high temperature and structural applications and yet have not achieved the full measure of reliability demanded by modern technology. Therefore, the entire spectrum of ceramic processing and in particular the ceramic

machining segment must be elevated to new scientific levels. In recognition of this need

NBS, the Office of Naval Research, and The American Ceramic Society organized this

symposium so that the groundwork for a science of ceramic machining could be established.

The National Bureau of Standards is publishing these Proceedings not only to make

an important body of knowledge available, but to stimulate further research efforts in the

field as well.

Lewis M. Branscomb

Director

• •• III
Preface

Presented here are the proceedings of the symposium on the Science of Ceramic Machining and Surface Finishing held at the National Bureau of Standards at Gaithersburg, Maryland on November 2-4, 1970. Approximately 190 scientists, engineers, and

others versed in ceramic machining attended the three-day meeting. Overall, seven

countries, including the USA, w^ere represented.

The machining and finishing process is of fundamental importance to the manufacture of many commercial ceramic products. Ultimate success and reliability often lies

with the nature of machining variables. Previously this facet of production was relegated to state-of-art procedures generally derived from knowledge attained in the processing of metals. Of late, however, many in the ceramic sciences have become cognizant

that a better understanding of machining and all its ramifications must be gained if

ceramics are to realize their full potential. Increased awareness of this is evidenced by

the recent report of the Materials Advisory Board on Ceramic Processing (Ad Hoc

Committee on Ceramic Processing, J. A. Pask, Chairman, National Academy of Sciences

Publication 1576 (1968) ). The Panel on Finishing surveyed various techniques, mechanisms, and resultant surface character and put forth recommendations for further

study. Significantly this symposium is a perfect adjunct to the work of the Panel.

The symposium had as its stated purpose "To survey the developing science of ce- ramic machining and to stimulate further progress by discussion of current problems and

research." These objectives were met through the presentation of a variety of invited

summary papers, supplemented with contributions on original research. Accordingly

the scope of the conference broadly encompassed the techniques, analysis, and effects of

machining and surface finishing by several basic techniques which include mechanical,

energy beam, and chemical methods. The known effects of such operations on ceramic

properties, especially mechanical properties, were covered with emphasis being placed

on the fundamental nature of these effects. Analysis and characterization of resultant sur- faces and subsurfaces are also covered since this information is necessary to an understanding of removal mechanisms and their effects on properties. The symposium fittingly

concluded with a panel discussion which effectively focused attention on significant

trends and problem areas. The symposium and the proceedings admittedly are not the last word in the complex

subject of machining. It is hoped, however, that the papers published here can lead to a

better understanding of material removal mechanisms and resultant surface structures

and properties, and thus to an eventual improvement in the technology of ceramic finishing. Furthermore, the symposium can serve as an aid in the development of a true

Science of Ceramic Machining and Surface Finishing.

S. J. Schneider, Jr. and R. W. Rice

Editors

Acknowledgment

The editors wish to thank all those who participated in the symposium either as attendees, speakers, session chairmen, or panel members. Special thanks are due the following for their assistance in organizing the meeting and for publication of the proceedings:

R. T. Cook, National Bureau of Standards

B, L. Oberholtz:er, National Bureau of Standards

K. M. Smith, Applied Physics Laboratory of Johns Hopkins University.

iv
Contents

Page

FOREWORD iii PREFACE iv

PARTICIPANTS viii WELCOME 1 E. Horowitz

I. INTRODUCTION

Session Chairman

S. J. Schneider, National Bureau of Standards

Science and the Machining and Surface Finishing of Ceramics 1 R. W. Rice, C. Bersch, and A. M. Diness

II. TECHNIQUES AND MECHANISMS OF REMOVAL AND SHAPING

Session Chairmen

R. W. Rice, Naval Research Laboratory

V. E. Wolkodoff, Coors Porcelain Company

A. H. Heuer, Case Western Reserve University

A. Mechanical Methods

Mechanical Methods of Ceramic Finishing 5

P. J. Gielisse and J. Stanislao

Discussion 35

Direct Observation of Material Removal Process During Grinding of Ceramics

by Micro-flash Technique 37

O. Imanaka, S. Fujino, and S. Mineta

Grinding—A Means of Shaping and Sizing Ceramics 45

J. A. Mueller

Discussion 52

Aspects of Machining Glass-Ceramic Materials 53

G. H. Allgeyer and L. V. Colwell

Discussion 57

The Princioles of Grinding 59

R. S. Hahn and R. P. Lindsay

Discussion 71

A Basic Study of the Diamond Grinding of Alumina 78

D. M. Busch and J. F. Prins

Discussion 87

On the Strength of Ceramics as a Function of Microstructure, Grinding Parameters, Surface Finish and Environmental Conditions 89

R. Sedlacek, F. A. Halden, and P. J. Jorgensen

Discussion 96

Grinding Alumina with Diamond Abrasives 99

R. J. Caveney and N. W. Thiel

Discussion 112

Ceramic Substrate and Specimen Fabrication 113

H. C. Leistner and W. A. Wilson

On the Shaping of Brittle Solids by Erosion and Ultrasonic Cutting 119

H. L. Oh. K. P. L. Oh, S. Vaidyanathan, and I. Finnic

Discussion 132

Sonic Machining of Ceramics 133

W. B. Campbell

Discussion 139

Environment-Sensitive Machining Behavior of Nonmetals 141

A. R. C. Westwood and R. M. Latanision

Discussion 153

V
B. Nonmechanical Methods

Page

Shaping or Figuring Ceramic Surfaces by Ion-Beam Bombardment 155

P. W. Levy

Discussion 168

The Effect of Sputtering on Surface Topography and Strength of Ceramics 171

R. W. Rice

Discussion 187

Computer Controlled Ionic Polishing of Optical Surfaces 189

J. W. Douglass

Arc, Laser, and Electron Beam Machining of Ceramics 193

R. W. Rice

The Techniques and Mechanisms of Chemical, Electrochemical and Electrical

Discharge Machining of Ceramic Materials 197

D. W. Lee and G. Feick

Discussion 211

III. TECHNIQUES AND MECHANISMS OF SURFACE FINISHING

Session Chairman

P. J. Gielisse, University of Rhode Island

Improvements in the Surface Finish of Ceramics by Flame Polishing and Annealing Techniques 213

M. J. Noone and A. H. Heuer

Discussion .' 232

Healing of Surface Cracks in Ceramics 233

F. F. Lange

Discussion 236

Flame Polishing of Flat Oxide Bars 237

P. F. Becher and R. W. Rice

Discussion 244

Continuous Flame Polishing of Sapphire Filament 247

J. T. A. Pollock

Discussion 256

Preparation of Smooth, Crystalline, Damage Free, Sapphire Surfaces by Gaseous

Etching 259

W. A. Schmidt and J. E. Davey

Discussion 265

The Strength of Gas Polished Sapphire and Rutile 267

R. W. Rice, P. F. Becher, and W. A. Schmidt

Discussion 269

IV. SURFACE AND SUBSURFACE ANALYSIS AND CHARACTERIZATION

Session Chairman

W. H. Rhodes, AVCO Corporation

Analysis and Characterization of Ceramic Surfaces for Electronic Applications .... 271

R. C. Sundahl and L. Berrin

Discussion 290

Surface Characteristics of Ceramic Substrates for Hybrid and Microwave Electronic Circuits 293

J. K. Emery

Ceramic Surface Texture by Reflective Replica Technique 301

W. C. Lo

Quantitative Surface Finish Characterization by CESEMI 309

E. W. White, H. A. McKinstry, and A. Diness

An Assessment of Surface and Subsurface Damage Introduced in Ceramics by

Semifinish Grinding Operations 317

B. G. Koepke

Discussion 332

Observations on Mechanically Abraded Aluminum Oxide Crystals by Transmission Electron Microscopy 333

B. J. Hockey

Discussion 339

Acoustic Emission Monitoring of Surface-Damaged Ceramic Materials 341

L. R. Bunnell, J. C. Crowe, and P. E. Hart

vi
V. MECHANICAL AND OTHER EFFECTS OF FINISHING

Session Chairmen

W. H. Rhodes, AVCO Corporation

J. B. Wachtman, Jr., National Bureau of Standards

Page

Effects of Surface Finishing on Mechanical and Other Physical Properties of

Ceramics 343

R. J. Stokes

Discussion 352

Strength Effects Resulting from Simple Surface Treatments 353

H. P. Kirchner, R. M. Gruver, and R. E. Walker

Discussion 363

The Effect of Grinding Direction on the Strength of Ceramics 365

R. W. Rice

Discussion 376

The Influences of Material Removal on the Strength and Surface of an Alumina . . 377

H. S. Starrett

Discussion 389

The Effect of Grinding Variables on the Strength and Surface Finish of Alumina . . 391

R. Sedlacek, S. A. Halden, and P. J. Jorgensen

Discussion 397

Edge Effect on the Modulus of Rupture of Ceramic Substrates 399

W. C. Lo

Discussion 400

VI. PANEL DISCUSSION

J. B. Wachtman, Jr., Moderator

F. Ernsberger, P. J. Gielisse, A. H. Heuer, W. H. Rhodes, R. W. Rice

Panel Members

Discussions 401

Author Index 417

Subject Index 419

VIl
Participants

Samuel J. Acquaviva

Army Materials & Mechanics

Research Center

Watertown, Massachusetts 02172

Jon J. Aho

Bausch & Lomb

635 St. Paul Street

Rochester, New York 14602

Guy H. Allgeyer

Owens-Illinois

1700 North Westwood

Toledo, Ohio 43607

Jan Anpelo

D. M. Stewart Manufacturing Co.

Post Office Box 510

Chattanooga, Tennessee 37701

F. Gary Augeri

MIT /Draper Lab

275 Massachusetts Avenue

Cambridge, Massachusetts 02139

Frank P. Bailey

Chemical Research Laboratories

C.S.LR.O.

Post Office Box 4331

Melbourne, Victoria, Australia

Jimmy D. Bailey

American Lava Corporation

Laurens, South Carolina 29360

Spencer H. Baker

Norton Company

New Bond Street

Worcester, Massachusetts 01606

Norris H. Barbre

Can-Tex Industries

Division of Harsco Corporation

Post Office Box 254

Weatherford, Texas 76086

Y. Baskin

Ferro Corporation

Technical Center

4150 East 56th Street

Cleveland, Ohio 44105

C. H. Bates

Bell Aerospace Company

Post Office Box 1 Buffalo, New York 14240

P. F. Becher

Naval Research Laboratory

Code 6136

Washington, D.C. 20390

Richard Behrens

Dow Coming Corporation

Saginaw Road

Midland, Michigan 48640

L. Berrin

Bell Telephone Laboratories

555 Union Boulevard

Allentown, Pennsylvania 18103

C. Bersch

Naval Air Systems Command

Jefferson Plaza, Room 1000

Arlington, Virginia 22901

R. E. Bickelhaupt

Southern Research Institute

2000 Ninth Avenue South

Birmingham, Alabama 35205

Joe Blaze

McDanel Refractory Porcelain Co.

510 Ninth Avenue

Post Office Box 560

Beaver Falls, Penna. 15010

Sherman D. Brown

University of Illinois

Dent. Ceramic Engineering

Urbana, Illinois 61801

W. R. Brown

Corning Glass Works

Fluidic Products Department

Corning, New York 14830

D. M. Busch

Diamond Research Laboratory

Industrial Division

Post Office Box 916

Johannesburg, South Africa

D. J. Cameron

Atomic Energy of Canada Ltd.

Whiteshell Nuclear Research

Pinawa, Manitoba, Canada

William B. Campbell

The Ohio State University

2041 North College Road

Columbus, Ohio 43210

John J. Capellman

PPG Industries

Creighton, Pennsylvania 15030

R. J. Caveney

Diamond Research Laboratory

Post Office Box 916

Johannesburg, South Africa

David W. Chalkley

Trans-Tech, Inc. 12 Meem Avenue

Gaithersburg, Maryland 20760

Jaime W. Colon

Potter Instrument Co.

Plainview, New York 11803

L. V. Colwell

University of Michigan

2046 East Engineering Bldg.

Ann Arbor, Michigan 48104

Donald G. Cooper

Avco Bay State Abrasives Division

Union Street

Westboro, Massachusetts 01581

Wade Cooper

McDanel Refractory Porcelain Co.

510 Ninth Ave., P. 0. Box 560

Beaver Falls, Penna. 15010

Kristofer T. Dahl, 2Lt. USAF

SMA.MA (MAQQLC)

McClellan Air Force Base

California 95652

A. U. Daniels

Air Force Machineability

Data Center

3980 Rosslyn Drive

Cincinnati, Ohio 45209

William H. Daniels

Westinghouse Research Lab

Beulah Road

Pittsburgh, Penna. 15235

N. S. Dempster

Embar,sy of Australia

1601 Massachusetts Avenue, N.W.

Washington, D.C. 20036

Frank V. DiMarcello

Bell Laboratories, Inc. Mountain Avenue

Murray Hill, New Jersey 07974

A. M. Diness

Office of Naval Research

Room 619, Ballstron #1

Arlington, Virginia 22217

Sam DiVita

U.S. Army Electronics Command

AMSEL-KL-EC

Fort Monmouth, New Jersey 07703

Louis P. Domingues

Trans-Tech, Inc. 12 Meem Avenue

Gaithersburg, Maryland 20760

David J. Donovan

Data Magnetics

355 Maple Avenue

Torrance, California 90503

S. Doroff

Office of Naval Research

Washington, D.C. 20207

James W. Douglass

Perkin-Elmer

Norwalk, Connecticut 06553

Bruce W. Dunnington

The Glennel Corporation

908 Chickadee Lane

West Chester, Pennsylvania 19380

Richard A. Edsall

Trans-Tech, Inc. 12 Meem Avenue

Gaithersburg, Maryland 20760

A.. E. Edwards

Material Dynamics

3151 Fostoria Way

San Ramon, California 94583

Shane A. Ellis

Ferroxcube Corporation

Saugerties, New York 12477

J. K. Emery

Accumet Engineering Corp.

25 Broad Street

Hudson, Massachusetts 01749

Fred Ernsberger

PPG Industries

Box 11472

Pittsburgh, Pennsylvania 15238

James W. Evans

Corning Glass Works

Process Research-Sullivan Park

Corning, New York 14830

George Eyerly

D. M. Stewart Manufacturing Co. Post Office Box 510

Chattanooga, Tennessee 37401

viii
Iain Finnic

University of California

Berkeley, California 94720

Ross Firestone

Case Western Reserve University

White Building

University Circle

Cleveland, Ohio 44118

P. L. Fleischner

National Beryllia

Haskell, New Jersey 07420

R. K. Frazer

Applied Physics Lab

Johns Hopkins University

Scaggsville, Maryland 20910

Nicholas Friedman

Duramic Products, Inc. 426 Commercial Avenue

Palisades Park, New Jersey 07650

Frank S. Gardner

Office of Naval Research

495 Summer Street

Boston, Massachusetts 02210

Erberto Francisco Gentile

Instituto de Energie Atomica

C. P. 11049

S. Paulo, Brazil

Peter J. Geilisse

University of Rhode Island

Kingston, Rhode Island 02881

William F. Gilbert, Jr. Bell Telephone Laboratories, Inc. 555 Union Boulevard

Allentown, Pennsylvania 18103

A. J. Gitter

Engineered Materials Division

Hi-Purity Materials, Inc. Box 363, Church Street Station

New York, New York 10008

Harry J. Green, III IBM-Components Division

Route 52, Bldg. 320-33, D/281

Hopewell Junction, New York 12533

Bernard A. Greene

Griffin Pipe Products Co.

Milledgeville, Georgia 31061

Victor A. Greene

Automated Business Systems

550 Central Avenue

Orange, New Jersey 07051

Irwin G. Greenfield

Mechanical & Aerospace Engr. Dept.

University of Delaware

Newark, Delaware 19711

William T. Greenfield

Coming Glass Works

Fluidic Products Department

Corning, New York 14830

William R. Griffin

Raytheon Company

130 Second Avenue

Waltham, Massachusetts 02154

Nelson Grimm

General Electric Company

24400 Highland Road

Richmond Heights, Ohio 44143

William E. Gurwell

Easton, Yale & Towne, Inc. 26201 Northwestern Highway

Southfield, Michigan 48075

Robert S. Hahn

Heald Division-Cincinnati Milacron

New Bond Street

Worcester, Massachusetts 01606

David M. Haines

INSACO, Inc. Post Office Box 422

Quakertown, Pennsylvania 18951

Joe N. Harris

Georgia Institute of Technology

High Temperature Materials Div.

Atlanta, Georgia 30332

Pat Hart

Battelle Northwest

Post Office Box 999

Richland, Washington 99352

D. P. H. Hasselman

Lehigh University

Bethlehem, Pennsylvania 18015

Howard Herzig

NASA

Goddard Space Flight Center

Code 284.1

Greenbelt, Maryland 20771

A. H. Heuer

Case Western Reserve University

Cleveland, Ohio 44106

Terrence P. Hobin

British Embassy

3100 Massachusetts Avenue

Washington, D.C. 20008

B. J. Hockey

National Bureau of Standards

Room A359, Building 223

Washington, D.C. 20234

M. W. Hoelscher

RCA Incorporated

New Holland Pike

Lancaster, Pennsylvania 17604

E. Horowitz

Institute for Materials Research

National Bureau of Standards

Washington, D.C. 20234

J. W. Hunt

Duplate Canada Limited

Ceramics Division

First Avenue

Oshawa, Ontario, Canada

Robert H. Insley

Champion Spark Plug Company

20000 Conner Avenue

Detroit, Michigan 48234

Virgil Irick, Jr. American Lava Corporation

Cherokee Boulevard

Chattanooga, Tennessee 37405

Lonnie F. Ivey

American Standard

Post Office Box A

Paintsville, Kentucky 41240

R. Jackel

Office of Naval Research

Washington, D.C. 20207

Paul J. Jorgensen

Stanford Research Institute

333 Ravenswood Avenue

Menlo Park, California 94025

Paul G. Juenemann

Trans-Tech, Inc. 12 Meem Avenue

Gaithersburg, Maryland 20760

Erwin E. Kipfer

Pfaudler Company

1000 West A.venue

Rochester, New York 14603

Henry P. Kirchner

Ceramic Finishing Company

Post Office Box 498

State College, Pennsylvania 16801

R. J. Kleinwassink

Philips Research Labs

Eindhoven, Netherlands

B. G. Koepke

Honeywell Corporate Research

500 Washington Avenue, S. Hopkins, Minnesota 55343

J. Koorneef

Philips Research Labs

Eindhoven, Netherlands

William Kovacs

Atomic Energy of Canada Ltd.

Whiteshell Nuclear Research

Pinawa, Manitoba, Canada

Don A. Kubose

Naval Ordnance Laboratory

White Oak, Maryland

Paul L. Kuzmick, Jr. Paul L. Kuzmick Company

271 Grove Avenue

Verona, New Jersey 07044

Richard Lane

Hamco

1000 Millstead Way

Rochester, New York 14624

Fred F. Lange

Westinghouse R&D

Churchill Boro

Pittsburgh, Pennsylvania 15235

Robert E. Lasson

1MB Corporation

2800 Sand Hill Road

Building 031, Dept. 201

Menlo Park, California 94025

D. W. Lee

Arthur D. Little Inc. Acorn Park

Cambridge, Massachusetts 02140

Henry C. Leistner

National Bureau of Standards

325 South Broadway

Boulder, Colorado 80302

James D. Lester

U.S. Army Frankford Arsenal

Philadelphia, Pennsylvania 19137

Newton Levy, Jr. W. R. Grace & Company

Clarksville, Maryland 21029

ix
p. W. Levy

Physics Department

Brookhaven National Laboratory

Upton, New York 11973

Jack Liker

Basic Ceramics Inc. 221 7th Avenue

Hawthorne, New Jersey 07507

Alan Little

Duplate Canada Ltd.

50 St. Clair Avenue, West

Toronto, Ontario, Canada

W. C. Lo

Bell Telephone Laboratories, Inc. 555 Union Boulevard

Allentown, Pennsylvania 18103

William D. Long

Kaman Sciences

1700 Garden of the Gods Road

Colorado Springs, Colorado 80907

Evan C. Luce

Norton Company

1 New Bond Street

Worcester, Massachusetts 01606

Howard L. McCollister

Owens-Illinois Inc. 1700 North Westwood

Toledo, Ohio 43607

J. F. McMahon

13 Sayles Street

Alfred, New York 14802

B. A. MacDonald

Office of Naval Research

Washington, D.C. 20207

William McDonough

Naval Research Laboratory

Washington, D.C. 20390

Karl R. McKinney

Naval Research Laboratory

Washington, D.C. 20390

H. A. McKinstry

Pennsylvania State University

161 Engineering Science Bldg.

University Park, Penna. 16802

I. Machlin

Office of Naval Research

Washington, D.C. 20207

Donald J. Maclntyre

Basic Ceramics Inc. Box 2090

Hendersonville, N. C. 28739

Frank Magnanelli

Trans-Tech, Inc. 12 Meem Avenue

Gaithersburg, Maryland 20760

Alan Maietta

Paul L. Kuzmick Company

271 Grove Avenue

Verona, New Jersey 07044

Joseph A. Martis

Metallurgy Product Department

General Electric Company

11177 East Eight Mile Road

Detroit, Michigan 48232

John Matsko

Division 5332

Sandia Laboratories

Albuquerque, New Mexico 87115

George Mayer

U.S. Army Research Office

Box CM, Duke Station

Durham, North Carolina 27706

George L. Metzger

James H. Rhodes & Company

1026 West Jackson Boulevard

Chicago, Illinois 60607

H. C. Miller

Super-Cut, Inc. 3418 North Knox

Chicago, Illinois 60641

Herbert R. Moeller

Union Carbide Corporation

Post Office Box 985

Union, New Jersey 07083

J. A. Mueller

The Carborundum Company

Post Office Box 403

Niagara Falls, New York 14302

M. K. Murthy

Ontario Research Foundation

Sheridan Park

Ontario, Canada

Donald R. Ness

Western Gold & Platinum Co.

555 Harbor Boulevard

Belmont, California 94002

M. J. Noone

General Electric Company

Space Sciences Laboratory

Room M-9126, P. O. Box 8555

Philadelphia, Penna. 19101

Theodore Nussbaum

Burroughs

Box 472

Piscataway, New Jersey 09854

Martin J. O'Hara

Battelle Memorial Institute

505 King Avenue

Columbus, Ohio 43201

William J. Otto

Ferroxcube Corporation

Mt. Marian Road

Saugerties, New York 12477

Hayne Palmour, III North Carolina State University

Department of Engineering Resea

Raleigh, North Carolina 27607

W. H. Parker

Kaiser Refractories

Post Office Box 870

Pleasanton, California 94566

William H. Parrish

Material Dynamics

Physics International Company

3151 Fostoria Way

San Ramon, California 94583

Edwin A. Pascoe

General Electric Company

6325 Huntley

Worthington, Ohio 43085

Damon S. Paulley

Coors Porcelain Company

600 9th Street

Golden, Colorado 80401

D. J. Perduyn

N. V. Philips

Eindhoven, Netherlands

H. A. Perry

Naval Ordnance Laboratory

Code 2301

Silver Spring, Maryland 20910

Raymond Pietras

Howmet Corporation

Roy Street

Dover, New Jersey 07801

J. T. A. Pollock

Tyco Laboratories, Inc. 16 Hickory Drive

Waltham, Massachusetts 02154

Kenneth Prince

Norton Company

1 New Bond Street

Worcester, Massachusetts 01606

W. G. Ranch

Office of Naval Research

Washington, D.C. 20207

Robert F. Rea

Champion Spark Plug Company

20000 Conner Avenue

Detroit, Michigan 48234

William H. Rhodes

AVCO Corporation

Lowell Industrial Park

Lowell, Massachusetts 01851

R. Rice

Naval Research Laboratory

Code 6136

Washington, D.C. 20390

John Robertson

American Standard Inc. 834 East Broadway

Louisville, Kentucky 40204

Kenneth S. Rolley

Hamco Division Kayex Corp.

1000 Millstead Way

Rochester, New York 14624

George W. Roust

Lawrence Radiation Laboratory

Post Office Box 808

Livermore, California 94551

W. A. Schmidt

Naval Research Laboratory

Code 5210

4555 Overlook Avenue, S.W.

Washington, D.C. 20390

S. J. Schneider

B214, Materials Building

National Bureau of Standards

Washington, D.C. 20234

Rudolf Sedlacek

Stanford Research Institute

333 Ravenswood Avenue

Menlo Park, California 94025

D. J. Shanefield

Western Electric Company

Post Office Box 900

Princeton, New Jersey 08540

X
H. Shapiro

Duramic Products Inc. 426 Commercial Avenue

Palisades Park, New Jersey 07650

Earl P. Shapland

Flo-Con Systems, Inc. Box 355

Fisher, Illinois 61843

Lawrence Sharper

Frankford Arsenal

Philadelphia, Pennsylvania 19137

William A. Shelton

National Lead Co. 4511 Hyde Park Blvd.

Niagara Falls, New York 14305

Charles R. Sherman

The Babcock & Wilcox Company

Lynchburg, Virginia 24505

C. R. Shuff

American Standard, Inc. Post Office Box A

Paintsville, Kentucky 41240

Melvin W. Shupe

Los Alamos Scientific Lab

Los Alamos, New Mexico 87544

Carl J. Smith

Dow Chemical Company

Rocky Flats Division

Post Office Box 888

Golden, Colorado 80401

K. M. Smith

Applied Physics Laboratory

Johns Hopkins University

Scaggsville, Maryland 20910

Leonard I. Smith

Norton Company

1 New Bond Street

Worcester, Massachusetts 01606

F. X. Sorg

Stackpole Carbon Company

Stackpole Street

St. Marys, Pennsylvania 15857

Joseph Stanislao

Department Chemical Engineering

University of Rhode Island

Kingston, Rhode Island 02881

John M. Stanley

Griffin Pipe Products Co.

Orchard Hill Road

Milledgeville, Georgia 31061

H. Stuart Starrett

Southern Research Institute

2000 Ninth Avenue South

BiiTningham, Alabama 35205

R. J. Stokes

Honeywell Corporate Research

500 Washington Avenue, South

Hopkins, Minnesota 55343

H. Nathan Stone

AVCO Bay State Abrasives Div.

Westboro, Massachusetts 01581

R. C. Sundahl

Bell Telephone Laboratories

555 Union Boulevard

Allentown, Pennsylvania 18103

C. J. Swafford

Georgia Tech

Atlanta, Georgia 30332

Jeff Templeton

James H. Rhodes & Company

1026 West Jackson Boulevard

Chicago, Illinois 60607

George R. Thompson

Commonwealth Scientific Corp.

500 Pendleton Street

Alexandria, Virginia 22314

Robert L. Tienken

International Precision Ferrite

141 Ulster Avenue

Saugerties, New York 12477

Ronald J. Tillen

American Optical Corporation

Southbridge, Mass. 01550

James J. Tuzzeo

General Electric Company

6325 Huntley Road

Worthington, Ohio 43085

John E. Tydings

Naval Ordnance Laboratory

White Oak, Maryland

Ray Van Pamel

Arnold Engineering

Post Office Box G

Marengo, Illinois 60152

J. B. Wachtman, Jr. B308 Materials Building

National Bureau of Standards

Washington, D.C. 20234

Basil E. Walker, Jr. Naval Research Laboratory

Code 6130

Washington, D.C. 20390

Thomas W. Wallace

Westinghouse

Cheswick, Pennsylvania 15024

Addison Watkins

American Lava Corporation

Cherokee Boulevard

Chattanooga, Tennessee 37405

G. Wells

International Business Machines

Dept. 204, 1701 North St. Endicott, New York 13760

Warren W. West

E. I. du Pont de Nemours & Co.

Eastern Laboratory

Post Office Box 67

Gibbstovni, New Jersey 08027

Rodgers Westlake

Diamonite Products Mfg. Co.

Shreve, Ohio 44676

A. R. C. Westwood

R.LA.S.

Martin Marietta Corporation

1450 South Rolling Road

Baltimore, Maryland 21227

E. W. White

Pennsylvania State University

161 Engineering Science Bldg.

University Park, Penna. 16802

Sheldon Wiederhorn

Inorganic Materials Division

National Bureau of Standards

Washington, D.C. 20234

James O. Williams

C-E Refractories

Valley Forge, Pennsylvania 19481

William A. Wilson

National Bureau of Standards

325 South Broadway

Boulder, Colorado 80302

Vladimir E. Wolkodoff

Coors Porcelain Company

600 9th Street

Golden, Colorado 80401

David S. Yang

C-B Refractories

2563 West 18th Street

Chicago, Illinois 60608

xi

The Science of Ceramic Machining

and Surface Finishing

Welcome

E. Horowitz

Institute for Materials Research National Bureau of Standards Washington, D.C. 20234

Good morning and welcome to the National Bureau of Standards. I can see from the

time on the NBS clock, and the program schedule, that this is going to be the shortest

welcoming address in the history of our symposium series. This conference is being co- sponsored by the American Ceramic Society (Baltimore-Washington Section), the Office of Naval Research, and the National Bureau of Standards, and NBS is pleased to be

both a host and a participant in these proceedings. We at NBS consider meetings of this

type to be extremely important because they are a prime mechanism for developing discussion, promoting communication, and for coupling science and technology. This comment is particularly appropriate in view of your chairman's remark that this is the first

truly scientific meeting dealing specifically with ceramic machining and surface finishing.

I would like to say to those of you who are here for the first time, and particularly

the visitors from the six foreign countries, that we will do all that we can to make this

meeting productive, and your visit to NBS enjoyable and informative. Please call upon

the NBS staiT if there is anything that we can do to be of assistance while you are here.

From reading the published schedule it is clear that the sponsors have prepared an ex- tremely interesting and a full program. I think we ought to get on with the business of

the symposium and let me wish you a fruitful meeting.

I. INTRODUCTION

Session Chairman

S. J. Schneider, National Bureau of Standards

Science and the Machining and Surface Finishing of Ceramics

R. W. Rice

Naval Research Laboratory

Washington, D. C. 20390

C. F. Bersch

Naval Air Systems Command

Washington, D. C. 20360

A. M. Diness

Office of Naval Research

Arlington, Virginia 22217

Historically, the technology of machining and surface finishing of ceramics has

advanced by empirical methods. Further advances appear limited without gaining basic

knowledge of the mechanisms of removing or redistributing material from or on

ceramic surfaces. Because of the diversity of ceramic behavior, this is a challenging

scientific problem. However, meeting this challenge can result not only in improved

or new machining and finishing processes, but also in greater understanding of the

character of the resultant surface and sub-surface, and hence of the surface-dependent

properties.

Key words: Machining; machining costs; methods of machining and finishing; surfacedependent properties; surface finishing.

1
1. The Prior Status of Ceramic

Machining and Surface Finishing

Modern methods of machining and finishing

of ceramics have grown empirically from such

earlier arts as machining and finishing of glassy

bodies, polishing of jewels, and applicable metals technology. The improvement of properties and behavior which depend on machining

and finishing has also been empirical, with the

possible exception of optical properties. While

some characterization of resultant surfaces

has been carried out, it has been limited. There

has been almost no effort to characterize the

subsurface of machined or finished ceramics.

Little effort has been directed to understanding the basis of surface generation or alteration processes which control machining and

surface finishing operations as well as the

character of the resultant surfaces and dependent properties. The materials removal

processes responsible for surface generation

are complex and arise from the interplay between grinding operations and microstructural features of the workpiece.

Chemistry and physics provide some insight into the mechanisms of chemical and

thermal (annealing and flame polishing) finishing. However, little basic understanding of

the mechanisms of most other machining and

surface finishing methods for ceramics, especially mechanical modes, now exists because

they have not had the benefit of fundamental

research.

2. The Motivation for Studying

Ceramic Machining and Surface

Finishing

There are two basic motivations for studying

ceramic machining and finishing: intellectual

and economic. The intellectual motivation is the

quest for knowledge and understanding which

stimulates the scientist's endeavors. The eco- nomic motivation is to improve the processes in

order to achieve lower costs, greater productivity, and improved products. Basic knowledge

of (1) the machining and finishing processes,

(2) the resultant character of the surface and

affected subsurface, and (3) the consequent

effect on ceramic properties is needed in order

to satisfy both of these motivations.

Intellectual motivation in its broadest sense

arises because of the importance and scope of

ceramic materials and the wide diversity of

kinds of behavior they exhibit. The importance

results from ceramics being one of the three

major classes of solid (ceramic, metallic, and

organic) materials. Thus, lack of knowledge

and understanding of them represents a major

scientific gap. The scope of ceramic materials

in terms of structure, bonding, defects, and

resultant properties presents a unique challenge and test of scientific theories. More specifically, machining and finishing are important

to ceramists and other scientists whose specialization may be far from this area because of

its importance in determining the character

of ceramic surface and subsurface regions and

thus many important properties. Optical and

mechanical properties are obvious ones; magnetic and electrical properties can also be affected. The latter become more important as miniaturization of many parts progresses (with

concomitant increases in surface-to-volume

ratio).

The economic motivation is very significant

because of two factors. The first is the impact

of existing machining and finishing practices

on the cost of ceramic components. Machining

costs are very commonly of the order of 50

percent, and not infrequently in excess of 90

percent, of the total cost. Thus improvements

in machining can have a major impact on the

cost of parts. The second factor is the increasing demands of technology. Better performance

is demanded of existing bodies, especially with

respect to mechanical factors such as the level

and reliability of strength. In addition, new

bodies that are more diflficult to machine than

conventional ones are being developed to meet

new needs. Finally, technology places new and

more severe demands on both old and new materials. Finer tolerances, e.g., in finishing of

gas bearings for gyroscopes and magnetic re- cording heads, are obvious examples. However,

control of special surface topographies can also

be important. These include controlled surface

depressions for fluid reservoirs (e.g., for lubrication) or controlled surface roughness (e.g.,

for friction).

The intellectual and economic motivations

are mutually dependent. Obviously, the immediate or ultimate cost of science must come

from economic gain. However, further eco- nomic gain from machining and finishing ce- ramics is increasingly dependent on gaining

basic understanding, the current lack of which

limits applied goals. Substantial gains have

been obtained empirically in the past, but cutand-try techniques are unpredictable and can

be excessive in cost when applied to ceramic

components fulfilling critical functions.

3. The Challenge

Gaining an understanding of both the processes and effects of machining and finishing

of ceramics presents many challenges. The multiple and often interacting parameters of the

tool and the work piece in grinding and polishing are complex, with very limited basic understanding from which to start. The situation is

compounded by the use of other mechanical

2
methods such as single point, ultrasonic, and

fluid-driven abrasive machining", and by nonmechanical means such as chemical, electrical

discharge, laser, etc. The same range of machining and finishing

methods has been applied to metals. However,

ceramics present additional challenges. First

is the general absence of macroscopic ductility

to inhibit fracture inward from the surface.

Second is the much broader range of structure

and bonding and related intrinsic properties

such as hardness, conductivity, and elastic modulus. Third is the broader range of microstructures that must commonly be dealt with. Porosity and second phases more frequently occur

in many commonly used ceramic bodies. Further, ceramic bodies, especially in recent years,

have been produced with substantially finer

grain sizes (e.g., 0.1-1 //.m) than are common

in most metals.

Because of these greater challenges with

ceramics and a lower level of research than in

metals, ceramics start from a poorer background of knowledge. For example, it is only

now coming to light that during mechanical

machining and finishing considerable microplasticity occurs in the surface of crystalline

ceramics and possibly some type of deformation

occurs in the surface of amorphous materials.

However, our knowledge of the details of these

important processes is very poor, much more

so than in metals.

4. The Rewards

The general rewards expected from the study

of ceramic machining and surface finishing

are, of course, the scientific understanding and

economic gain that result as well as the attainment of improved property levels with higher

degrees of reliability and reproducibility. The

recognition by, or reminder to, many scientists

that complex problems of practical importance

in this area can involve stimulating, challenging, and fruitful research is hopefully also to

be realized. More specifically, it seems reasonable to predict that an understanding of the

mechanisms of removal and the character of

the resultant surface can be gained, in part,

by one aiding the other. Understanding these,

in turn, can increase our knowledge of surfacedependent, e.g., mechanical, properties and

hence provide the basis for improving these

properties. Any or all of these advances may,

in turn, lead to the develoument of new machining tools, some possibly designed explicitly for

ceramics, and to new or broader apnlication of

ceramics, especially under demanding mechanical conditions or for electromagnetic components.

Among the specific trends that appear is the

use of greater removal rates during grinding

of many ceramics, and hence possible increased

productivity. It seems clear that improvement

or evolution will occur in properties, economics,

rates of removal, tool and wheel life, and so on. On a broader scale, there should be a greater

choice between mechanical and nonmechanical

means of machining and finishing, as well as more choices within these two categories for

many materials and their applications. The

technology involved in the field of ceramic

machining and finishing should mature rapidly

if necessity, which often prompts invention,

is recognized, and it can draw upon a growing

base of sound scientific research and understanding.

This conference and the resulting volume

should provide a good basis and a substantial

first step for further scientific study and practical imnrovement of ceramic machining and

surface finishing.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

3

11. TECHNIQUES AND MECHANISMS OF REMOVAL

AND SHAPING

Session Chairmen

R. W. Rice, Naval Research Laboratory

V. E. Wolkodoff, Coors Porcelain Company

A. H. Heuer, Case Western Reserve University

A. Mechanical Methods

Mechanical Methods of Ceramic Finishing

Peter J. Gielisse and Joseph Stanislao

University of Rhode Island, Kingston, Rhode Island 02881

Mechanical methods of finishing ceramics have been reviewed based on current

practice as well as past and present research results. A total systems evaluation

involves technical process and economic parameters. State of the Art information

indicates a need for scientific evaluation of the basic stock removal process.

Experiments simulating grinding with single point diamond tools have generated

quantitative and qualitative data on mechanical, dynamic and thermal aspects of the

stock removal process. Generation of the data involved specially designed techniques.

Analysis of the results have been made in terms of mechanical, geometric and

material parameters. A mechanical and thermodynamic model is presented.

Key words: Abrasive; ceramics; diamond; fracture; grinding; grinding forces; grinding temperatures; machining; polishing; residual stresses; stock removal; surface

analysis ; theory of grinding.

1 . Introduction

1.1. General

The objective of this presentation is to review

the mechanical methods of ceramic finishing

present, the state of the art, and explore the

knowledge to-date on the theory of ceramic

finishing and the effect of the finishing system

on the final product.

Mechanical removal or grinding of ceramics

has long been an art, the progress of which

has almost entirely been based on experience.

True scientific knowledge in this area has been

meager. This is surprising since a major factor

in the manufacture or fabrication of ceramic

products is detailed knowledge of the variables

and their interrelation in the processing and

finishing cycles. Moreover, the critical link

between the processed part and the characteristics of the end product is provided by the

finishing cycle.

1.2. The Ceramic Finishing System

The two criteria for the selection of a ceramic

finishing system are: (a) the economic considerations infiuenced by stock removal rates,

tool life, initial cost, and grinding ratios (ratio

of stock removal to wheel wear), and (b) the

technical considerations. Economic factors have

been and still are of major concern in the finishing cycle. For this reason, and also because

they are more easily evaluated, the major emphasis in ceramic finishing studies has been on

economics.

However, a better understanding of the technical aspects has lately become very desirable,

if not imperative, especially with the many

types of present-day high performance ceramics. Fundamental technical know-how in this

area is particularly needed for: • elucidating the basic stock removal process

in ceramics, aimed at the development of

generally applicable ceramic finishing

data, • understanding the nature of the influence

of the finishing cycle on the finished product, • accelerating improvement in the application and methods of finishing systems, • guaranteeing reproducibility of each step

of the finishing process and end product, • evaluating the pros and cons of the various

available methods of ceramic finishing, • exploring the possibilities for preventing,

minimizing or removing deleterious results.

439-921 O - 72 - 2
• enhancing knowledge of the behavior of or resistance of ceramics to various types

of stress systems peculiar to the finishing

mode,

• improving the economic factors which will

enhance the feasibility of wider application of ceramics in society's technology.

With demands for specific properties in the

finished product, it is imperative that all the

capabilities, as well as the limitations of finishing techniques, be examined and evaluated.

This evaluation can only be achieved when all

available scientific and engineering knowledge

is applied to establish optimum systems in order

to gain the desired reliability in ceramic products.

Of all presently available stock removal processes—chemical, mechanical and other techniques—grinding is the furthest developed and

most widely applied in the ceramic industry

[1] \ Grinding has become a highly effective

means of stock removal because of the continual improvements which have been made

in controlling the various parameters in this

process.

In the choice of abrasive, a diamond abrasive

system appears most versatile and, in many

cases, except for very soft ceramics, is the only

choice. A diamond abrasive system can handle

almost all degrees of finishing on practically

every type of ceramic material. It is available

for finishing in every grinding, cutting, or

drilling mode on most sizes and with standard

equipment. Its real cost, i.e., cost per finished

part, is low. Furthermore, the diamond abrasive

system can combine most of the "ideal" re- quirements in one stock removal system. Silicon

carbide can be quite effective when applied in

the appropriate mode, but is not generally competitive when stock removal rate has to be

optimized. Aluminum oxide is used in finishing

"green" ceramics.

Grinding, specifically diamond grinding, is

thus recognized as a highly effective means

of stock removal in ceramic systems. Present

insufficient knowledge and lack of availability

of information contribute to an often inadequate set of criteria being used in the selection

of a grinding wheel for a specific application

[2]. Moreover, this approach of random selection totally ignores the need for a full systems

analysis of the complex energy (thermal, mechanical, electrical and chemical) mass interaction in ceramic finishing, to produce a fully

characterized end product (workpiece). The

choice of mode of application and abrasive does

not define the system. A total systems study

may in this context be regarded as the ultimate

goal. Unfortunately, the behavior of many of

the variable parameters as a function of total

systems performance is virtually unknown. As

an example, it is known that different surface

finishes can be obtained under similar sets

of conditions on various bodies having different

structural and or compositional characteristics.

Also, the result of the mass-energy interaction

of the finishing cycle may be very desirable in

one case but no so in another—particularly

when it is realized that recrystallization, melt

formation, changes in hardness, polymorphism,

embrittlement, chemical solution, micro- and

macrocracking, grain growth, residual stresses,

formation of high dislocation densities, twining,

slip, and the like may occur and change the

characteristics of the original body thus influencing its performance.

The mode of material removal in grinding of

ceramics affects the strength of the end product

[2]. Dry grinding may be less economical, but

wet grinding alters strength. One type or size

of diamond in a specific wheel matrix may or may not act similarly to another type in a different matrix. Chemistry of the process or en- vironment may influence the results.

In almost all cases, a highly selective, qualitative picture has emerged but no generally

applicable quantitative data is available or has

been developed. A lack of knowledge about the

mechanism of abrasive material removal [3]

in these high energy rate transfer processes

has placed the selection of grinding parameters

on an entirely arbitrary and pragmatic basis.

Figure 1 schematically shows the simplified

PROCESS

SPECIFICATIOM

/TECHNICALX COMPUTER

J PROGRAM

^MODEL,

ECONOMIC

PART

CHARACTERIZATION

TECHNICAL

INTERNAL EXTERNAL

' Figures in brackets indicate the literature references at the end of this paper.

T

Figure 1. Schematic presentation of a

ceramic finishing system.

6
ceramic finishing cycle, beginning with the

needs as expressed in a set of finished piece

specifications. The system can, via the economic

loop and the technical loop, provide information

of importance to the economics or the technical

aspects of the process. The significant differ- ence between the two components of the system is that the output provides the economic

picture directly, while the technical information

is not provided until the part characterization

has been performed. It may further be noted

that the technical information subdivides into

two parts: the internal and external components. The external aspects would include such

matters as geometry, finish, tolerances, roughness and waviness; while the internal aspects

primarily refer to micro- and/or macrocracking, residual stresses, chemical contamination,

alteration of transformation, resolidified or re- deposited material, i.e., surface and possibly

bulk alterations. A full systems analysis of the

ceramic grinding process would produce a technical and economic model, close the cycles and

provide for the possibility of entering particular specifications directly into the system to

produce a ceramic product, fully characterized,

in the economic and technical sense. With the present state of the art, such a

cycle of information cannot be expected to be

completed or a complete theory or model of

ceramic finishing devised in the very near future. However, it is possible to devise correlations which describe certain facets of the components of the system. Most important of all is the need to start an organized effort in the

systems analysis of the ceramic finishing cycle,

which should provide information generally

available and applicable to random material or

application requirements. Some knowledge is

available for metallic finishing systems, in cutting as well as grinding, but virtually none has

been developed for ceramic systems. It is at

present impossible to consult one single source which gives normalized information significant

in the determination of the choice of grinding

wheel to be used in any one ceramic finishing

application.

Any such study should start with the generation of information which will further the un- derstanding of the fundamental and basic as- pects of the ceramic finishing process. Such

studies cannot escape the need for basic analyses with both single-point diamond tools and

multipoint grinding implements (grinding

wheels).

Again, a further choice indicates the primary

need to be for single-point cutting tool studies

since single-point studies unequivocally establish the true nature of many of the phenomena

which are part of the ceramic finishing cycle.

This can be done without interference from the

many uncontrollable factors intrinsic to grinding wheels (variations in hardness, shape, size,

concentration, bonding, etc.,) or need for re- course to complex statistical analysis. Statistical analysis is generally required when multipoint cutting tools are used. Such information

by itself can, however, never establish the

basic physical nature of the stock removal process. The research reported here deals, therefore, with the elements of a qualitative and subsequent quantitative study of the stock removal

phenomena in ceramics by means of singlepoint (diamond) tools of known geometry.

A full systems analysis of a finishing cycle

on ceramics is difficult due to several factors

which comprise its very nature: • the multipurposed objective in finishing • the nature of ceramics

• the nature of the finishing cycle • the inability to properly characterize ce- ramic surfaces internally and externally • the lack of knowledge of a ceramic's re- action to conditions unique to the material

removal system • the still inadequate ways of specifying, in

terms of physically recognizable and measurable quantities, the properties of a finished product.

1.3. Nature of Ceramic Finishing

The purpose of stock removal is generally

to obtain a simple or complex final shape; to

reduce to size and to specific, often highly ac- curate, dimensions; and to generate desirable

surface finishes or properties as well as eliminate undesirable surface or bulk characteristics.

It is thus clear that the requirements put on

a bulk stock removal process are quite different

from that which is called upon to produce a

less than one micron finish in a fragile micro

circuit substrate. The repeatability needed in

the latter operation is not necessarily significant if gross shape features are the only objective. Detailed technical knowledge of the function and interaction of each parameter in a

variety of stock removal systems employed is

thus required to obtain the desired result.

Ceramics vary widely in characteristics

which normally influence the stock removal

process: hardness, compressibility, chemical

reactivity, brittleness or ductility, elastic impedance, and resistance to thermal or mechanical stresses. The physical nature of ceramics

further influences their behavior in stock re- moval. They can be homogeneous as in single

crystals or heterogeneous as in the majority

of technical and structural ceramics. Great differences in mode of stock removal may be observed depending on the ceramics grain size,

degree and type of porosity, impurity level,

presence or absence as well as nature of a sec- ondary amorphous phase, and most important7
ly, the microstructural features which are the

result of an often inadequately controlled or understood processing history. Much could be

gained if properly characterized ceramics in

which only a specific parameter such as grain

size were a variable could be made available.

This would considerably enhance the possibilities of investigators carrying out their re- searches on typical ceramic systems rather than

on what is often alleged to be the less important,

more idealized controllable systems as are, for

instance, found in single crystals. This is not

to say that work on such systems is not veryimportant in terms of fundamental characterization of the basic principles but transition to

the more complex and more generally encountered ceramics should be made and efforts

should be expanded to facilitate this work.

Finishing of ceramics, through grinding, is

in many ways similar to other commonly applied removal processes such as milling, cutting,

and shaping, and has many of the facets of the

thoroughly investigated field of finishing metallic materials. Yet there are significant diflfer- ences. Grinding per se is probably one of the

least understood processes. This is due to ran- domness of physical characteristics such as

grit geometry or morphology, friability or hardness of the grain, degree of bonding to the

matrix, and the nature and influence of the

matrix on the process as well as the concentration and grain size of the abrasive itself.

From the point of view of the finishing "tool"

no one set of characteristics can be specified as is typical in, e.g., a metal cutting tool. Furthermore, significant interactions are hard to

observe or study. Parameters common in metal

finishing studies such as rake angle, depth of

cut, shear angle, and chip geometry are not

fixed but are variable in process which often

necessitates the application of statistical analysis. Although notable, knowledge regarding

the physical process often goes undetected.

This is particularly true with regard to the

cutting action of individual grains. The machine process creates further differences between abrasive machining and classical machining. Abrasive finishing is presently done

at tool speeds about an order of magnitude

higher than in cutting. Depth of cut, of the

individual grain, is considerably less. Circular versus linear tool motion, the possibility of "up

grinding" or "down grinding," and continuous

versus discontinuous tool-workpiece interaction

all add additional factors to be considered.

Although wheel speeds in grinding are con- sidered high as compared to speeds of cutting

tools, present research indicates the desirability

of increasing speeds well beyond those presently employed. Further complications arise

when the different modes of abrasive machining are considered, each of which is distinctly

different from the other in the mechanism

whereby material is removed from the workpiece. This topic will be further discussed in

section 2 of this paper.

Proper characterization of a generated sur- face is, as figure 1 shows, a very important

part of a full ceramic finishing system. Characterization of the surface of a processed material in a particular application can, in general,

be classified into one of three modes, according

to chemical, physical, or geometrical parameters

(see fig. 2). This is not necessarily synonymous

SURFACE CHARACTERIZATION IN CERAMICS

METHODS-CONCEPTS

Physical Chemical (a) Geometric

Others

icy

APPLICATION PERFORMANCE

Substrates

Seals

Bearings

Nozzles Others

PROCESS PARAMETERS + CHARACTERISTIC FEATURES + PERFORMANCE PARAMETERS

MACHINING PARAMETERS

Tool

Machine system Workplece

MACHINED PRODUCT

RMS, BA, Tp, %

ni>J , V , ,

' OPTIMAL MACHINED ' PRODUCT

Figure 2. Surface characterization in ceramics. Schematic of interrelated

components.

8
with a subdivision of the techniques being geometrical, physical, oi' chemical, i.e., a chemical

technique may point toward the nature and

extent of physical inhomogenieties in a surface, such as etching reveals dislocations. Discussion of the physical and chemical techniques

of surface analysis is beyond the scope of this

article. The topic will be extensively discussed

in other sections of this conference.

At the moment geometrical characterization

of a surface is readily carried out, widely investigated and developed, and generally ac- cepted for comparative purposes. It should be

pointed out, however, that the function of a

surface is invariably influenced by its chemical

and physical parameters. In many, if not most

cases, the geometrical constitution of a surface

will markedly influence the chemical and physical aspects, as well; a physical property, such

as the coefficient of friction, is greatly influenced by the "roughness" of two mating surfaces, the reaction (chemical) between the tool

and the workpiece (tool life) in a cutting operation strongly depends on the surface geometry

of the tool as it engages the workpiece. Analysis

of the functionality of a surface in a particular

environment invariably requires geometrical

analysis.

It is important to point out here that the

development of a surface test, which will characterize and guarantee the performance of a surface in a particular application, presupposes

that a surface analysis has been made. Attempting to obtain surfaces of lower and lower roughness readings of the peak (overall roughness)

or averaged (AA or RMS) variety, and equating this with surface improvement, constitutes

a positive result only if it has been determined

that there is a direct correlation between this

measured parameter and the functionality of the

surface. At present, it must probably be as- sumed that, in the majority of surface

test is attempted ahead of the more critical

surface analysis. Surface analysis therefore

attempts to establish the parameter or parameters most directly connected with the optimum functionality of a surface.

The objectives of a complete surface analysis

may be thought of as consisting of four parts

the improvement of the surface function, the

influence of every processing step, the guarantee

of part reproducibility, and the provision of

meaningful criterion on surface conditions.

The first objective is the improvement of the

surface function or functionality. The solution

of this problem requires the determination of

the often missing link between the surface

conditions of the processed part and its ultimate

performance in the application. At present no

standards exist in even the most general of

cases. The ultimate solution lies in the close

cooperation between the producer and user in

determining the critical parameters of the process under identical conditions. Specifications

on parts will only have meaning when such

specifications are determined under identical

test conditions. This is of particular importance

when it is realized that the functionality of a

surface may critically depend on the processing

mode. Standards on a ground surface may have

to be altered when such a surface is found to be

more economically produced by a turning operation.

The second objective is the determination

of the influence of each and every step in the

processing on the final functionality of the part.

Considerable economic advantage may be obtained if the sequence of operations leading

most advantageously to the desired surface

can be determined. It might thus be possible

to replace particular operations with others, or eliminate them completely. A final polishing

or lapping operation could be prevented by

judicious choice of the previous grinding parameters (speeds, feeds, hardness of wheel, etc.)

The third objective is the guarantee of part

reproducibility. Without the ability to guarantee part reproducibility, uniform part performance will not be possible. Part reproducibility

requires the investigation of the influence of

the individual process parameters of any one

processing mode on the final surface conditions.

Again, in the case of a grinding operation the

influence of the variation of parameters, such as

speed, grain size and hardness of the wheel,

and often even the type of grinder, should be

known. If such parameters can be or are already being controlled, it suffices to relate their

variation to the surface conditions of the single

part and thus guarantee reproducibility.

The fourth objective is to provide a meaningful criterion on surface conditions. The importance here lies in the establishment of an objective measurement to replace the often

subjective opinions or criteria (reflectivity,

appearance, etc.) which easily lead to misunderstanding. An objective measurement eliminates such misunderstandings.

It should be apparent that the task of surface

finishing is not a simple one and cannot be

carried out without intensive surface analysis.

Significant advances have been made during the

last decade in determining the functionality of a finished ceramic workpiece. This is primarily

true with regard to the influence of surface and

subsurface characteristics on mechanical properties.

1.4. Effect of Ceramic Finishing on the

Workpiece

Machining operations, such as grinding, lapping and polishing, leave residual surface defects and, as a result, these defects impair the

mechanical strength of a ceramic. The nature

9
of the damage, the physical mechanism by

which it is produced, and above all the extent

of the depth below the surface of such damage

is not very well known. Surface damage (cracks

or flaws left in the machined surface) is generally distinguished from subsurface damage

(introduced lattice defects). Brittle materials

show the development of surface damage more

than subsurface damage which appears to dominate in semibrittle ceramics. The two types

of damage, however, can occur simultaneously

and their relative extent depends on the stock

removal mechanism (speeds, feed, wheel or abrasive characteristic) and the nature of the

polycrystalline ceramic. Koepke and Stokes [4]

have studied the nature and extent of subsurface damage in MgO after grinding with ALO^

and diamond wheels. Their etched ALO3 ground

specimen shows a distinct plastically deformed

layer, the thickness of which increases with the

depth of cut. Slip bands in (101) and (101)

planes penetrate beneath the plastically deformed layer at high depth of cut and feed rates, as may be seen in figure 3. The extent of the

damage is higher in the dry grind than the wet

grind mode. Diamond grinding produces a less

well defined deformed layer with evidence for

formation of screw dislocations below and parallel to the ground surface. No simple relationship was found between the depth of damage

and the material removal rate as in the case

of AI2O3 grinding. Lower limit of depth of damage was found to be about 50 /^ms or 1.5 times

that with an alumina wheel. This seems to clearly indicate a difference between the reaction of

ALO., and diamond particles when interacting

with the same workpiece. Diamond tends to

deform less, grain penetration is higher and

energy absorption and stress concentration is

higher for the workpiece in the diamond case. Surface damage showed again a diiference

in behavior for AI2O3 and diamond. AI0O3 grinding produced a network of surface cracks nor- mal to the grinding direction. This is due to

Figure 3. Etch pitted damage layers in magnesium

oxide single crystals as a result of grinding magnification app. 95X. From Koepke and Stokes [4].

tensile stresses generated by thermal quenching. AI2O3 grinding can be high temnerature

grinding due to the low thermal conductivity of

the grain and possibility of wheel loading. At

these high temperatures material removal can

be made by plastic flow. The material often becomes burnished. In diamond grinding of MgO

crack formation is limited because of cool

grinding with unloaded wheels. Gielisse et al. [1] have previously indicated that this should

be the case. Section 3 of this paper has worked

out this concept further. Whereas the work of

Koepke and Stokes indicates a definite influence

on surface and subsurface properties as a result

of grinding, the nature and extent of the damage depends on the parameters of the grinding

mode. Cool grinding with nonloading wheels,

containing discrete cutting points of which the

penetration depth and stress concentration for

any one ceramic can be controlled, appear indicated. Rice [5] has made similar investigations and observed dense layers of dislocations,

slip bands and sometimes cracks in MgO, with

the extent of the phenomena again depending

on machine and material parameters. Fine grain

size and high yield stress limits the depth of

the altered layer. There also is a direction dependency (anisotropy). Inhibition of motion

and surface work hardening were observed.

Annealing of the finished specimen can reduce

the grinding damage considerably. The effects

have been observed in both single and polycrystalline samples. Rice's work also includes investigations on the lower yield strength CaO.

X-ray techniques have been used by Cutter and

McPherson [6] to examine handlapped or sanded single crystals of hot pressed MgO. Dense

layers of dislocations (10 dislocations/cm^)

were observed below the surface. Harrison [7]

,

as well as Evans and Davidge fS], have observed cracks in diamond cut MgO crystals. Ion

beam thinning in conjunction with transmission

electron microscopy is presently being used to

study subsurface structure in polycrystalline

ceramics. Hockey [9] has used these techniques

to provide direct evidence of plastic deformation occurring as a result of abrading aluminum

oxide. Single and polycrystalline specimens of

AI2O3 show high dislocation densities, basal as

well as nonbasal, within the near surface areas when polished with diamond as fine as one

quarter of a micron. This may indicate a considerable amount of plastic deformation during

abrasive wear. Regions of the highest dislocation densities defined the traces of surface

scratches produced by individual abrasive particles. This is shown in figure 4. Rice has used

R.F and ion beam sputtering techniques to ex- plore the extent of damage in surface finishing.

Figure 5 shows the effect of R.F. sputtering of

scribe marks on cleaved MgO crystals. The na- ture of Hockey's techniques enabled him to

10
Figure 4. Dislocation produced as a result of diamond

polishing polycrystalline AI2O3 magnification lO.OOOX.

Courtesy B.J. Hockey, National Bureau of Standards,

Washington, D.C.

observe the introduction of relatively large

compressive residual surface stresses [9].

It is thus rather well established that grinding of ceramics is a mutative process. The question is whether the mutation is detrimental or advantageous and what factors play a major

role in their generation or prevention. The

formation of regions of high density dislocations as a result of thermal and/or mechanical

effects in surface finishing should lead to sur- face strain or work hardening. This has been

observed by Rice [10] to be particularly pronounced in large grain size, low yield strength

materials. Fine grain size materials can be

weakened due to the more limited degree of

workhardening and the higher stresses reached

in the finer flaws of these materials. The mutative effect of surface finishing processes can

be more easily quantified when expressed in

terms of the surface residual stresses resulting

from the combined physical processes that

cause them; thermal, chemical or mechanical.

Plastic deformation due to non-uniform dilation, volume changes as a result of chemical

reactions, phase transformations, precipitation,

melting and nonuniform plastic deformation

due to mechanical working are respective ex- amples of such processes. When considering

strength, compressive residual stresses are generally regarded favorable in ceramics. The

compressively stressed surfaces prevent the

Figure 5. R.F. sputtered scribe marks on cleaved MgO

crystals. Two hours (top) and seven hours (bottom).

Courtesy Roy W. Rice, U.S. Naval Research Laboratories, Washington, D.C.

propagation of cracks. Both compressive and

tensile stresses must be controlled when a high

degree of flatness in thin, low cross-section

parts, such as in substrates, is required. Residual stresses due to finishing can lead to failure

in thermally cycled ceramics such as in ruby

laser applications. Above all, a convenient nondestructive method for determining subsurface

residual stresses and stress distribution is needed. No satisfactory method is available as yet.

Hara et al., [11], using the x-ray stress measurement technique, have recently shown that

grinding induces compressive residual stress

(>100 kg/mm 2) in cemented WC which de- creases with increasing distance from the sur- face and finally become tensile. They also measured the lattice distortion and came to the

remarkable conclusion that residual stresses can

extend beyond the depth zone of lattice distortion. Residual stresses can be reduced by anil
GRIT SIZE

Figure 6. Average tensile strength of 99.5 alumina

as a function of diamond grit size. After data from

Sedlacek [12].

nealing. Annealing reduces dislocation density

as conclusively shown by Rice [5] and Hockey

[9] in their investigations on MgO and ALO.j

respectively. Knowledge of residual stress patterns should enable a more direct quantitative

way of studying mutative effects due to subsurface alterations in the workpiece.

Since grinding establishes the surface and

subsurface characteristics of the workpiece,

the influence of the various system parameters

should be examined. The work by Sedlacek [12]

has made significant contributions in showing

that diamond grit size has a definite and systematic effect on strength. Figure 6 is a plot

of average ultimate tensile strength as a function of diamond grit size. Data are from Sedlacek's investigations on 99.5 alumina using the

internal grinding mode. The effect is found re- gardless of previous grinding history. The

relationship is systematic and shows a decrease

in strength with increase in particle size. Similar results have been reported by McKinney and

Herbert [13] on SiC ground AlSiMag 614

(94% pure alumina near zero porosity). A re- lationship seems to exist between strength and

the depth of surface flaws if one assumes that

larger grain sizes leave larger surface cracks.

This can be further examined in figure 7 which

was plotted using data from Sedlacek [14]. It may be observed that larger grit size produced

a wider spread in and higher values of surface

roughness. The boundary lines were drawn in

to accentuate the apparent trend. The data do

not, however, indicate whether the effect may

be due to subsurface cracking generally observed with larger grit sizes. Laboratory results

[15] indicate a considerable increase in grinding temperature with increase in depth of cut

(i.e., increase in grit size). What appears very

significant is the influence of the matrix of the

grinding wheel. A 100 grit resin bonded wheel

was found to produce the same surface finish

and strength as a 32 grit metal bonded wheel.

A larger grit size can cause a much smaller

effective scratch when it is allowed to absorb

part of the grinding force in a low modulus

matrix. Sedlacek [14] further found a considerable effect of the rate of material removal

but distinguishes between infeed and table speed

both of which determine rate of removal. Heavy

infeeds are detrimental to strength. Dry grinding lowers strength presumably due to an in- crease in thermal upset. In general, the mode

of material removal influences the strength of

the final product. The extent of this influence

is considered by Sedlacek to be less than the

effect of environment on the strength of alumina. He holds a stress activated chemical reaction (stress corrosion) between alumina and

surface absorbed water responsible for the decreased strength of alumina in a wet environment. Our experiments in 0^, N^, air and a water mist show a considerable decrease in both

tangential and normal forces for a single grain

cutting through various types of ceramics in

the case of the H,0 mist (see fig. 8). A limited

difference is found between the gases. A larger

difference is apparent from material to material, pointing towards the importance of

microstructure (grain size and composition) of

the workpiece itself.

Careful monitoring of the dynamic aspects,

i.e., force pattern, will be relevant with regard

12
70,

60

50

2 40

<n 30- u

o

§ 20J

MATERIAL AO 94

TOOL 120'

DEPTH or CUT 003" CUTTING FORCE Fn TA8LE FEED 9

70,

60

50

240

CO 30

UJu§20

MATERIAL AO 94

TOOL l2(fO

DEPTH OF CUT OOJ' CUTTING FORCE F, TABLE FEED 9

MECHANICAL STOCK REMOVAL PROCESSES

1500 2000 2500 3000 3500 WHEEL SPEED (RPM)

1500 2000 2500 3000 3500 WHEEL SPEED (RPM)

1500 2000 2500 3000 3500 WHEEL SPEED (RPM)

1500 2000 2500 3000 3500 WHEEL SPEED (RPM)

MATERIAL AD 995

TOOL l2Cf O

DEPTH OF CUT <X)3" CUTTING FORCE Ff TABLE FEED 9

2000 2500 3000 3500

WST (Hf 0)

isbo idoo 2*0 sdoo 35bo

Figure 8. Influence of environment on grinding forces

in single point cutting of three types of alumina.

to the study of subsurface plastic deformation

as a result of the stresses imposed by the wheel.

These forces depend primarily on the properties

of the ceramic as will be shown in another sec- tion of this paper. Control of the thermal as- pects will influence the extent of the surface

damage.

2. Mechanical Finishing Processes

2.1. General

Stock removal by mechanical means can be

accomplished in one of two ways: on a microscale where material is removed by abrasive

interaction, atom by atom or molecule by molecule, or alternatively on a macroscale in which

the dimensions of the resultant particulate matter are orders of magnitude larger than that

of the basic molecular constituent (see fig. 9).

The former process leads to a perfectly smooth

GRINDING

macrofinishing

(Particulate concept)

POLISHING

microfinishing

(Molecular concept)

BAM

Cylindrical

Centerless

Surface

Tool and cutter

CAM

Point contact

Line contact

Area contact

FAM

Single mode

Figure 9. Subdivision of mechanical stock removal

processes.

surface if the process is carried on long enough

and a homogeneous material is assumed. In

principle, perfectly smooth surfaces could be

obtained in macrofinishing if the material were

removed by an internal shearing process as a

result of total and perfect plastic deformation.

Broadly speaking, mechanical stock removal in

metals is a process of this type albeit not totally

perfect, while stock removal in brittle or semibrittle ceramics is primarily the result of an

internal fracturing process. The macrofinishing

mode requires penetration of the surface which

is accomplished only if sufficient energy is imparted by the system on the workpiece to over- come its yield strength. The concepts of macroand microfinishing leads to the distinction

between the two fundamental types of abrasive

mechanical surface finishing—grinding and

polishing. Grinding requires penetration of the

grit into the workpiece surface and material

of macroscopic dimensions is removed. It may

be referred to as macrofinishing. In grinding

ceramics, as a result of insufficient plastic deformation, penetration of the abrasive produces

a crack field extending in front of the abrasive.

By multiple interaction the grinding particles are produced. Because of the need for penetration, grinding is a high pressure process.

Polishing does not require penetration, and

material of microscopic dimensions is removed.

This will be referred to as microfinishing, and

it is inherently a low pressure process. Polishing

can, of course, be done by other than mechanical

processes. Thermal processes in which surface

layers are melted have been applied; flame

13
polishing is an example. Chemical processes

such as chemical etching have been successfully

used. Combinations of mechanical and chemical

action are also possible. Polishing of boron car- bide surfaces with alumina or boron carbide

itself is an example. Several of these processes

will be discussed in detail in other sections.

2.2. Mechanical Methods of Surface

Finishing in Ceramics

Virtually all methods of ceramic surface finishing to-date are multipoint abrasive operations, as distinct from single point tool operations. Therefore, only multipoint processes will

be discussed.

Both grinding and polishing processes fall

into three fundamental categories based on the

way in which the abrasive is held: Bonded

Abrasive Machining (BAM), Contained Abrasive Machining (CAM), and Free Abrasive

Machining (FAM). A diagrammatic presentation of this subdivision is given in figure 9.

a. Bonded Abrasive Machining

Bonded abrasive machining can be equated

with standard wheel grinding. The abrasive is

fixed in a matrix of a vitreous, metallic or or- ganic nature. By feeding the rotating wheel

into the moving workpiece, the abrasive is

forced to penetrate it. The individual abrasive

grain intermittantly impacts and moves through

the workpiece at speeds of thousands of surface

feet per minute. Recent research indicates that

desirable speeds differ with materials [15]. Low

speeds can quickly generate very high forces.

Indications are that high speeds are generally

beneficial if the resultant high temperatures

are not detrimental. Truly high speed processes

(> 10,000 SFM) in ceramics have not been investigated sufficiently. This is discussed further

in section 3 of this paper. The nature of the

bond is all-important. The matrix should have

mechanical properties which enable the individual grains to effectively penetrate the workpiece without premature breakdown of the grain

itself. The mechanical characteristics of the

workpiece primarily dictate the choice of matrix

properties. Thermal properties must be compatible with the temperatures developed during

stock removal. The optimum elasticity of the

wheel under any one set of machining parameters is primarily determined by the choice

of matrix material but can be further influenced

by adjusting abrasive concentration, size, morphology, filler type and content and processing

parameters. Generally insufficient attention is

being paid to the possibilities of optimizing

the process by varying the above parameters.

The primary reason for this is lack of fundamental understanding and knowledge in this area. The performance of the abrasive in the

BAM process is determined by its abrasion re- sistance. The physical characteristics which

are known to play a major role are: (hot)

hardness, strength, morphology, chemical re- activity, frictional and thermal properties.

These have recently been reviewed [16] and

need not be further reviewed here. The above

considerations lead to the inescapable conclusion that for finishing ceramics with present

BAM techniques, diamond is the preferred abrasive. Some reports indicate that manufactured

diamond may, on ceramics, be preferred over

its natural counterpart [1]. Physical characteristics are more important than the abrasive's

origin, particularly since natural diamond can

be processed to yield closely resembling characteristics. A classification of diamond types on

the basis of physical characteristics is important since all types are used in distinctly different finishing: operations on ceramics. A brief

overview follows. See figure 10.

Low impact strength diamond (LTSD) is a

relatively friable diamond crystal essentially

anhedral, or without prominent crystalline

faces. Its large surface area and irregular geometry make possible excellent bonding and

retention in the resin matrix in which it is

almost exclusively used. Because of its relatively

high friability, it will generate new cutting

points on materials with a high modulus of

resilience (MOR), but would create undue wheel

wear on low MOR materials. Its applications in

ceramic systems is on very high yield strength

aluminas (99.9), Lucalox, or hot pressed polycrystalline alumina. The large number of as- perities make for low pressure per contact

point and possibilities of effective polishing.

Modified low impact strength diamond

(MLISD) is basically a friable-type crystal

surrounded by a high-modulus metal shell. It generates new cutting points on high MOR materials, but does so in a controlled manner. Retention of the abrasive and thermal conductivity are both enhanced. The net result is usually a

significant increase in grinding efficiency. A

nickel coated product is used in wet grinding

and a copper coated product in dry grinding

of more ductile ceramics such as cemented

tungsten carbide.

Medium impact strength diamond (MISD)

is a less friable or semifriable diamond in which

crystal faces are moderately well-developed.

Geometry is decidedly more blocky, generating

higher pressures per contact point and thus

more efficient penetration. MISD is a tougher

diamond than the low impact strength diamond

and is primarily used in metal bond systems

and is ideally suited for grinding glass and

ceramics. On low or high MOR materials it will cause undue wheel wear.

High impact strength diamond (HISD) is

a very strong crystal with large, well-developed

14
LOW IMPACT STRENGTH DIAMOND

\ DIAMOND L / CflVSTM I

I

WILti

• CONTROl RATE OF CUnih4G

POINT GENERATION ON HIGH MOR MATERIAL

• STIll GENERATE UNDUE

WHEEL WEAR ON LOW

MOR MATERIAL

LOW MODULUS BOND

PROVIDES

• MECHANICAL SUPPORT

• GOOD IMPACT ENERGY

TRANSFEB

HIGH MODULUS SHEU MEDIUM IMPACT STRENGTH DIAMONDWILL: • GENERATE NEW CimiNG

POINTS ON MODEBATEIY

HIGH MOR MATERIALS

• GENERATE UNDUE WHEEL WEAR ON LOW OR HIGH MOB

MATERIALS

HIGH MODULUS BOND

PROVIDES

• MECHANICAL SUPPOfiT

• ABRASION RESISTANCE

HIGH IMPACT STRENGTH DIAMOND

WILL: • WITHSTAND HIGH IMPACT lOADS ON lOW MOB

MATERIAIS

• GENERATE EXCESSIVE GLAZING AND PUHOUT ON

HIGH MOB MATERIALS

HIGH MODULUS BOND

PBOVIOES: • MECHANICAL SUPPOBT

• ABBASION BESISTANCE

Figure 10. Principal Types of Diamond and Their Major Applications.

cubic and octahedral crystal faces. It will withstand high impact loads on low MOR materials

because of its strength, but alternately (mainly

because of its smoothness and relatively low

surface area) it will generate excessive glazing

on high MOR materials and thus promote pullout from its metal bond. Its application in ce- ramics is in cutting and slicing of both natural

and manufactured products.

Relative grinding performance and efficiency

of the various types of diamond are shown in

figure 11. Generalized workpiece characterization as a function of the modulus of resilience

is given in figure 12. Other abrasives used in

BAM processes are SiC and AI2O3. Silicon car- bide is used in machining "green" ceramic

parts. Alumina has the advantage of not introducing materials of different chemical composition. The same operation is often performed

with carbide cutting tools in conventional turning or drilling operations. Bisque fired parts

could be handled by the same abrasives. Finish15
ing of bisque fired parts is not commonly practiced because of the additional processing step

involved in manufacturing the part. In addition

to the true matrix type vi^heels, plated wheels

MODIFIED LOW IMPAa

STRENGTH DIAMOND

MEDIUM IMPAa

0.1 1.0 10 100 1000

- MODULUS OF RESILIENCE -

2E

Figure 11. Relative grinding performance of principal diamond types as a function of modulus of

resilience. After data from reference 1.

.1 1.0 10 100 1000

- MODULUS OF RESILIENCE -

2E

Figure 12. Workpiece characterization [1] as function of the modulus of resilience.

in v^hich a monolayer of diamond is deposited

on a metal base are finding more v^^idespread application. They are particularly useful v^hen

complex shapes have to be imparted (form

wheels) or when process parameters require it, as in rotary ultrasonic machining.

The bonded abrasive machining method can

be applied in a variety of forms depending on

the finishing mode and the geometry of the

workpiece. Various types of BAM techniques are shown diagrammatically in figure 13. In

CYLINDRICAL GRINDING the workpiece

periphery of a rigidly supported rotating workpiece held between centers is ground by a wheel

revolving at high speeds while the work ro- tates at much lower speeds of 75—150 surface

feet per minute. The wheel is fed into the work,

and either workpiece or wheel is traversed.

Variants are the universal center type grinder

in which tool post and stock holder can be moved

around a vertical axis enabling bevels and

angles to be ground. Roll grinders are designed

to handle large diameter rolls such as those used

in the textile, paper and printing industry. In

chucking grinders used on short pieces, the

work is held in a chuck or collect rather than

between centers. Tool post grinders carry the

wheel for external or internal work on a tool

post or the compound rest of a lathe. The wheel

is directly driven from a motor shaft. CENTERI I HORIZONTAL SURFACE fiRINDING

I (rotary)

vertical surface srindiwg

(rotary or reciprocating)

Figure 13. Diagrammatic presentation of the major

bonded abrasive machining techniques.
LESS GRINDING can be done internally and

externally by supporting- the workpiece on a work rest blade and revolving it between a drive

wheel and the grinding wheel mounted on opposite sides of the cylindrical post. Additional

pressure rolls and supporting rolls are often

used in internal grinding of this type. This regulating wheel is often mounted at an angle to

the plane of the grinding wheel to accomplish

through feed. These grinders are limited to

grinding cylindrical parts and are most commonly used in mass production. INTERNAL

GRINDING can be of the chucking or centerless

type with relatively small wheels mounted on

spindles revolving at high speeds and traversing the workpiece by hydraulic movement of

the grinder bed on which the spindle is mounted.

The wheel is simultaneously fed into the workpiece. SURFACE GRINDING is used to produce flat surfaces or grind irregular surfaces

by means of shaped or formed wheels in form

grinding operations. The various types are based on the movement of the wheel into vertical or horizontal spindles while the table holding the workpiece may be reciprocating or ro- tary. The reciprocating horizontal surface

grinder is the all purpose shop and tool room

machine, while all other types are more generally applied where many parts need to be

produced in a mass production fashion. Vertical

spindle rotary table grinding is very common

in nonmetallic or ceramic applications. Frequently more than one grinding head is used

in vertical spindle, rotary table grinders. Work

is held by magnetic, electrostatic or vacuum flat

chucks or is cemented or clamped down on

the grinder's table. TOOL AND CUTTER

GRINDERS are used to sharpen the edges of

milling cutters, reamers, drills, etc. They are

outfitted to move the workpiece horizontally

and vertically. The wheel can move up or down

with a 360 degree rotation around a vertical

axis. By employing beveled or cup shaped

wheels, virtually any tool configuration can be

accommodated. NONCONVENTIONAL MECHANICAL TECHNIQUES can take several

forms. Abrasive jet machining is the removal

of material through the action of a focused

stream of fluid generally containing abrasive

particles. Ultrasonic abrasive machining cuts

material by high frequency repetitive impact

between a tool and workpiece with an abrasive

slurry in between. Rotary Ultrasonic Machining

(RUM) is a relatively new technique with particular promise for the ceramic area. It is,

therefore, discussed in somewhat greater detail. In RUM the actions of high speed rotary

grinding and reciprocation of the rotating tool

at ultrasonic frequencies are combined. Conventional ultrasonic machining does not employ

rotary motion and uses an abrasive slurry between a preformed tool and the workpiece.

Tapering of the hole, imbedding of abrasive

into the workpiece and limitations on stock re- moval are cited as the disadvantages of this

method, which are overcome by simultaneously

rotating and ultrasonically vibrating. The RUM

technique employs a bonded (plated) diamond

abrasive rotating at variable speeds up to 5000

RPM. The application of ultrasonics to the

rotating diamond tool assists the grinding by

reducing friction between the tool and the

workpiece, preventing clogging of the wheel

through ultrasonic cleaning action of the coolant and reducing the required grinding pressure. Disadvantages are primarily derived from

the limitations of weight of the diamond tool

(70 gram maximum) which can be effectively

coupled with the present 20 KHz sonic vibrator

and coupled horn.

In principle, the method lends itself to drilling, threading, grinding and milling operations.

Material limitations are those commonly con- nected with the use of diamond, although operation on carbides, tungsten, titanium, as well as boron carbide, has so far not been successful.

In general, application is limited to hard, brittle,

nonmetallic materials such as sapphire, ruby,

beryllium oxide, aluminas, glass, ferrites,

quartz, zirconium, pyroceram, lucalox and boron composites where chemical interaction with

the carbonaceous abrasive does not play a significant role.

No clear advantage has been developed for

the use of fluids other than water. The fluid

functions as a lubricant, coolant and cleaning

agent. Ultrasonics may add to the operation by

atomizing the fluids. Dry operation leads to appreciable heat generation resulting in premature tool failure. Fundamental knowledge of

his technique is still lacking, probably more so than for any other BAM method. Tool pressures

are said to be less than for other grinding operations. Tool loads range from 5—60 lb depending on tool type and operation. A force of

5 lb is considered typical on a 0.040 in core

drill. Surface finish data are not available for

direct and objective comparison. Finishes of

^32 rms with a 120 grit diamond tool are con- sidered standard.

Material removal rates and optimum machine

parameters have not as yet been determined.

Where applicable, the method is invariably

faster and performs the operation with less

chipping, tapering or break-through problems.

Figure 14 compares the feed rates in drilling 99

percent alumina employing RUM and with the

ultrasonic vibration turned off. As the figure

shows, the feed rates drop off as the drilling

depth increases for both methods although the

average drill rate appears twice as fast with

RUM for the conditions of this particular test.

17
I-3

CC

[ll Q.

CO

kJX

o

<

tr.

WITH ULTRASONIC

0 12 3 4 DEPTH OF HOLE 10"' INCHES

Figure 14. Comparison of performance data in drilling 99 percent

alumina with and without ultrasonic

assist. Data generated by the authors.

b. Free Abrasive Machining

Free abrasive machining is different from

bonded abrasive machining (standard wheel

grinding) or contained abrasive machining

(lapping) in that the abrasive is suspended

in a vehicle and held between a very hard

(Rc 60-62) horizontally rotating alloy steel

wheel and the workpiece. In this fashion the

abrasive grain is in contant surface contact

with the workpiece and is rolled over the wheel

surface without being embedded in it. It is an

inherently cool machining method with the

wheel reaching temperatures around 200° F.

The successful operation is predicated on the

fact of two favorable hardness ratios: the wheel

hardness to abrasive hardness ratio and the

abrasive hardness to workpiece hardness ratio.

The former is ideally held at a value of 1 :1 or

2:1 and primarily governs the degree to which

the abrasive is effectively rolling over the sur- face without embedding. The second ratio is

ideally regarded as having a 1.5:1 or higher

value and generally determines the degree to

which the abrasive can penetrate the workpiece and subsequently remove stock. The term

hardness is again used as expressing a combination of static physical characteristics such

as a combination of physical hardness and

toughness reflecting the degree of inpenetrability of a sharp point. A peculiar situation

exists in which on the one hand one desires as "hard" an abrasive as possible to remove stock

whereas the "hardness" of the abrasive at the

same time has to be limited to a value which

prevents embedding in the wheel to facilitate

the mechanism of free abrasive machining in

the first place.

Silicon carbide is generally used as the abrasive and most generally approaches the above

requirements for a wide variety of applications.

It is advantageous from an economic point of

view. Very little difference in performance, if

any, can be detected between different grades

of SiC (black or green). Boron carbide is effective on certain very hard materials such as tungsten carbide. The economic disadvantages

are offset by the possibility of recycling or re- covering the abrasive for further use. Aluminum oxide can be applied on soft materials but

with a much reduced stock removal ratio when

compared to SiC.

Maintaining the abrasive film is all-important. Several types of vehicles are used such

as diesel fuel oil of various viscosities as well as water. The type of vehicle depends on the operation. The abrasive film requirements limit

the grain size in an effective abrasive machining operation to approximately 5-10 ^ with subsequent limitations on surface finish. With very

small grain size the forces required to affect

penetration would be high enough to break

the grains down by shear or imbed them into

the rotating wheel. It is at this point that free

abrasive machining merges into a CAM operation in which abrasives of any particle size can

be employed.

The effectiveness of the operation is governed

primarily by speed, pressure, concentration and

size of the abrasive. The rotational speed of

the wheel varies with its diameter (3 in—87

rpm, 48 in—63 rpm, 64 in—45 rpm, 84 in

31 rpm) to give an average constant surface

speed. Higher speeds would mean higher stock

removal rates but speed is limited by the

properties of the vehicles (evaporation and

decrease in viscosity) and the centrifugal forces

which would remove the abrasive prematurely.

Pressure increases stock removal rates. Pressures of from 3 to 5 psi are most favorable, the

exact value depending on the material ground.

This is considerably higher than the pressure

in an average contained abrasive machining operation which is generally held at from V2-2 psi

but lower than the (dynamic) pressures in

bonded abrasive machining. This becomes an

important consideration when surface alterations due to pressure or temperature are ex- pected to have an adverse effect on the workpiece. Pressure is required to force the

abrasive into the workpiece causing deformation or brittle fracture as a result of the high

pressure in the actual contact zone. Subsequent

stock removal follows as a result of the rolling or cantilevering action of the abrasive grain.

Figure 15 top shows a generalized pressure

versus stock removal curve. The sharp drop

off after 5 lb/in is primarily due to premature

breakdown of the particles at pressures higher

than their friability limit. A certain minimum

concentration of abrasive is required to estab18
1 . I I I—.—I—'—I—'—

12 3 4 5 6 STOCK REMOVAL RATE (SRR)

Figure 15. Performance criteria for free abrasive

machining. Stock removal rate versus pressure

(top) and WGS/AGS ratio versus stock removal

rate (bottom).

lish the abrasive film before any cutting action

commences. Stock removal rate than increases

rapidly as concentration is increased up to a

maximum which in the absence of more detailed

information is presently stated as "1 lb of

abrasive per gallon of vehicle." Actual optimum

concentration should be a function of abrasive

particle density, a maximum being reached at

the point where free unrestrained rolling is

significantly limited by too high a particle density. The size of the abrasive determines the

stock removal rates. In general, a larger grain

size results in higher rates. This is not so much

due to size alone, but primarily in polycrystalline material, to an effective workpiece grain

size (WGS) to abrasive grain size (AGS) ratio

(see fig. 15 bottom). A small AGS can be con- sidered to be a small single point tool which

volumetrically speaking is not as effective in

stock removal as the larger grain which is able

to remove complete workpiece grains or aggregates rather than only small grain particles.

There is yet another reason demonstrated by

the abrasive action in this method on composite materials such as cemented tungsten carbide. Material removal does not become appreciable until the AGS is larger than the WC

particle spacing. Below this value the abrasive

grain becomes to a large extent embedded in

the softer (cobalt) matrix. The need for larger

particle sizes is probably not entirely governed

by this phenomenon but to a certain degree also

by the fact that small grain size high MOR

materials are harder to grind, requiring larger

particle sizes for appreciable removal rates. Because of the randomness of the lay of the abrasive and subsequent abrasive action, the surface

of FAM finished workpieces is invariably matte

or has a sandblast like appearance. In contained

abrasive machining a multidirectional-lay and

in bonded abrasive machining a roughly parallel lay are affected. Roughness is directly proportional to size, the larger sizes giving the

rougher surfaces, but is inversely proportional

to the hardness of the work, the harder surfaces

generally giving a lower reading for roughness

for any one abrasive size. This is directly re- lated to the extent of penetration of the abrasive particles. Stock removal rate is governed

by grain size as previously indicated, but primarily by the workpiece characteristics. Figure

16 shows the variation of this rate as a function of grain size. The materials with the lower

strength to modulus ratios (Modulus of Resilience) show the lower rates. In increasing or- der of favorable grindability and by group only,

the following general series is established:

Hard refractory elemental and polycrystalline

materials (WC, Ti), hard to soft steels, nonferrous materials, plastics, and finally the very

brittle or very soft inorganic materials such as

glass and ferrite. It may be noted that metallic

aluminum is essentially on par with the plastics,

that brittle plastics (phenol formaldehyde) perform better than the low modulus plastics, and

that the influence of AGS is most noted in the

multiphase materials with appreciable difference in physical properties between grain and

matrix (Alumina ceramics and WC). Members

of the same material type are primarily differentiated on the basis of hardness (see e.g., the

steels in figure 16). As may be seen from figure

16, free abrasive machining is not competitive,

on a stock removal basis, with bonded abrasive

machining, but may be conducted on almost any

material with various degrees of efficiency. It

is limited to flat surfaces and finds its best application when very large amounts of parts have

to be finished without appreciable stock removal

requirements, or where conventional grinding

and machining are impractical or impossible.

An example is on very thin workpieces where

magnetic or vacuum distortion could produce

warping. Of all FAM applications it is estimated that approximately 30 percent is done

on nonmetallic materials. In terms of material

ground, in decreasing overall quantities, it primarily handles cold rolled steel, non-ferrous

metals, cast iron, stainless steel, ceramics, glass,

19
lOO lOOr

10 20 30 40 50 60 GRAIN SIZE

10 20 30 40 50 60 (MICRON)

Figure 16. Stock removal rate versus abrasive grain size for various

material types ground in the FAM mode.

quartz, plastics, ferrites and boron carbide. The

biggest single research need in this area is for

further information on the mechanics of stock

removal.

c. Contained Abrasive Machining

This mode of stock removal accomplishes

grinding or polishing by holding, or generally

containing, an abrasive between tool and vi^orkpiece in such a way that the abrasive's movement is neither fully restricted as in BAM or

entirely free as in FAM. Limited movement due

to the elastic properties of the matrix in the

BAM mode is not of the magnitude considered

important for the above distinction. CAM is

different only in that it combines the features

of BAM and FAM. It cannot be directly equated

with lapping which merely connotes the use of

a flat revolving disk. Lapping per se can be

done in several modes. The use of the word lapping to indicate fine (abrasive grinding) finishing is entirely arbitrary with respect to

the degree of finish, minimum or maximum.

The action between a revolving disk, charged

with an abrasive, and a workpiece can be either

grinding or polishing. True polishing will ensue

only if conditions are such that no penetration

of the workpiece occurs and particulate matter

of macrodimensions is removed.

The CAM method generally entails the use

of an active element to which the abrasive is

applied. The abrasive is retained in the element by special vehicles (compounds), a re- tainer (felt, cloth, wire mesh) or by the mere

action of penetration of the abrasive into the

active element, usually a combination of the

above. The abrasive is retained but not rigidly

localized. Application modes may be divided as

20
INTERNAL DIAMETER

Figure 17. Subdivisions of the application modes in

contained abrasive machining (CAM).

shown in figure 17 into area contact, line contact, and point contact. It is understood that

similar operations may be carried out with the

principles of BAM. Contained abrasive machining is generally applied when finer finishes are

desired. The use of successively smaller grain

size abrasive leads to a natural transition from

the predominantly grinding action to ultimate

polishing. Abrasive requirements are allimportant. Ideal conditions are uniform size,

shape, strength and purity of the abrasive. Application of CAM to ceramics uses diamond almost exclusively as the abrasive. The active

element should be relatively hard. Ceramic or ceramic composite materials may be used successfully. This eliminates the possibility of con- tamination of the workpiece with material from

the active element. CAM is usually practiced in

the presence of a coolant-lubricant. Its desirable

action is in part dictated by chemical requirements such as possible detrimental effects on

the workpiece but is primarily governed by its

ability to effectively disperse the abrasive, act as a lubricant and coolant and keep the process

residue in suspension.

A diagrammatic overview of mechanical

stock removal processes is given in figure 9. The

BAM mode offers the highest stock removal

rates per part and consequently is the most

widely used method in ceramic finishing. Subdivision into types is influenced by the nature

of the products produced. Table 1 gives a breakdown of machine usage by type differentiated

on the basis of the nature of the operation.

Table 1. Machine usage by type and operation in

ceramic finishing

Machine usage percent

Type of Operation Specialty techni- General ceramcal ceramics ics operations

operations

Surface grinding

(vertical)

34 10

Cylindrical grinding 20 5 Centerless grinding

Surface grinding

(Horiz.-"Blanchard")

10

8

60

Universal grinding

(O.D.-I.D.,

radius)

7 15

Specialty grinding

Jig

Drilling

Slicing

4

16

1

10

The economic importance of the finishing step

in ceramic part processing becomes apparent

when it is realized that approximately 75 percent of the direct cost of the part is due to finishing in speciality technical ceramics. The

figures are 60 percent up to 90 percent in general ceramics manufacture for large and small

parts respectively.

3. Science and Theory of Ceramic

Finishing

3.1. General

In grinding, the kinetic energy of an abrasive

particle at the periphery of a rotating grinding

wheel is transferred to the workpiece and

dissipated in a variety of ways : kinetic energy

of the chips, difference in energy between the

old and the new surface, possible residual

energy of the material and the chips due to

distortion by the grinding process, and energy

required to remove material, and thermal energy. Quantitatively, the last two are the most

important.

Analysis of the mechanical stock removal

energy requires a knowledge of the forces

imparted to the workpiece under the dynamic

loading conditions of the grinding system. The

conditions under which such forces should be

determined must be characteristic of a typical grinding cycle. The grinding conditions to

- Figures are a result of a survey conducted in the ceramic

industry by the authors.

21

439-921 O - 72 - 3
which the workpiece is subjected creates material behavior aspects which are not normally

found under static loading conditions. The most

important aspects are the highly transient

localized nature of the loading, the dynamic

quality of the rapid or impulsive loading, and

the spatial and temporal distributions of the

load. Furthermore, the possibility for change of

the material properties in the relatively small

region to which the high pressure, rapid-rate

loading is applied, may influence the level and

distribution of the stresses. There are additional factors which set apart the analysis of

material behavior under grinding conditions.

The material reaction is influenced by its physical properties as well as those of the imparting

member (load). As will be seen later, this latter

determines to a large extent the efficiency of the

grinding process. Secondly, the impulsive loading imparts motion to the material upon which

it impacts and each loading cycle is followed by

a removal phase. Thirdly, as a result of geometric restrictions, the loadbearing areas at

first increase and then decrease providing cyclic

pressure dependence in the contact area. Finally, the highly complicated nature of the

impact phenomenon takes place at appreciably

elevated temperatures.

High temperatures are generated during the

contact time (tc) of load and workpiece. A

cooling cycle ensues during time (t, —tc) where

t,. is the time required for a full revolution of

the load. With increasing time, the temperature

will continue to rise and eventually, in the

absence of forced environmental cooling, attain

an equilibrium temperature the exact value of

which will be determined by the thermal properties of all system components. Heat will be

distributed primarily into the abrasive and

wheel, the workpiece, and the grinding residue.

This involves conduction convection and radiation mechanisms. Exact determination of the

extent of the heat flow into all three major

components would be required to draw a complete heat balance on the grinding process. An

approach to the solution of this problem is

presented in one of the following sections.

Primarily for the reasons outlined above,

initial fundamental knowledge on the dynamic

and thermal aspects of stock removal in ce- ramics can only be obtained when data are

collected using the most simple experimental

simulation, while at the same time retaining all operational aspects of the process. The authors'

experimental procedures and the results are presented in this report.

3.2. Experimental Design

Grinding action on ceramics has been simulated using a single point diamond "tool" of

varied geometry mounted on the periphery of a

Figure 18. Perspective drawing of single point

diamond wheel indicating diamond, thermocouple wires and slipring assembly mount.

7 in diam. one-half in thick alumina disk. See

figure 18. A basic grinding machine was modified to conduct the experiments. A special calibrated strain gauge type dynamometer designed similar to the original concepts by N. H.

Cook [17] recorded tangential and normal

grinding forces. A plastic housing encased the

entire grinding table enabling control of grinding environment. The collection of grinding

debris was by a specially designed tray. This

is shown in figure 19. The speed of the force re- corded (Sanborn Model 321) was sufficient to

register the force level at each pass of the

diamond through the workpiece. It was thus

possible to examine the variation of the force

level during each pass. Data indicate that for a

perfectly flat surface, the force level remains

constant during each pass across the workpiece.

The force level in both the normal and the tangential component decreases as a function of the

time required to impart the full tool geometry

to the workpiece. Even with a perfectly rigid

diamond, more than one pass was required to

impart the full diamond geometry to the workpiece. Under certain conditions, a continuous

grinding action of forty passes was observed.

Further details will be discussed in the following section.

The initial temperature measurements during the grinding action were made by means

of chromel-alumel thermocouples. One of the

thermocouples was mounted against the back of

the diamond body while the other thermocouple

was mounted as close to the tip of the diamond

as possible without interfering with the cutting

action. The thermocouple wires were brought

through the metal diamond housing, through

the groove in the wheel and the shaft and onto

a slipring assembly. With a known volume and

22
Figure 19. Two dimensional grinding dynamometer designed and used

for the studies reported here.

path length of the diamond, two thermocouples

provide the input for heat transfer calculations

and possible extrapolation to the actual temperatures at the interface between the diamond

and the workpiece. Infrared photography, microscopy, and liquid crystal targets were used

to obtain further temperature data.

The workpiece materials initially consisted of

four different conventional alumina types: 85,

94, 99.5, 99.9 percent alumina.^ They were

chosen to be distinctly different in terms of

mechanical as well as thermal properties. The

characteristics of the four workpieces are tabulated in table 2. All workpieces were preground to essential flatness and a finish of 10

microinches so as to ensure a constant depth of

cut during table traverse. Except for the 99.9

percent alumina, which had to be stacked, the

workpiece dimensions were 7 1/2 -in long, 2-in

'Source: Coors Porcelain Company, Golden, Colorado 80401. Certain commercial products and instruments are identified in

this paper in order to specify adequately the experimental procedure. In no case does such identification imply recommendation

or endorsement by the National Bureau of Standards, nor does it imply that the products or equipment identified are necessarily the best available for the purpose.

wide and 1-in thick. Materials from other

sources but with the same nominal composition

were used for comparison experiments.

In single point experiments and for a given

downfeed, tool geometry and workpiece material, the only variables in the process are wheel speed and table speed. The experimental

approach was such that output data would yield

force levels and temperatures as a function of

both wheel speed and table feed.

Observations on material behavior were carried out by reflected light microscopy. Use of

scanning electron microscopy has been initiated.

3.3. Results and Discussion of

Experimental Work

a. Force Measurements

Results of force measurements as described

in the previous section are shown in figures 20,

21, and 22 for a 90 degree cone angle diamond

cutting the four materials and for a 120 degree

diamond in the case of AD 85. Experiments on

AD 99.9 have in no case been carried beyond

23
Table 2. Mechanical and thermal properties of selected ceramics*

Alumina

X ropci by U ill to OO /C yy.o vo QQ OtCi yy.y

Specific gravity g/cc 3.42 3.62 3.84 3.96

Compressive strength psi 240,000 300,000 300,000 650,000

Tensile strength psi

70°F 18,000 27,000 -- 60,000

2000 °F 9,000 10,000

Modulus of elasticity psi 33X10° 41X10° 52X10° 56X10°

Shear modulus psi lo X lU 1 ( X lU ZZXIU OO w 1 A 6 2a X 10

Bulk modulus psi 20x10' 24X10" 30X10° 33.3X10°

Poisson's ratio 0.22 0.21 0.21 0.22

Hardness Rockell 75 78 81 90-91

45N

Color White White Pink White

Specific heat Cal/G/ °C 0.22 0.21 0.21 0.21

@ 100°C

Thermal conductivity Cal/CM V

20°C CM/Sec/ °C 0.035 0.043 0.075 0.0744

100°C 0.029 0.035 0.065 .0550

400 °C 0.016 0.017 0.028 .0409

800 °C 0.010 0.010 0.017 .0310

Thermal coefficient of Per °C

Expansion as noted

-200 to 25°C 2.3X10 3.4X10 -° 3.4x10 -° 3.6X10 -' 25 to 200 °C 5.4X10 6.3X10 6.8X10 -° 6.8X10^

200 to 500 °C 6.8X10 ^ 7.7X10-^ 8.0X10 -° 8.0X10-°

500 to 800 °C 8.3X10 8.5X10 -° 8.9X10-° 8.9X10 -° 800 to 1200 °C 9.0X10 9.2X10-° 9.7X10 -° 9.7X10-°

25 to 1000 °C 7.5X10 ^ 8.0X10 8.4X10-° 8.4X10 -°

* Compiled after data provided by Coors Porcelain Company, Golden Colorado on their products designated

with the nominal alumina percentages as indicated.

speeds of 2000 rpm due to premature failure of

the diamond. Initial impact forces are such that

physical integrity of the diamond is destroyed.

The grinding forces during the material re- moval phase are, however, well within the

capability of the diamond.

The results show that for any table speed,

the force level is directly related to the wheel

speed. The forces decrease rapidly as the wheel

speed increases. The table speed was found to

have a considerable influence on the force level

which increased as the table speed was increased. Increasing table speed is indicated in

the figures by an increase in number from six

through fourteen. It should be noted here that

the indicated forces are the sums of the forces

registered during each pass of the diamond

across the ceramic, until the full geometry of

the diamond, at a particular downfeed, had

been imparted to the workpiece and no further

forces could be measured. This is the only way

in which a total force for the energy required

to remove a given amount of stock from the

workpiece, could be successfully obtained. As

the amount of material removed per revolution

is increased by increasing the depth of cut, the

force level is raised. See figure 22.

If the ceramic grinding forces could be

viewed as being governed by strictly geometric

considerations, an analysis of the force levels

would be relatively simple. For an initial model

we could assume that the amount of material

removed would be dependent on the shape and

depth of penetration of the diamond and the

machine parameters only, i.e., no densification,

plastic deformation or extent of deformation

beyond the tool workpiece contact zone. Direct

visual observation of cutting grooves reveals

that this does not reflect the actual process. The

use of a strictly geometric model does, however,

assist in understanding some of the force variations as a function of changing grinding conditions. Let the number of revolutions of the

wheel per unit time equal V, the depth of cut

equal d, the table speed v, and the width of the

groove at the surface Wa. The volume removed

per unit time. A, will be

A = {vdWa)/2. (1)

and the volume per revolution of the wheel. A'

A' = {vd Wa)/2V (2)

The exact shape of the removed volume under

grinding conditions is complex and its exact

24
QO

5.0

4.5

4.0

as

O 3.0

<" 25

1500 2000 2500 3000 3500

WHEEL SPEED (R PM)

2.0

1.5

1.0

0.5

00

MATERIAL AD 94

TOOL 90° CUTTING FORCE Fn DEPTH OF CUT OOl" TABLE SPEED 6 to 14

14

12 \

10 \ \

8\\ \

6 Yy. \

oo

50

4.5

4.0

3.5

m30 O

w2.5

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1.5

1.0

05

00

1500 2000 2500 3000 3300

WHEEL SPEED (RPM)

MATERIAL AD 94

TOOL 90° CUTTING FORCE F| DEPTH OF CUT OOl" TABLE SPEED 6 to 14

WHEEL SPEED (RPM)

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

Figure 20. Force versus wheel speed for 85 and 9^ percent aluminas,

.001 in depth of cut, 90° angle.

length and area are not easily determined. It has been shown [18], in the case of sur- face grinding metals, that the length of the

grain path L for v < < F may be expressed as L = {Ddy- where D is the diameter of the

grinding wheel. If we assume the force to be

proportional to the cross-sectional area of the

removed material, then the relative single point

force (F) would be

F = (vdWa)/2VL (3)

or using L = {Ddy^

F = (v Wa){dDy/2V (4)

For a given diameter of the wheel, the relative

single-point force is therefore expected to de- crease as the wheel speed (V) increases and to

increase as the table speed and depth of cut are

increased. This is supported by the data in

figures 20 through 22. The increase in force

level when the cross sectional area, i.e., Wa, is

25
5.0

4.5

4.0

3,5

o 3lO

V) 2.5

UJo

£2.0

1.5

1.0

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4.0

3.5

MATERIAL AD 99-5 TOOL 90* CUTTING FORCE Fn DEPTH OF CCn' 001" TABLE SPEED 6 to 14

O 3.0

UJo

I

2.0

1.5

1.0

0.5

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1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

MATERIAL AO 99-9 TOOL 90' CUTTING FORCE Fn DEPTH OF CUT 001" TABLE SPEED 6 to 14

5.0

4.5

4.0

3.5

83.0

to 2.5

£2.0

1.5

1.0

05

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50

45

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3.5

MATERIAL AD 99-5

1 UUL 90* CUTTING FORCE Ft DEPTH OF CUT 001° TABLE SPEED 6 to 14

I4v

12

10 \

8

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

g3.0

2.5

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

iij o

cc£2.0

15

1.0

05

0.0

MATERIAL AD 99-9 TOOL 90° CUTTING FORCE Ft DEPTH OF CUT 001° TABLE SPEED 6 to 14

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

Figure 21. Force versus wheel speed for 99.5 and 99.9 percent alumina,

.001 in depth of cut, 90° cone angle.

increased such as in going from cone angle 90

degrees to 120 degrees, may readily be seen when comparing figures 20 and 21, and 22.

Most efficient grinding, in terms of expended

cutting force, is, therefore, done at high wheel

speeds, low table speed and depth of cut with

diamonds of small cross-sectional area. Even

the simplified geometric model is further complicated by the fact that feed and wheel speed

are not independent variables.

Consider a surface as shown in figure 23,

where the relative speeds are such that individual cuts are made per revolution of the wheel.

If the table speed is kept constant but the

revolutions per minute increased by a factor of

two, the result will be as indicated in the center

part of figure 23. The amount removed per pass

will have increased. The number of revolutions

per inch will be V/v. The more revolutions per

in, the more material will be removed per unit

length. In reality, an increase in table feed (v)

not only reduces the number of revolutions per

26
70

60

50

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u>30

111 o

(££20

10

0

70

60

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§40

[2 30 u^20

10

0

70

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MATERIAL AD 85

TOOL 120' CUTTING FORCE DEPTH OF CUT •001" TABLE SPEED 6 to 14

14

"'^^^^

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

MATERIAL AD 85 TOOL 120° CUTTING FORCE Fn DEPTH OF CUT 002" TABLE SPEED 6 10 14

14.

12 \

.10 \\

8 \\\

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1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

CO

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£

20

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MATERIAL AD 85 TOOL 120° CUTTING FORCE Fn DEPTH OF CUT 003" TABLE SPEED 6 to 14

"1550 266023003556~3300'

70

60

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£ 20

10

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MATERIAL AO 85

TOOL 120° CUTTING FORCE F, DEPTH OF CUT 001" TABLE SPEED 6 to 14

1412^-^

6

tM 40 o

(O 30

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£ 20

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20

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WHEEL SPEED (RPM)

MATERIAL TOOL

CUTTING FORCE DEPTH OF CUT TABLE SPEED

AO 85

120°

Ft002!' 6 to 14

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

MATERIAL AD 85 TOOL 120° CUTTING FORCE Ft DEPTH OF CUT 003" TABLE SPEED 6 to 14

1^ 2obo zioo 3600 iSoo'

Figure 22. Force versus wheel speed for 85 percent alumina, three depths

of cut, 120° cone angle.

unit length, but also increases the length of the

cut path as shown in the lower part of figure

23. There will thus be a table feed at which the

number of passes will level off. This tendency is

apparent from figure 24 which shows the dependence of the number of passes required to

remove a unit volume of stock as a function of

the various grinding parameters. This behavior

is characteristic of the single-point grinding

action only, since in multipoint cutting tools

(wheels) the path of one diamond is followed

by many others. The action of the single-point as a function of the various grinding parameters thus influences the nature of the

generated surface. An increase in depth of cut

also increases the length of the cut, which for a given wheel diameter was found to be proportional to \/d. The total amount of revolutions

per inch and hence the amount of material re- moved per pass is thus proportional to V/v^d,

27
WHEEL SPEED V rev./min TABLE SPEED v in/min

"A"

WHEEL SPEED V, rev/min TABLE SPEED v in/min

(A)

WHEEL SPEED .V rev/min TABLE SPEED v, in/min

V, > V

(B)

Figure 23. Diagrammatic representation

of the amount of material removed per

pass as influenced by increase in wheel

speed (top two) and increase in table

speed (bottom).

or the number of passes required N ccV\/d/V.

These findings are borne out by the experimental results as shown in figure 24. Figure 24

may indicate a trend in showing an increased

number of passes required, when going from

low to high modulus materials. The nature of

the fracture phenomena must influence the ac- tual amount of material removed per pass.

Although no two ceramics show identical

force patterns, there are several important

common features. A sharp increase in force

level occurs at low speeds. The speed at which

the sharp increase in slope occurs is primarily

governed by the type of material and secondarily by factors influencing the total amount of

material removed per unit time, namely table

speed and depth of cut. For the materials

examined here, the inflection point lies at lower

speeds the lower their respective (fracture)

strengths. Figure 25 shows idealized forcespeed plots for two hypothetical materials.

Material removal forces are higher for the

100 160 220 280 340 400 100 160 220 280 340 400 100 160 220 280 340 400

TABLE SPEED (in/min) TABLE SPEED (in/min) TABLE SPEED (in/min)

100 160 220 280 340 400 100 160 220 280 340 400 100 160 220 280 340 400

TABLE SPEED (in/min) TABLE SPEED (in/min) TABLE SPEED (in/mir)

Figure 24. Number of passes required to remove unit volume of stock as a function of

table feed for various wheel speeds at three depths of cut.

28
WHEEL SPEED (rpm)

Figure 25. Idealized grinding force versus wheel speed plots for two

hypothetical ceramics.

stronger material although surprisingly little

difference was found between the various types

of aluminas. A sharper differentiation between

materials can be found at lower speeds where

the material type more drastically influences

the rate of increase. Based on the position of

the inflection point, a cutoff speed for lower

force grinding can be indicated which differs

for each material. As figure 22 indicates, the

volume parameters (feed and depth of cut) also

shift the inflection point. This is not shown in

the idealized curves of figure 25. Tangential

and normal forces follow the same trend. The

magnitude of the forces as well as the ratio of

tangential to normal force is primarily influenced by the geometry of the cutting diamond. The tangential forces are higher than

the normal forces for cone angles of 90 degrees

or lower whereas the value F,/F„ moves closer

to unity at higher cone angles. Experiments

with cone angles greater than 120 degrees were

not conducted. The magnitude of the tangential

force may be regarded as more directly reflecting the amount of energy required to remove

the material. The value of the normal force

component is particularly important with re- gard to the extent of surface or subsurface

damage. In actual cutting experiments the

diamond is only weakly supported by the fracturing workpiece so that the bulk of the normal

force field is concentrated at the small area of

direct contact between diamond tip and the

newly generated surface of the workpiece. Very

high stress fields are thus set up. At sufficiently

high pressures the contact area becomes plastic

and leads as a result of permanent deformation

to residual stresses. The initially high compressive stresses lead on relaxation to substantial

tensile strains which may cause subsurface

cracking. Recent studies of fracture phenomena in statically and dynamically loaded sodalime glasses by Peter [19] have shown the

'development of in-line subsurface cracking

("tiefenriss"). Koepke [4] observed subsurface

cracking in severely workhardened surface

layers when grinding MgO single crystals with

alumina wheels. Finely polished ceramic sur- faces often reveal small surface cracks on

profilometric traces. Peter [19] further indicated relaxation times for the development of

subsurface cracking of the order of days. This

phenomena has to our knowledge never been

studied in ceramics. It certainly deserves attention particularly in glass-phase bonded ce- ramics. The extent of subsurface cracks in- creases as a function of both pressure and

velocity of the diamond.

The foregoing discussions provide qualitative

insight into the dynamic process of diamondworkpiece interaction, based on geometric considerations. It leaves, however, some very important areas uncovered. The force expressions

developed thus far do not take into account the

specific differences created by the physical

properties of the workpiece. The (linear) dependency of the parameters shown in eq (4) is

not borne out by experimental observations.

This is in part due to the nature of failure of

the ceramic. As may be easily observed, ce- ramic failure is at least in part effected by side

cracking or divergent fracturing which as a

result of propagation and interaction ejects

flake like particles from the surface. This

(mechanical) spalling effect can extend in dimension to up to three or four times the width

of the diamond at the workpiece surface. Nature and extend depend on material type. Furthermore, rigidly mounted diamonds fail on

grinding high modulus aluminas in spite of

tangential and normal stock removal forces of

the same order as those on lower modulus

materials.

In ceramic grinding the impact phase creates

the initial fracturing which is followed by the

removal phase for which the energies are determined by the work required to propagate

and remove the material which has failed in a

brittle or ductile fashion or both. In grinding,

the propagation of the initial impact fracturing

is aided by the extreme forces which continue

to work on it as the wheel revolves. The work

to extend the fracture will be larger and the

velocity of propagation less for ductile failure.

The magnitude, durations and distribution of

the forces that are set up will be determined by

the way the diamond and the ceramic react on

one another. Fracture formation within the

material which is intimately associated with its

mechanical properties will determine the kind

of load imparted to the diamond. The main

areas of interest are, therefore, the distribution

of forces at the impact interface as a function

of time, and the way in which the load distributes itself within both ceramic and diamond.

29
High velocities or increasing strain-rates generally raise the flow stress of a material so that

in a high speed grinding operation on already

brittle ceramics, little plastic deformation is

expected to precede fracturing.

In high velocity high pressure impact experiments (such as in shock waves) use is made of

the one dimensional mass and momentum con- servation relationships [20],

and

V = V,{\-U,/Us)

P = Po + poUpUs

(5)

(6)

to calculate pressure and specific volume behind

the shock front from the experimentally determined Shockwave velocity (Us) and the "particle velocity" or bulk material velocity (Up)

behind the shock front.

A first order approximation of the normal

fracture stress occurring in the region of

fracture at the moment of impact may be

obtained from eq (6). Equating the particle

velocity with the wheel velocity at 6000 ft/min

(3x10^ cm/s) and taking the average of alumina density at 4 g/cm^ and the longitudinal

sound speed at 1 x 10® cm/s, the fracture stress

computes at 12x10" dyne/cm^ (12 kbar). Eq

(6) strictly holds only for infinite bodies which

do not sustain shear and both elastic and inelastic processes should be considered. Wear or

deterioration of the diamond will occur if the

prevailing stress level falls above its yield

strength. Accurate data on the strength of diamond is scarce and unsatisfactory. Careful

measurements by Hull and Malloy [21] come to

the conclusion that, if extraordinary precautions are taken to avoid stress raisers, the

failure shear stresses in clear natural diamond

will fall around 43,000 jssi and a compressive

strength of 1.261x10® psi is indicated. The latter value should be compared with data from

Schwarzkopf and Kieffer [22] which give a compressive strength of 284x10^ psi. Their

transverse rupture strength falls at 42,6x10^

psi. Considering the above figures, diamond will

fail in compression between 20 and 100 kbar in

shear, and will transverse rupture at about 3

kilobars. If fluxural strength is taken as an

indication of failure and a value of one sixth

the compressive strength is assumed, the fail- ure limit must be set at 13 kbar. It is doubtful

that the theoretical (hydrostatic) compressive

strength of diamond will be the governing

factor in the grinding mode. It may, therefore,

be concluded that the stress levels of a rigid

diamond impacting on the hardest of the aluminas at speeds of 100 ft/s or up is in the

neighborhood of the critical failure stress.

Experimental observations indicate premature

failure on the ceramic with the highest density

and highest sound velocity, namely the 99.9

percent alumina. For any given force, size,

shape, and orientation, the rate of loading will

influence the stress experienced by the diamond.

The degree of support or energy absorption

characteristic of the bond will thus largely

determine the erosion characteristics of the

diamond at impact on any one type of ceramic.

These bond characteristics will, for optimal

grinding conditions, usually have to be adjusted

for each type of ceramic. Single point diamond

grinding tests may be a good way to determine

the failure strength of diamond under dynamic

loading conditions.

The effect which the transient disturbance of

the impacting diamond has on the workpiece is

largely determined by its physical characteristics. At least the qualitative aspects of the

behavior may be analyzed from the point of

view of an elastic approximation.

A knowledge of the extent of the damaged

area per unit kinetic energy may be obtained

from the following simplified model. At impact

the kinetic energy per unit area, Ejt will be

E, =

~2A

where m = diamond mass

A = impacting area

V = velocity of the wheel

Energy consumption will be the product of the

energy density and the extent of the damage

Ec = creS (8)

where a = stress

c = strain

S = depth of damage

At constant stress, the value of the strain at

any point into the solid may be expressed as

[23],

o-tx

) (9)

where x = -q/E, the relaxation time and rf characterizes the viscosity of the deformation

process and E is the elastic modulus. Substituting eq (9) into eq (8) and equating with eq

(7) leads to

E

(1

assuming o- = <ti, or

^*^) S =

E

'2A

2A crMl - e-*^)

(10)

The extent of the cracking can thus be re- garded as a product of a term expressing the

level of the impacted energy as determined by

systems parameters, and an energy density

term reflecting the strength of the material or

30
Figure 26. Schematic presentation

of the extent of fracturing at impact at increasing wheel

speeds for a hypothetical aluthe resistance to failure. The second term combines the deformation impedance and yield

strength characteristics of the material into a dynamic damage resistance parameter. The

extent of the damage for materials with very

small relaxation times will be inversely related

to their modulus of resilience (MOR = 2E ).

A minimum stress (o-o) is required to propagate

a crack of length Co according to the Griffith

theory,

^ 2 -

>'

(To — —E

Co

where y = surface energy of the material.

Substitution into eq (10) leads to,

S = mV^ Co

2A (1 -tx

) (11)

mina.

The smaller the initial crack length (grain size)

and relaxation time and the larger the surface

(fracture) energy of the material, the more

limited the extent of the damage at impact

becomes. Among the investigated materials

these properties are most enhanced in the 99.9

percent alumina. Figure 26 gives a schematic

1500 2000 2500 3000 3500

WHEEL SPEED (RPM) 1500 2000 2500 3000 3500WHEEL SPEED (RPM)

1500 2000 2500 3000 MOO

WHEEL SPEED (RPMT^

1500 2000 2S00 3000

WHEEL SPEED (RPM)

Figure 27. Temperature recorded by tip and back thermocouples for various ceramics as a function of wheel speed,

0.000 in depth of cut, 90° cone angle and several table feeds.

31
interpretation of the increase in damage at

impact as a function of increasing wheel speed

for a hypothetical material. Forces required to

remove the entire volume at the lowest speeds

(top fig. 26) would be the highest as a result of

the need to induce crack propagation during

completion of the interaction cycle.

b. Temperature Measurements

Figure 27 shows the results of grinding

temperature measurements. The recorded re- sults are the highest temperatures observed

during each experiment conducted. They are intended to show the temperature trends as a function of the various grinding parameters.

An increase in temperature is obtained with an

increase in wheel speed and a decrease of the

table speed. The temperature is also greatly

influenced by the type of material ground. The

highest temperatures were recorded on the 99.9

material. Temperatures increased considerably

when depth of cut was increased, with an in- crease of about 30 degrees per mil in the case of

99.5 percent alumina. The temperature trends

can be explained by the thermal model as discussed below. The temperature increases as a function of speed, appear to contradict the in- verse relationship shown for the force levels

(see fig. 22). The explanation may be found in

the fact that the temperatures indicated are

those generated during each grinding event for

a given set of grinding parameters. As such,

the temperatures will be directly related to the

number of passes required to remove the total

volume from the groove. As may be seen from

figure 23, higher speeds required fewer passes

and, therefore, a greater amount of heat was

dissipated per pass resulting in higher temperatures. The same analogy may be used in

evaluating the influence of the table feed. The

higher the table feed, the greater the number

of passes and subsequently the lower the temperature per cutting event. The temperature,

therefore, is directly proportional to the amount

of energy expended per unit time.

The grinding temperature generated at the

diamond workpiece interface will be dissipated

to the wheel, the ceramic, and the chips. A

knowledge of heat flow distribution is as much

needed as the magnitude of the grinding temperature itself since it determines how and to

what extent the various system components will

be thermally affected. The nature of the grinding process necessitates an analysis in terms

of an intermittent transient (unsteady) heat

transfer model. Only the essential concepts can

be given here. Further details may be found

elsewhere [24].

The general heat conduction equation which

governs the temperature distribution and the

conduction heat flow in a solid, without internal

heat sources,

a 8(9 (12)

where, a = k/c,„ = thermal conductivity,

c„ = specific heat and p = density, and T and

0 indicate temperature and time respectively,

has for the case of a one-dimensional transient

heat conduction in a semi-infinite solid a solution according to the following equation

T - T„ X

To-T^ - ^ 2V^ ^^^^

where (T — T^) = temperature difference at

any one point and {To -T^) = initial potential

difference x = distance into the surface. The

quantity on the right hand side of the equality

is known as the Gaussian error integral. The

rate of heat flow at the surface, q, may be obtained from eq (13), by evaluating the temperature gradient at the surface. The total change

of internal energy or heat flow, Q, may then

be found by integration

Q

9

qde

from which

Q = 2kA(T^ - To)\na

(14)

(15)

and

Q. =

e 2A "p'

1/2 (To-TJ (16)

Vn ^

Setting T = (To - T^),t'he heat flow into the

workpiece material will be

1/2 Cnm \ 2A

/ Vn

and similarly into the wheel

(17)

_ _rp / KmpmCpd \

'^'^ 2A dg)

where subscripts m, d, and w refer to material,

diamond, and wheel respectively. Heat ex- change to the grinding chips is related to the

mass rate of removal (m), the heat capacity

{€„„,) and the temperature through

Qr = mCpmT (19)

Mass rate of removal for a 120° diamond cone

and table speed V will be,

m = tan60p„,y (20)

The area of contact is simply given by the half

cone area at a particular depth of cut.

Using eqs (17), (18), and (19) and multiplying by the factor

32
V:^ Q 1/2 2AT '

•^1 — {K-mPmCpm)^^

Fz = (KdpdCpd) 1/2

pm

Vn

2A

a 1/2

(21)

(22)

(23)

The contact time follows from

Ow —

* 60

360 N

where N = revolutions per minute of the wheel

and * is the number of degrees of arc over

which the diamond is in contact with the workpiece. The value of the latter quantity can be

obtained from a rather complex formulation

[18]. It varies generally between three and five

degrees for 0.001 in and 0.003 in depth of cut.

Contact times range subsequently from 1 x 10-^

to 6x10-3 seconds. Eqs (21), (22) and (23) afford determination of the heat flow ratios such

as

Qt

and similarly for — and —

Qt Qt

where

Qt = Qm + Qc + Qw

The results were obtained with the aid of a

small computer program and are shown in fig- ure 28. Material differentiation is primarily on

the basis of the thermal properties. The largest

amount of heat goes into the wheel followed

by chips and workpiece material. From a thermal point of view, diamond grinding ceramics

is a gentle operation. Cooling of the wheel as a

whole, not only at the grinding interface, is

bound to result in very low thermal effects on

the workpiece. This is particularly true at the

more efficient higher wheelspeeds, because of

the steep slopes indicated in the upper right

hand graph of figure 28.

On account of process limitations the thermocouple mounted on the cone part of the diamond

is a finite distance away from the actual grind27-

25-

23

17

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

1500 2000 2500 3000 3500

WHEEL SPEED (RPM)

MACHINE SETTING WHEEL SPEED — 2500 RPM

TABLE SPEED — 9 DEPTH OF CUT - 003 INCHES TOOL (DIAMOND) - O

% DISTRIBUTION OF HEAT INTO WORKPIECE (Q^/Oj)

% DISTRIBUTION OF HEAT INTO CHIP (Qc/Qj)

% DISTRIBUTION OF HEAT INTO WHEEL (Qw/Qy)

Figure 28. Percentage heat flow into the major components of a ceramic finishing system.

33

I
ing interface. Using a steady state heat con- duction model [24] it has been possible to

find an expression relating the measured temperature {To) to the grinding temperature To

To =

1.99

a

(25)

where "a" varies in value from 0.79 to 0.94 for

M^heelspeeds of 1500 and 3500 rpm respectively.

The model takes care of heat losses by conduction and convection during the time that the

vi'heel revolves in air. As a result the temperature at the single point diamond-ceramic grinding interface on AD 99.9 at 1500 rpm, 0.001 in

depth of cut with a cone angle of 90°, is estimated at 275° F. (See also fig. 27).

Attempts at determining the swarf temperature directly by means of a small field of view

(0.010 in) infrared microscope have not led to

quantitative results. Conversion of the millivolt reading of the infrared instrument is predicated on the field of view being entirely filled

with the source to be measured. During grinding the percentage filling of the field of view is

a constantly changing function of time. Qualitatively we have been able to determine relative

temperature variations in the area of the diamond tip and its surroundings by training the

center of the field of view on various X-Y co- ordinates. The results are shown in figure 29.

The highest temperatures (highest numbers)

are found directly in front of the moving diamond.

If the temperature of a chip in flight can be

determined, the temperature at the point of

- INFRARED MICROSCOPE

EXPERIMENTAL SET-UP

NO. LOCATION MV REMARKS

1 1.2 33

2 2, 3

3 2,4 3.2 4 2,5

5 6, 1 1.2 6 6, 3 7 6,4 2.2 8 6, 5 2.2

Figure 29. Relative grinding temperatures in

the area surrounding the diamond as obtained by infrared microscopy.

origin can be calculated. Particulars are given

in another communication [24] . Thermally sensitive targets such as surfaces coated with

liquid crystals have been used. The thermal

mass of the particles is generally too small to

create areas of color change large enough to

be photographed. Targets which leave a permanent record are presently being investigated.

Infrared photography has been applied and

works in principle. Calibration of the film colors

is hampered by a curious phenomena. It has

been observed during grinding tests in different

gaseous atmospheres that the color of the

"sparks" is different from one environment to

another. It is felt, therefore, that the color of

the sparks is partially thermal and partially

due to gaseous discharge as a result of high

electric fields during the rubbing action of the

two contacting dielectics. Special filtering will

be required to separate the various components.

4. Conclusions

It is apparent that considerable progress has

been made in recent years in the science of

ceramic finishing. One of the fundamental objectives of this paper was to identify the state

of the art in ceramic machining. A total systems

approach was undertaken to analyze the mechanics of the process, determine the behavior

of the ceramic material during grinding, and

identify the thermodynamic and mechanical

parameters. An understanding of the machine

system, which delivers the motion and energy,

is necessary to clarify the overall grinding process.Some major findings are: • Mechanical stock removal can be classified

into three categories: Bonded Abrasive

Machining (BAM), Free Abrasive Matching (FAM), and Contained Abrasive

Machining (CAM).

• Cutting forces that are generated between

diamond (tool) and material (workpiece)

are directly related to the impact energy,

rigidity of the system, and physical properties of the workpiece. • Thermodynamic properties of the interacting materials are very important in the

grinding process. A thermodynamic model

shows that the greatest percent of the heat

is absorbed in the wheel, and the remaining heat flow is distributed between the

chips and the workpiece. • Machine parameters are an important consideration in determining the optimal

grinding process. Wheel speed generally

had the greatest influence on the process

behavior. Table speed has a lesser effect,

while depth of cut primarily determines

diamond life. • Studies have indicated that environment

34
has an effect on the cutting- force. A water

mist has the largest effect, and reduces

cutting forces when compared to grinding

in oxygen, nitrogen, air, and compressed

air environments. • In general, high volume of material re- moved combined with low cutting speeds

requires disproportionately large cutting

forces and reduces diamond life. With a

decrease in depth of cut and an increase

in cutting speed, the diamond has a long

life time.

In general, continued emphasis should be

directed to the mechanics of stock removal in

all three modes of abrasive machining. Specifically, in the Free Abrasive Machining (FAM)

mode, scientifically justifiable information appears to be totally lacking. Quantitative information on surface and subsurface damage,

while grinding, needs to be developed and ex- pressed in terms of measurable and controllable

parameters.

One of the most significant indications derived from this research is the necessity for inprocess control. In-process feedback control,

will emerge as an absolute necessity to repeatably produce finished ceramics with characteristics specified prior to the finishing cycle.

5. References

[1] Gielisse, P.J., Mathewson, W.F., Martis, J.A., and

Ratterman, E., Ceramic Finishing with Diamond, Ceramic Industry, Feb., March, April

issues (1968).

[2] Sedlacek, R., Processing of Ceramics—Surface

Finishing Studies, Final Technical Report,

Contract #N00019-67-0494, November (1968).

[3] Gross, G.E., Physical Parameters Affecting

Fracture Strength and Fracture Mechanisms

in Ceramics, Technical Report, Contract

#N00019-68-C-0127, December (1968).

[41 Koepke, E.G., and Stokes, R.J., A Study of

Grinding Damage in Magnesium Oxide Single

Crystals, Technical Report, Contract #N00014-

69-C-0123, July (1969).

[5] Rice, R.W., Machining and surface workhardening of MgO, to be published.

[61 Cutter, I., and McPherson, R., Examination of

abraded MgO by x-ray diffraction line broad-

[7

[8

[9

no

[11

[12

[13

[14

[15

[16

[17

[18

[19

[20

[21

[22

[23

[24

ening, Phil. Mag., 20 (165), 489-94 (1969).

Harrison, W.B., Influence of surface condition on

the strength of polycrystalline magnesium

oxide, J. Am. Ceram. Soc, 47 (11) 574-79

(1964).

Evans, A.G., and Davidge, R.W., The strength

and fracture of fully dense polycrystalline

magnesium oxide, Phil. Mag., 20 (164), 373-88

(1969) . Hockey, B.J., Observation of plastic deformation

in alumina due to mechanical abrasion, Am.

Ceram. Soc. Bull. 49, (4) 498 (1970). idem

Plastic Deformation of Aluminum Oxide by

Indentation and Abrasion, J. Am. Ceram. Soc,

to be nublished.

Rice, R.W., personal communication.

Hara, A., Megata, M. and Yazu, S., X-ray study

of residual stresses in ground cemented carbides, Powder Metallurgy International, 2 (2),

43-47 (1970).

Sedlacek, R., Processing of Ceramics—Surface

Finishing Studies, Final Technical Report

Contract #N00019-67-C-0494 (1968).

McKinney, K.R., and Herbert, CM., The Effect

of Surface Finishing on Structural—Ceramic

Failure, NRL Report 6942, Naval Research

Laboratory, Washington, D.C.

Sedlacek, R., Processing of Ceramics—Surface

Finishing Studies. Final Technical Report,

#N00019-68-C-0388 (1969).

Gielisse, P.J. and Stanislao, J., this paper.

Gielisse, P.J., and Stanislao, J., Abrasion, in Encyclopedia of the Solid State (Marcel Dekker,

Inc., New York, New York (1970).

Cook, N.H., Loewen, E.G., and Shaw, M.C.,

Machine tool dynamometers, American Machinist 98, 125-129 (1954).

See for instance Mathews, S , Thermal Aspects

Related to Stock Removal of Ceramic Systems,

M.S. Thesis, University of Rhode Island

(1970) . Peter, K., Sproedbruch und Mikroplastizitaet von

Glas in Eindrucksversuchen, Glastechnische

Berichten 37, (7), 333-345 (1964).

Duvall, G.E., and Fowles, G.R., Shockwaves in High Pressure Physics II, R. S. Bradley ed. (Academic Press. New York, New York, 1963).

Hull, E. H., and Malloy, G. T., The strength of

diamond, Presentation Annual Meeting American Society of Mechanical Engineers, Chicago,

Illinois (1965).

Schwarzkopf, P., and Kieffer, R., Cemented Carbides, p. 236 (MacMillan Book Company, New

York, New York, 1960).

Steverding, B., Brittleness and impact resistance,

J. Am. Ceram. Soc. 52, (3), 133-136 (1969).

Gielisse, P.J., and Stanislao, J., A Thermal Model

for Ceramic Grinding Temperatures, to be

published.

Discussion

O'Hara: You mentioned that 70 percent of

the heat goes into the workpiece rather than

the tool itself ; now that's primarily governed

by the fact that the tool was a higher thermal

conducting material than the workpiece. Is that

not true ? Gielisse: That's correct. As a matter of fact

our wheel is essentially our diamond and the

large difference in the thermal conductivity

and heat transfer processes is, of course, due to

the very thing you indicated.

O'Hara: In this case you are using a single

crystal. Depending upon the bond, heat transfer

conditions may change if you are using an ac- tual wheel having a composite structure.

Gielisse: That's right, and specifically so when you use a poljoner bond.

35
NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).
Direct Observation of Material Removal Process During Grinding of Ceramics

by Micro-Flash Technique

Osamu Imanaka, Seiji Fujino, and Shin-ei Mineta ^

Electrotechnical Laboratory, Tokyo, Japan

In order to make the mechanism of grinding ceramics clear, chip removal processes have been observed during grinding with a specially-built setup which consists

of a flash lamp, synchronizer, microscope and motor-driven camera. Duration time of

the flash is about 0.8 fis in half-value width, and a storage capacitor in the flash lamp

circuit can store an energy of about 10 W-s. By choosing the flashing point, grinding

phenomena at diff'erent positions on the contact arc can be photographed at certain

time intervals.

The experiments were done mainly by using one grain of diamond attached to a

disk which was mounted on a shaft of a surface grinding machine. Workpiece

materials examined were glass-ceramics and several types of oxide ceramics, e.g., alumina, forsterite, and steatite ceramics. The main results obtained are as follows:

Under moderately large depth of cut, the chips generated from ceramics are generally

of a fragment type and are distinguished from those of ductile materials. Part of

these chips were observed to be removed, of course, directly by the grain, but a con- siderable part of them were found to splinter out of the workpiece surface in a short

time after the grain has passed over. Possible explanations are proposed for those

phenomena. The above-mentioned facts suggest that a damaged layer may remain

under the ground surface of ceramic materials. It is also demonstrated that chips of

continuous type are occasionally observed at high grinding speed and under moderately

small depth of cut. The chip formation of this type depends delicately on the materials

to be ground.

Keywords: Alumina; brittle fracture; chip removal process; forsterite; glass-ceramics;

grinding of ceramics; micro-flash technique; observation of chips during grinding;

steatite.

1. Introduction

In research into the cutting mechanism of

abrasive grains on the grinding wheel, direct

observation of chip formation during grinding

is thought to be one of the most effective approach to the object. In metal grinding,

Schv^artz - photographed the surface of a grinding v^heel at certain intervals during grinding

with a specially-built equipment, which consisted of a substandard film camera, a stroboscopic light, suitable synchronization setup, and

optical compensation. However, little has been

reported concerning the cutting behavior of the

abrasive grain during grinding, and the information has been limited to the occurrence on

the peripheral surface of the grinding wheel.

The purpose of this study was to observe the

chip removal processes directly during grinding and to provide with new insights on the

cutting performance of abrasive grains in ce- ramics grinding. Photographs of grinding chip

removal from different materials taken by

means of a synchronized flash lamp are presented and discussed.

2. Experimental Setup and Procedure

The experiments were done mainly by using

one grain of diamond attached to a aluminum

' Chief, Senior Scientist, Junior Scientist of Fabrication Technology Section, respectively.

2 Schwartz, K.E., Industrie-Anzeiger 80, 1391 (1958).

disk of 200 mm diam, which was mounted on

a shaft of a surface grinder as shown in figure

1. The diamond grain used had the shape of a truncated quadrangular pyramid of 90° facing

angle. Diamonds were chosen with edges between 20 and 30 ^^m to simulate flat worn sur- faces on actual grinding wheels. Grinding was

done by pyramidal face (not by an edge).

Workpiece materials examined are shown in

table 1. Before an experiment, the surface of

the specimen was finished to a roughness of

less than 1 jjm maximum height by lapping

with silicon carbide abrasives of 2000 mesh

grit size.

Table 1. Workpiece materials used

Workpiece materials

Microhardness, Hv

{kg/mm

Steatite ceramics (77% 674

MgO-SiO)

Forsterite ceramics (97% 773

2MgO-SiO)

Alumina ceramics (94% AI2O3) 1860

Insulator porcelain 651

Glass-ceramics 685

Synthetic sapphire 2130

Bearing steel, SUJ 2 (hardened) 623

439-921 0-72-4 37
Figure 1. Experimental set-up. A, abrasive grain; C, camera; L,

synchronized flash lamp ; M, microscope ; W, workpiece.

The grinding speed, U, and the table speed, v, adopted in the experiments were 12.5 to 37.5

m/s and 0.042 to 0.125 m/s, respectively. The

depth of cut, g, was chosen to be 3 to 4 fim

except the case shown in figure 15. For the con- venience of observation, all the experiments

were carried out without supplying grinding

fluid.

By means of a parabolic reflector and Fresnel

lens, the light beam from the flash lamp is con- verged on the grinding part to be photographed.

In order to photograph the chip removal process during the grinding at a specific, chosen

position of the grain (refer to fig. 4), a synchronized illumination system shown in figure

2 was adopted. When the shutter of the camera

is opened, the flash lamp lights up in agreement

with a preset angular position of the disc.

Namely, the beam reflecting from the synchronizing mark on the disk comes to the synchronizer, and makes the photocell emit an electric

pulse which is amplified and gives the signal

to trigger the flash. The synchronization is good

to within 0.2 fis at each time.

The light-emitting spark gap of the flash

lamp is enclosed by a fused quartz cylinder

which is filled by high-pressure argon gas. Input voltage to the flash lamp is 20kV and the

built-in storage capacitor can store an energy

of about 10 W • s. Restoring charges of the ca- pacitor needs 10 s after the preceding flash

has finished.

The duration of a light pulse of the flash was

measured at about 0.8 /xs in half-value width

as shown in figure 3. Provided that an exposure

time to the film is equal to 0.8 /^s and that the

Gate circuit.

synchronizer

of photo- conductive tvoe

Synchronizing mark

X-contaot

of camera

?la3:i la.Tip

oulse transformer -Lnw

\—

w

-11- O.C25l-i5''*VQl

II o

' 0.005-1?

SyncliTO- trigger pulse circuit

Figure 2. Schematic arrangement of illumination system.

38
Figure 3. Oscillogram showing the duration of a light

pulse. Time scale : 1 /is/div.

deb

Figure 4. Flashing point.

grinding speed of the grain is 25 m/s, the movement of the grain is 0.02 mm in the time of ex- posure. When a blur on the film of 0.2 mm is

allowed, photographs can be taken at magnification of 10 diam without the picture's quality

being noticeably degraded. A flashing point

(fig. 4) can be altered by moving the position of

the beam emitter of the synchronizer or by

using a time-delay device which presets a pulsedelay between an input pulse and triggering

of the flash.

3. Results

A series of photographs taken at different

flashing points during grinding forsterite ce- ramics are shown in figure 5, where the photographs (a), (b), (c), (d), and (e) were re- corded at the positions a, b, c, d, and e in figure

4, respectively. On the right side of each photograph, is arranged a schematic sketch of chips.

Each photograph was taken approximately 10

s after the preceding one had been done. (The

same procedure was used in figs. 6 through 9.)

Hence the chips in different photographs are

not identical. However, since the wear of the

diamond tip can be regarded as negligible during such a sequence photographs (fig. 5), they

can be regarded as an accurate composite picture of the chip forming process.

(e)

Figure 5. Chips from forsterite ceramics at

different grinding stages. U= 12.5 m/s, 1; =

0.042 m/s.

(e)

Figure 6. Chips from insulator porcelain at

different grinding stages. U= 12.5 m/s, v =

0.042 m/s.

39
(e) Figure 7. Chips from glass-ceramics at different

grinding stages. [7 = 12.5 m/s, 1; = 0.042 m/s.

(e)

Figure 8. Chips from alumina ceramics at different grinding stages. Z7=12.5 m/s, v = 0.042

m/s.

Figure 9. Chips from steatite ceramics at different grinding stages, f7 = 12,5 m/s, 11 = 0.042

m/s.

0 0.5nrB

Figure 10. Chips from synthetic sapphire.

Photographed at the position "c" in figure 4, C7=12.5 m/s, v = 0.042 m/s.

Figure 11. C/iip /rom hardened steel. Photographed at the position "c" in figure 4. [7=12.5 m/s, y = 0.042 m/s.

Figures 6, 7, 8, and 9 show typical chips

observed at different stages in the grinding

of insulator porcelain, glass-ceramics, alumina

and steatite, respectively. From the figures

illustrated above, it is clearly seen that the

chips from these materials under the grinding

conditions examined are commonly of a fragment type. This characteristic is also seen in

grinding synthetic sapphire (fig. 10).

In the case of grinding ductile materials,

the removed chips have been observed to curl

40
(e) Figure 12. Chips from forsterite ceramics at

different grinding stages. [7 = 37.5 m/s,

v = 0.125 m/s.

(e)

Figure 13. Chips from steatite ceramics at different grinding stages. [7 = 37.5 m/s, v = 0.125

m/s.

(e) Hardened steel Figure 14. Chips from various materials. Photographed at the position "c" in figure 4,

[7 = 37.5 m/s, v = 0.125 m/s,

Figure 15. Continuous chip formed in grinding

glass-ceramics under small depth of cut. Photographed at the position "c" in figure 4. U= 37.5 m/s, '!; = 0.125 m/s.

in front of the abrasive grain, which bear some

resemblances to the one formed in turing with

a bit.' The shape of chips varies, of course,

according to the shape of abrasive grain tip,

but the chips from ductile materials are generally of a "continuous" type.* An example

of the chip during grinding hardened steel with

the diamond model grain is shown in figure 11.

On the other hand, it is seen from figures 5

through 10 that a considerable part of chips are flying off in front of the grain. From the

blur of the photographs, their flying velocities are thought to be considerably higher than that

of the peripheral speed of the grain. Moreover,

part of chips are observed to splinter out of the

specimens surface in a short time after the

grain has passed over. This phenomenon is

pronounced, especially in the grinding of forIraanaka, O., Fujino, S., and Mineta, S., Bull. Electrotech. Lab.

32, 907 (1968).

41
(a) Porsterite I I I I I I 0 0.05nmi

(b) Steatite

Figure 16. Ground tracks on forsterite and steatite surfaces. Z7=12.5 m/s, v = 0M2 m/s.

sterite, insulator porcelain and glass-ceramics.

The amount of such chip splintering after

passage of the grain is fairly small in the case

of steatite.

The chips formed at higher grinding speed

are shown in figures 12 through 14. The photographs in figure 14 were taken at the position

of "c" in figure 4. General characteristics are

almost the same as those seen at lower grinding

speed. But the velocities of chips flying out in

front of the grain are observed to be much

higher than those at lower grinding speed.

In addition, it should be pointed out that,

even in grinding brittle materials, a sort of

continuous type chips are likely to be formed

at high grinding speed and under moderately

small depth of cut as seen in photographs (d)

and (e) in figure 13. Under the conditions ex- amined, "continuous" chips were apt to be

observed in grinding steatite. Another example

of this type chip is illustrated in figure 15,

where the depth of cut adopted was less than

2 ixm.

Fig. 16 shows the ground tracks on the forsterite and steatite surfaces. In grinding forsterite, which is the most brittle ceramic

u O.lnim Q o.05nm'

(a) Porsterite (b) Steatite

Figure 17. Chips sticking to pyramidal face of

grain. U=12.5 m/s, i; = 0.042 m/s.

examined, local fractures in the workpiece sur- face and deep fissures extending into the workpiece body are clearly seen. Existence of these

fissures suggests that a damaged layer may

remain under the ground surface. In grinding

steatite, though brittle features are observed,

the disorder of the cracks is much less than in

the above case. Fig. 17 illustrates the chips sticking to the

pyramidal face of the grain. Measurement

shows, a number of chips are larger in diameter or thickness than the grain depth of

cut. It should also be noticed that figure 17(b)

shows a "continuous" type chip (indicated by

arrow) similar to those in figure 13(d) or (e).

4. Summary

In order to make the cutting behavior of the

abrasive grain in ceramics grinding clear, the

chip formation processes have been observed

during grinding with a model grain by using

the synchronized illumination setup developed.

The main results obtained are as follows:

In grinding ceramics under moderately large

depth of cut, fragment type chips generally

appear, and part of them splinter out of the

specimen surface a short time after the grain

has passed over. Considering the fissures on the ground sur- face which extend into the workpiece body, a model of chip formation process in grinding

brittle ceramics may be proposed as shown

schematically in figure 18. The sketch illustrates

the cracks produced under ground surface, the

extent of which depends on the properties of

workpiece material and on stresses caused by

the abrasive grains. When the load at the contact area of the workpiece and the abrasive

grain is suddenly released at the instant that

the grain has just passed over, the portion

deformed elastically while being loaded will

tend to recover its deformation. This portion,

42
Chips splintering

from v/orkpiece surface after

Figure 18. Schematic model of chip formation process

in grinding brittle materials.

being prevented from recovering by tiny fragments of the vi^orkpiece material or by air compressed in the cracks, may make the shoulder

parts above the cracks splinter off. This is a

possible explanation for the splintering of chips

along the locus of the grain after its passing.

The type of chips also depends on the properties of workpiece materials and on the grinding conditions. In grinding steatite and glassceramics at higher grinding speed and under

moderately small depth of cut, chips of "continuous" type are occasionally formed. This

phenomenon is thought to be due to the microplasticity of the materials.

Concerning the relationship between the

grinding characteristics and the material properties of ceramics, a great deal of further work

is required.

NATIONAL BUREAU OF STANDARDS SPFCTAT ptir p/iq t c Proceedings of a Symposiun, Sponsored by tL American C^^^^^ Offi "^'"f Mac„,^:^c and Surface Finishing.

of Standards, held at NBS Nov. 2^, I9ll GaitheX^ MrSsLS MaVl972) ' ' '"'^ '"'^

43

Grinding—A Means of Shaping and Sizing Ceramics

J. A. Mueller

The Carborundum Company

Bonded Abrasives Division, Niagara Falls, New York 14302

The shaping and sizing of ceramic bodies can be done with abrasive grinding

wheels. From a grinding wheel point of view, ceramic bodies may be classified into

two groups. One group that is harder than 800 Knoop and the other group that is softer than 800 Knoop.

The materials that are harder than 800 Knoop generally require the use of diamond wheels. Diamond wheels have the capability to remove stock, to generate precise

geometry, and to produce finishes as fine as those that are measured in light bands.

The selection of the diamond grinding wheel is critical because improper selection can cause excessive grinding costs, poor quality, and low productivity. The selection of

the diamond grinding wheel includes not only the selection of the specification of the

wheel formulation but also the selection of the geometry of the wheel.

Case histories of finishing and shaping ceramic bodies illustrate the capability of

the grinding wheel in these areas. Suggestions for starting speeds, feeds, and wheels

together with a list of practical everyday "Rules of Thumb" for efficient grinding are documented.

Key words: Capability; ceramics; grinding wheels; selection; shaping; sizing.

1. Function of Grinding Wheels

The sizing and shaping of ceramics involves

practically every knovi^n kind of grinding operation. Offhand grinding, surface grinding, both

horizontal spindle and vertical spindle, disk

grinding, internal grinding, centertype and

centerless cylindrical grinding, thread grinding,

and cutting off are used in the manufacturing

of ceramic bodies. Grinding, therefore, is an

essential and important part of the manufacturing process.

It is the object of this article to define the

capability of abrasive grinding v^heels to shape

and size these ceramic bodies.

In a general classification, from a grinding

wheel point of view, these ceramics include tv^^o types of material. Prefired ceramics together

with ceramics softer than 800 Knoop hardness

represent one classification and fired ceramics

harder than 800 Knoop hardness represent the

second classification.

The materials softer than 800 Knoop are

often ground with silicon carbide wheels and a

satisfactory product can be produced. Traditional grinding practices are used and generally no difficulties are encountered.

The materials in the second classification

that are harder than 800 Knoop require the use

of diamond wheels and at times present some

difficulty in grinding, particularly if good

grinding practices are not used. It is to this

group of materials that the remainder of the

article will be directed.

The function of a grinding wheel includes

the following:

(a) Capability to remove stock.

(b) Capability to generate geometry.

(c) Capability to produce fine finishes.

In today's manufacturing many times one

grinding wheel is expected to perform all three

functions. In a great number of cases it can

and does. However, to efficiently perform all

three functions with one wheel necessarily demands that a compromise be effected between

one or more of the functions.

By documentation of case histories in specific areas of grinding, hopefully an awareness

of the capability of the grinding wheel will be

developed, and with this awareness, then, practical use can be made of the grinding wheel in

everyday operations to produce parts to size,

shape, and finish at minimum cost and satisfactory productivity.

2. Selection of Grinding Wheels

2.1. Grinding Wheel Formulation

The selection of the grinding wheel is important and critical if the lowest cost of grinding is to be effected without sacrificing quality

of grind or dimensional accuracy. Evidence of

this is shown dramatically in table 1. High

purity alumina pieces were surface ground on

a vertical spindle rotary table grinder. By substituting a coarser 80 grit wheel in place of a

finer 100 grit wheel, the cost of grinding was

reduced 28 percent. Then by substituting a 50

concentration diamond wheel in place of a 100

concentration diamond wheel, the cost was

reduced 61 percent.

Diamond concentration is defined as the number of carats of diamonds per cubic centimeter

of wheel. Lower concentration means less diamonds and therefore less cost of wheel.

Table 1 defines and demonstrates a rule of

45
Table 1. Effect of diamond grit size on the cost of grinding

Machine Vertical spindle rotary table grinder

Material

Type

Size

Grinding fluid

96% Alumina

0.75 cm wide X 15.2 cm long

Water soluble synthetic fluid 1-80 concentration.

Operating conditions

Total area ground

Table speed—r.p.m. Down feed rate

46.5 square cm

50

0.13 mm per minute

Results

Tolerance—flatness

Tolerance—squareness

+ 0.08 mm or -0.08 mm

+ 0.08 mm or -0.08 mm

Abrasive specification

Size

Type

Speed

Abrasive specification

Abrasive usage—mm

Number of dressings

45.7 cm X 1.9 cm X 40.6 cm - 1.9 cm face

2A2T

890 rpm.

MDIOO-PIOO-M MD80-P100-M

0.020/100 pes. 0.014/100 pes.

1—2 per shift 1 per shift

MD100-N50-M

0.008/100 pes. Seldom

diamond grinding that states, "It is advisable

to use the coarsest grit size diamond possible

that will satisfy the finish and Quality of grind

requirements." Coarser grit sizes will produce

a better finish when larger areas are ground

than when smaller areas are ground.

It also defines the rule that it is advisable

to use the lowest concentration diamond wheel

that will satisfy the quality requirements of

the finished product.

Until experience has been generated on a

specific grinding operation, it is advisable to use 100 concentration diamond wheels for centerless, surface, and cylindrical grinding.

Lower concentrations are adaptable for use in

offhand grinding operations and in cutting off.

Included in the selection of the grinding

wheel is not only the grit size, hardness, and

concentration but also the bond type. Metal

bond diamond wheels are generally the choice

for ceramic grinding because ceramics possess

high abrasion properties and consequently tend

to wear other bonds quickly. Resinoid bond can

be used particularly for producing fine finishes.

Resin bond wheels, however, wear away rapidly

and generally are expensive to use. Because of

their high wear characteristics, at times resin

bond diamond wheels have difficulty in generating dimensional accuracy.

Table 2 compares wheel life and grinding

rate as a function of the bond of the wheel.

2.2. Grinding Wheel Geometry

Recent performance data have documented

the impact that diamond wheel geometry can

have on performance and productivity.

Figure 1 illustrates a diamond wheel with

radial slots in it. This wheel compared to a

wheel without slots and having a solid face of

diamonds differs widely in performance. Table

3 describes this performance. The solid diamond

Table 2. Effect of diamond wheel bond type on

grinding performance

Wheel

Wheel Bond Wheel life Grinding rate

Metal Longer Slower

Vitrified

Resinoid Shorter Faster

face wheel was able to remove 1.83 mm from

an area of 2916 square cm of high purity alumina and then the wheel required dressing. By

slotting the wheel with a varying number of

slots, the amount of stock that could be re- moved without dressing ranged as high as 192

percent more. Translated into productivity, this

Figure 1. Slotted face resinoid bond diamond

grinding wheel.

46
Table 3. Grinding performance of solid face versus slotted face grinding wheel

Machine Vertical spindle rotary table grinder

Wheel size 45.7 cm X 1.9 cm X 40.6 cm-2.5 cm face.

Wheel speed 720 rpm—17.2 m per second

Grinding fluid Water soluble synthetic fluid 1-80 concentration.

1 able speed 33 rpm

Downfeed rate 0.10 mm per min.

Material ground / 96% Alumina—4 pieces 30.48 cm diameter X 2.22 cm thick.

\ Total area—2916 sq. cm. Wheel specification CMD180-L50-B8

Type of Wheel

Amount of infeed—^millimeters before

dressing was necessary

Percent increase in length of infeed

before dressing was required

Solid Rim 1.83

8 slots 2.64 44

16 slots 2.79 53

32 slots 3.35 83

64 slots 4.11 125

128 slots 5.33 192

NOTE: All slots 0.16 cm wide.

means that the solid wheel required three dressings to one dressing with the wheel having 128

slots. Dressing automatically means a loss in

productivity because of the time lost in dress

and in addition means some loss in total life

of the wheel because diamonds are dressed away

without producing any useful work. Many

times abrasive life is controlled by the amount

of dress. Frequent dress produces short life.

A softer wheel permitted to wear away somewhat faster often times has outproduced a wheel that wears away more slowly and re- quires frequent dressing.

After the wheel for grinding has been se- lected, it remains to define what the wheel

can do.

3. Grinding Wheel Capability

In stock removal capacity, diamond grinding

wheels have removed high purity alumina at

rates as high as 16.2 cubic cm per minute.

The ability to generate a surface finish of

high reflectivity is demonstrated in figures 2

and 3. Figure 2 is a picture of a 15.2 cm diameter flat made from high purity alumina that

was ground to a reflective finish such that the

scale can be seen reflected in the face of the

disc. Figure 3 illustrates the same capability

on a much smaller area of grinding.

Both pieces were ground on a vertical spindle

rotary table surface grinder using diamond

wheels and a water soluble grinding fluid. A

120 grit metal bond diamond wheel rough

Figure 2. 96 percent Alumina disc. Ground to a

reflective surface with a roughness of 0.08 /jl.

ground the material to a finish of 0.13-0.15 /x and the finish grind was produced by the 500

grit metal bond segmental wheel shown in

figure 4. A surface roughness of 0.05-0.08

was generated.

In grinding this material, caution needs to be

exercised so that the grinding wheel does not

produce an impact shock on the piece. Too

great an impact may crack it. In order to avoid

this, the part to be ground should be placed

as close to the center of the chuck as possible

so that the wheel remains on the work piece

all during the grinding cycle.

Grinding wheels, therefore, possess the capability of generating surface roughness on ce- ramic parts to as fine as 0.05 fi. On the other

47
Figure 3. 96 percent alumina rectangle

ground to a reflective surface with a

roughness of 0.08 jx.

hand, rougher surfaces may be generated with

no difficulty.

Table 4 broadly defines the range of surface

roughness that can be produced with various

grit sizes of diamond grinding wheels. The

quality of the material to be ground will change

the ability of the grinding wheel to produce a

finer finish. Generally finer grained bodies permit finer finishes than coarser grain ceramic

bodies.

Manipulation of the operating conditions of

the grinder will permit some range of finish

that can be generated by the same wheel.

Figure 4. 500 Grit metal bond segmental type

diamond wheel.

Table 5 demonstrates the change in surface

roughness that was produced by lengthening

the time of spark out. This example shows a refinement of finish that appears to be small,

yet this change was difficult to achieve because

the actual finish is so fine. On parts that have

a rougher finish, refinement of finish shows

more dramatically the eff'ect of increased spark

out time.

Figure 5 shows the body whose finish was

Table 4. Effect of diamond grit size on surface

roughness

Diamond wheel Finish

grit size /I

100 0.38—0.76

320 0.13—0.38

500 0.05—0.13

Table 5. Effect of spark out on surface roughness

Machine Vertical spindle rotary table grinder

Wheel speed 720 rpm—17.2 m per second

Wheel type Segmental

Material ground High purity alumina

Grinding fluid Water soluble synthetic fluid 1 —80 concentration

Table speed 33 rpm

Abrasive specification

Type Segmental—60 segments 0.95 cm X 2.54 cm

Grading D500-N100-M

Diameter 45.7 cm

Total infeed Rate of infeed Spark out Finish

mm

0.13

mm/min min

0.03 5 0.15—0.20

0.30 0.03 10 0.15—0.18

0.30 0.03 15 0.15 Newtonian rings visible

48
Figure 5. 96 Percent alumina ceramic valve

disc ground to a finish of 0.15 fi.

refined by spark out to such a degree that Newtonian bands appeared on the ground surface.

The 500 grit metal bond segmental wheel shown

in figure 4 was used for this grinding operation.

The third function of a grinding wheel is to

generate geometry and dimensional accuracy.

Dimensional accuracy capability has been demonstrated by the ability to produce a fine finish.

Finishes in this range necessarily mean accurate dimension.

Generation of geometry on the other hand

involves shaping of complex forms and shapes

that require ability to grind angles, radii, contours, and at the same time blend in contour

with angles.

Figure 6 illustrates such a part. This part

necessitates the forming of a radius, an angle, or taper, and a blend of the taper and the

radius.

Grinding wheels possess the capability of

performing work of this kind. This particular

form can be generated by using a straight lAl

grinding wheel or a formed wheel.

If a lAl wheel were used it would need to be

a small wheel mounted on a mandrel and inserted into a spindle. The entire spindle mount

would then need to be swung in a radius equivalent to the radius being formed. The taper

would be ground in a conventional manner.

This kind of part can also be ground by using

a formed wheel such as shown in figure 7. The

advantage of this method over the method using

an lAl wheel is that substantially higher productivity can be achieved. Formed diamond

wheels are becoming widely used and offer an

opportunity to grind complex shapes at fast

rates of productivity.

Threads can be ground into ceramic bodies

using conventional thread grinding practices

with a diamond grinding wheel.

Another example of the ability of the grinding wheel to fabricate a complex form is shown

in figure 8. This cone was ground with a res- inoid bond diamond wheel to a wall thickness

of 0.51 mm without cracking or breaking the

MAX. i.yj.7" MIN. 1.1|27"

BREAK CORNER

i|5 X .03

SECTION " A - A

"

Figure 6. Alumina nozzle insert.

49
Figure 7. Metal bond diamond formed

grinding wheel.

\\\ mm m ^ 1 2

Figure 8. High Purity alumina cone ground to a wall thickness of 0.5 mm.

cone. Grinding wheels possess the demonstrated

capability of being able to grind shapes and

forms into ceramic bodies.

4. Grinding Machine Availability

In addition to the grinding wheel that is

used, the grinding machine on which the wheel

is mounted can play a critical role in the manufacture of ceramic bodies. Because of their

inherent characteristics, the grinding of ceramics generally require rigid machines to prevent or minimize wheel vibration and its

consequent effect on finish, quality, and even

chipping or breakage. Fortunately, machine

design has kept pace with the advance in grinding wheel capability.

Figure 9 typifies a machine that can grind

a variety of shapes with the use of a grinding

wheel. The machine essentially is a center type

Figure 9. Vertical spindle template grinder

using a metal bond diamond grinding wheel.

cylindrical grinder. The workpiece, which in

this illustration is the cone, rotates on its center. The grinding wheel is driven and traverses

across the work. The direction of rotation of

the wheel and the work are the same so that

at the point of contact, the wheel and work

direction are opposed. Infeed motion is provided for sizing and provision can also be made

for automatic size control.

The path of the grinding wheel is controlled

by a template. In general terms, the grinding

wheel will be able to follow whatever path

the template provides.

A similar operation on a horizontal spindle

surface grinder is shown in figure 10.

Machines are available to permit the grinding wheel to grind a multiplicity of shapes in a variety of sizes.

The fabrication of ceramic bodies necessitates many times the use of an abrasive wheel

to cut off pieces.

Table 6 illustrates the difference in chipping

that was produced by using a segmental type

lAlRSS cut off wheel and a continuous rim

type lAlR cut off wheel. The segmental wheel

produced chipping, the continuous rim wheel

was able to cut without chipping. This is particularly significant because the continuous rim

wheel that was used was 46 grit which was

substantially coarser than the segmental wheel

and in addition, the continuous rim wheel was

50 concentration compared to 100 concentration for the segmental wheel.

50
Tablk 6. Effect of solid rim versus segmental rim cut off wheels on chipping

Machine Cut off

Wheel speed

Cutting fluid

Material cut

Procedure

Wheel size

3080 rpm—57.3 m per s Tap water

Dense porcelain—Size 5.08 cm X 6.35 cm

Plunge cut full depth.

35.6 cm X 0.32 cm X 2.54 cm

lAlR lAlRSS lAlRSS

Wheel grading MD46-N50-M MDIOO-NIOO-M MDIOO-LIOO-M

Time per cut (seconds) 61 135 237

Quality of cut No chipping Chipping Chipping

Figure 10. Horizontal spindle template grinder

rising a formed diamond grinding wheel.

To provide chip free cuts, it is advisable to use continuous rim cut off wheels in as coarse

a grit size as the quality w^ill permit.

The specification of the grinding v^^heel and

the conditions under v^^hich that wheel should be

used are most difficult to define in a general

way. However, to provide a starting point so that experience may be gained, suggestions are

offered in table 7 for grinding the hard dense

ceramic bodies.

5. Rules of Thumb

In addition to these suggested starting points,

and for the same purpose, the following "Rules

of Thumb" are listed.

(a) Use the coarsest diamond grit size

wheel that can satisfy the finish and

quality requirements.

(b) It is advisable to use 100 concentration diamond wheels on any specific

job until experience has been gained.

(c) Metal bond diamond wheels produce

long life and slow grinding rates.

(d) Resinoid bond diamond wheels produce

shorter life and faster grinding rates.

(e) The longer the time of spark out, the

finer the finish and the flatter the

surface.

(f ) For fine finishes, use fine grit wheels.

(g) For coarse finishes, use coarse grit

wheels.

(h) For chip free quality cuts use a lAlR

continuous rim cut off wheel.

(i) When grinding large areas, coarser

grit diamond wheels may be used to

improve performance.

(j) Ceramic grinding uses grinding practices similar to those used in metallic

grinding and consequently knowledge

of metallic grinding can be transferred

for use on ceramic grinding.

6. Summary

(a) Abrasive grinding wheels, including

diamond wheels, possess the capability to:

1. Remove stock rapidly.

2. Generate precision geometry.

3. Produce fine finishes.

(b) Machines are available to utilize the

capability of the grinding wheel.

(c) The limiting factor in sizing and shaping ceramics is often the composition

and geometry of the body itself.

(d) Grinding of ceramic bodies essentially

is the same as the grinding of metal.

Knowledge of metal grinding, therefore, can be transferred for use in

the grinding of ceramics.

51
Table 7. Starting grinding wheel recommendations for cylindrical and surface grinding

Grinding wheel specification Cylindrical grinding Vertical spindle surface grinding

Grain type Diamond Diamond

Roughing 100—150 100—150

Finishing 220—500 220—500

Lrracie "NT IN "VT JN Bond preferred Metal Metal

Bond to eliminate chipping Resinoid Resinoid

Speeds and feeds

Grinding wheel speed—m per sec. 27.9—33.0 22.9—27.9

Work speed— per sec. 0.51

Table speed—rpm 33

Traverse Speed—cm per min. 7.6—15.2

Crossfeed—mm 0.25—0.50

Infeed—mm 0.01—0.08

Infeed Rate—mm per min. 0.01—0.25

Grinding fluid Water soluble Water soluble

NOTE: These specifications are starting points. They may need refining depending upon the area of contact,

the type of material that is ground, and other specific characteristics of the local grinding operation.

Discussion

soft grade vitrified bonded wheels;

slow wheel speed—2000 to 3000

S.F.P.M.;

light infeeds— 0.0002 in to 0.0005

in;

fast table speeds—80 F.P.M. and

higher

sharply dressed grinding wheel with

a sharp diamond

dress grinding wheel before spark

out ; and

selected straight grinding oils for use as the grinding fluids.

(c)

(d)

Heuer: What is "spark out" ? (a)

Mueller: "Spark out" is simply a term for

permitting the wheel to grind without any

additional in-feed. It is a method of relieving

the pressure on the grinding wheel.

Perry: If the goal is to minimize structural

damage to the surface how would you modify

your rules of thumb for the selection of grinding parameters ?

( f

)

Mueller: To minimize surface damage, the

following parameters of grinding are sug- (g)

gested:

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-^, 1970, Gaithersburg, Md. (Issued May 1972).

52
Aspects of Machining Glass-Ceramic Materials

G. H. Allgeyer

Corporate Research Department, Owens-Illinois, Inc.

Toledo, Ohio 43651

and

L. V. ColwelP

Department of Mechanical Engineering, University of Michigan

Ann Arbor, Michigan 48104

The manufacture of optical and mechanical parts from glass-ceramic materials

usually requires the removal of unwanted stock resulting from casting, hot forming,

or the cutting of blocks from large ingots or blanks. Desired surface finishes at

specified dimensions are the usual requirements. Conventional metalworking cutting

tools did not produce satisfactory results. Scratching, splintering and brittle fracture were responsible for the material removal action in the sawing, drilling, turning and

milling operations attempted. Very rough unsatisfactory surfaces resulted.

The physical and mechanical properties of CER-VIT®" material classify it as a

brittle, crystalline material in the usual machining state. The best controllable means

of material removal was found to be abrasive machining. The high hardness (550

Knoop) and abrasive character of these materials conclusively recommends diamond

as the most satisfactory cutting medium. Machinability studies were conducted on CER-VIT® glass-ceramic material using diamond abrasive tooling while adjusting

the controllable machining variables. The results and relationships of these techniques

as applied to CER-VIT® glass-ceramic are discussed.

Significance of the experimental results are explored with particular emphasis

being placed upon the predictable performance of diamond tools and the feasibility

of programming of machining variables for utilization of numerically controlled

machining processes.

Key words: Abrasive machining; brittle chip formation; brittle fracture; chip accomodation; glass; glass-ceramics; machining and grinding technical ceramics; material

removal rate ; specific energy ; wheel structure.

1. Introduction

Mechanical parts can be machined as readily

from glass-ceramics as from metals providing

adequate consideration during the planning

stage is given to the unique properties of these

materials.

The most relevant of these properties is the

brittleness in chip formation which manifests

itself persistently over the entire range of useful cutting speeds and sizes of cut. Thus the

development of quantitative guide lines for

planning machining operations on technical

ceramics involves the measuring and analysis

of several operational factors v^ith emphasis

upon identifying the role of brittleness.

Early attempts to use the more conventional

turning and milling operations met with failure

not because of excessive tool wear but because

the brittle chip formation at the sizes of cut

peculiar to these operations with conventional

cutting tools resulted in substantially unacceptable conditions of surface finish and size con-

' Mr. Colwell is Professor of Mechanical Engineering at the University of Michigan, and Consultant to Owens-HIinois, Inc. '® A registered trademark of Owens-Illinois, Inc.

' Figures in brackets indicate the literature references at the end of this paper.

trol. The search for more productive techniques

narrowed rapidly toward abrasive machining

with various forms of diamond tools.

The authors do not claim originality for much

of the material discussed in this presentation,

but they have applied the results of basic re- search reported in the technical literature from

various parts of the world and have succeeded

in developing reliable guidelines for planning

machining operations on technical ceramics.

The nature of these guidelines and several

other important aspects of the problem are

discussed under the sections carrying the headings: Theoretical Considerations, Experimental

Results and Priorities for Extending Productivity.

2. Theoretical Considerations

The theoretical aspects of the machining of

glass ceramics are dominated by the chip formation almost to the exclusion of all others

except for fracture stress and surface energy.

Tanaka and Ikawa [1] investigated the grinding of relatively hard and strong sintered car- bides with both conically and spherically pointed

439-921 O - 72 - 5 53
diamonds and found that the normal stresses

were not only of the same order of magnitude

as the flow stress of the work material, but

also that they could be predicted mathematically

by the application of the principles of theory

of elasticity along with the geometry of the

individual diamonds. This has been explored

also by Colwell [2] in connection with the physical properties of the tool-work material combination and the changes in tool geometry arising out of tool wear.

These results, as well as related efforts by

many other investigators, have established that

it is necessary to build up a stress between the

tool and work material at least equal to the

flow stress in the case of brittle materials like

glass.

In the cutting of ductile metals the total

energy requirements arise from four sources. Shear-flow energy is dominant by requiring

70 to 90 percent of the total while most

of the remainder is needed to overcome friction. The energy required to create new surface

is almost negligible by comparison.

Glass by contrast requires no energy for

shear flow, which results in the requirements

being dominated by that required to create new

surface and to overcome friction. As a conse- quence chip formation and related particle size

become of prime importance.

2.1. Chip Formation

Investigators at the Electrotechnical Laboratory in Tokyo, Japan, carried out studies of

chip formation at actual grinding conditions

with the microflash technique. The results [3]

prompted them to characterize the mechanism

of chip formation in glass and ceramics as shown in figure 1. It was believed that the chips

or glass particles found in the grinding debris

come primarily from the shoulders designated

by the cross-hatched segments at the right in

the figure.

They noted also that:

a. Chips are most frequently larger than

the attempted thickness of cut, and

that

b. Chips were found to be formed and leave

the surface after the grain had passed

over the point of origin of the chip,

which lead them to suggest that most

such chips crack off upon elastic relaxation of the load created by the indenter.

The fact that the chips are brittle fragments,

and that they can be larger than the attempted

thickness of cut places emphasis upon grain

size of the abrasive and especially upon those

aspects of the structure of the grinding wheel

which determine its ability to accomodate the

chips; i.e., chip space.

2.2. Grinding Wheel Structure

Prokhorchik [4] of the U.S.S.R. investigated

the structure of diamond grinding with the

objective of developing techniques for predicting the capacity of any given wheel for accomodating the chips, since he had found that this

was the major limitation on the rate of removal.

The experience of the authors of this paper

has lead to substantially the same conclusion.

Prokhorchik concerned himself with the

model illustrated in figure 2 and tried to relate

(b)

Binder

Figure 1. Schematic diagram of splintering of glass

scratched by hard indenter (from 0. Imanaka,

et. al. [5]).

Workpiece

Figure 2. (From Prokhorchik [4])

Normal operating conditions will

accomodate all of the chips as at

(a). Too little projection of the

diamonds, too high a concentration

of diamonds, or too high a feed rate

will overcrowd the available chip

space as at (b)

54
the maximum capacity for chip accomodation

to grain size, diamond concentration and the

average projection from the sintered metal

matrix. This could be accomplished only empirically and was complicated by the fact that the

matrix material frequently projects into the

grinding surface, especially with fine grained

wheels as reported by Tanaka and Ikawa [1].

Despite these difficulties, Prokhorchik was

able to establish that the capacity for chip

accomodation was the most important deterrent

to greater removal rates, and further that such

capacity varied directly with diamond grain

size and exhibited an optimum in relation to

concentration. Experience gained by the authors

of this paper strongly support these correlations and point further toward the potentially

superior chip accomodation capacity of platingbonded, single-layer diamond tools compared to

the more conventional metal and resinoid

matrix-bonded, multi-layer tools.

3. Experimental Results

Against the background of theoretical considerations found in the technical literature,

the authors of this paper initiated a research

program to develop necessary guide lines for

applying basic principles in the machining of

CER-VIT® and companion materials in a

line of technical ceramics. The prime objective

of this effort was not only to meet the usual size

and surface finish specifications, but above all

to achieve a maximum of productivity while

fulfilling these other conditions.

The research program sought initially to find

the limits of removal rate and to identify or confirm its determinants. Much of this has been

carried out with conventional surface grinding

on the periphery of plain circular wheels, but

the established principles and guide lines have

also been confirmed in other operations which

must be performed in machining parts from

these materials.

It is always advisable to use the most sophisticated investigative techniques available when

searching for new information, but frequently

such simple parameters as average spindle

power and cutting forces may yield all one

needs to know to greatly improve a process.

This was the case in surface grinding CERVIT® glass-ceramic materials.

Figure 3 is a plot of net spindle power as a

function of rate of removal at constant wheel

surface speed but for ranges of both table speed

and depth of cut. The wheel was a single-layer

diamond wheel with a grain size of 60/80 grit

(177/250 X lO -'m). It will be noted that all

of the experimental points for removal rates

below about 2.5 cu in per minute (6.83 x 10"" cu m/s) are closely grouped about a single

straight line while those at higher removal rates

rr?/s (xlO^) 0 2 4 6 8 10 12 14

1.4

L2

10 • • *" /

Horsepower y^HPu=.22 -

•

. /(602 X 10^ W/m%)

Z .6

• / •

.4

.2

D 1 2 3 4 5

1000

800

400

200

Rate of Removal (inVmin)

Figure 3. Net spindle power versus rate of re- moval while surface grinding. Wheel: 60-80 grit

(177-250 X lO^m) single-layer diamond; Work

material CER-VIT® glass-ceramic.

are above the line because of greater friction

due to varying degrees of incipient overcrowding of the available chip space in the wheel.

The straight line in figure 3 is labeled HP,,

with a numerical value of 0.22. The designation

HP,, is commonly used in technical literature

on metal cutting and usually is referred to as "Unit Power" (specific energy), since it is

obtained by dividing net spindle horse power

by the rate of material removal. Consequently

the units are Horsepower per cu in per min

(the SI unit for specific energy is watts per

cubic meter per sec (W/mVs), and 0.22 hp/

inVmin = 6.02xlO « W/mVs). If the value of

HP,, is multiplied by 33,000 foot pounds per

minute (the equivalent of one horsepower), the

result is the number of foot pounds of energy

required to convert 1 cu in of the work material

into chips (The horsepower equivalent of 33,000

ft-lb/min is 746 joules per second (J/s), or 746

watts). Thus it would appear that the energy

requirements are not critically sensitive to the

variations of surface energy peculiar to the

variation of particle size as between light cuts

and heavy cuts. On the other hand, the specific

energy for fine grained wheels has been found

to be as much as two times greater.

The overcrowding tendency indicated by the

points above the line in figure 3 is delineated

more sharply when the unit power is plotted

55
c3

Q

1 .2

Log (X 10")

8 10 20 30

^HPy=Q22±l47o

=[6.02 X I0°]

20 40 60

Log Load Density

100 200 400 600

in^/min (Removal Rate)

inVmin (Supply of Tool Surface)

1000

x|0'

20

CO'Q

108

"1

6 o

_l

4

Figure 4. Specific energy requirements normalized against rate of

supply of tool surface area for a constant wheel speed and wide ranges

of depth of cut and table speed in surface grinding. Wheel and work

material same as in figure 3.

against the degree of utilization of available

chip space in the wheel. This has been done

on double logarithmic coordinates in figure 4.

The degree of utilization of chip space has

been designated as "Load Density" and has

been calculated by dividing the removal rate

by the area of wheel surface which has passed

over the surface being ground in the same

period of time. Thus the value of 1000 on the

abscissa represents the loading of one thousandth of a cu in of work material over each

sq in of wheel surface (25.4 X lO-^mVm-).

Similarly the figure 100 represents only one

tenth as much loading.

It is to be noted especially that the wheel

becomes seriously overloaded or overcrowded

at load densities above 400 microinches (1.016

xlO 'm). Dulling of the wheel will shift the

entire band of points upward but will have

little effect on the 400 microinch breaking point.

Similarly, wheel speed should have little effect,

but variations in the effectiveness of coolant

application can shift the breaking point substantially.

4. Priorities for Extending Productivity

As mentioned earlier, Prokhorchik established that chip accomodation was a most important limiting factor in achieving greater

material removal rates. The authors observed

that the size and shape of chips produced by

the finer grain sizes of diamond were much

more persistent in packing tightly in the available chip space than those produced by the coarser grain sizes. As the "Load Density" increases, friction increases, and the unit power

curve rises sharply as shown in figure 4. In

figure 5 the data in figure 4 has been plotted in

n?/rr?{^\d)

10 15 20

orX

3

.500

400

300

,200

.100

— Actual Data — Hypothetical

/ 200 Grit

\ /325 Grit (74«l(^«m) 100 Grit y(44 « IO«m) / (149 ^lo^^,

' /60/80 Grit

'(177/250 « 10 m)

Projected Breaking

Point

100

10 'b

I

300 500 700 900

Load Density -(inVin^« 10*)

Figure 5. Specific energy requirements versus Load

Density projected for a range of diamond grain sizes.

linear coordinates with an expanded "Unit

Power" scale. The sharp rise in unit power at

the origin indicates a high amount of frictional

energy due to rubbing caused by insufficient

depth of cut. The sharp rise to the right indicates the approach of full utilization of available chip space. The solid line was derived from

actual test data; the dotted lines are hypothetical projections for finer diamond grains.

As the diamond grain sizes are reduced, an

increase in the unit power may be anticipated

while the critical load density may be expected

to decline. The slope of the dashed line connecting the others may be relevant to the different

characteristics of a particular bond, or the diamond concentration. Large diamond grain sizes,

"open" bond, and effective coolant application

point the way to higher material removal rates

56
and greater productivity. Knowing the breaking points of the curves in figure 5 for the diamond tools to be used can point to an optimum

choice of variable factors for a particular

machining application.

5. Conclusions

The knowledge of predictable tool performance is essential to success in economical metalworking operations. The data resulting from

the authors' experimental effort in this instance

indicates convincingly that this is equally true

in the machining of glass-ceramics. Many machine operations involve the removal of a non- uniform depth of material from the workpiece.

This presents a problem in selecting a suitable

depth of cut, feed rate and tool surface cutting

speed for any individual cut.

All of these elements combine to control the

"Load Density" mentioned earlier. Consequently predictable performance of a diamond tool

suggests the feasibility of programmed feeds,

speeds and depth of cut essential for the application of Numerical and Adaptive Control

techniques to the machining of glass-ceramics.

6. References

[1] Tanaka, Yoshinobu, and Ikawa, Naoya, Cutting

mechnism of the bronze-bonded diamond Wheel,

Annals of the C.I.R.P. XIII pp. 443-440. Ms.

78/3. (1966).

[2] Colwell, L. v., Methods for sensing the rate of

tool wear, Preprint of a presentation to the

C.I.R.P. in Torino, Italy (1970).

[3] Imanaka, 0., Fujino, S., and Mineta, S., Observation of chip removing process in grinding by a microflash technique. Preprint of a presentation by A. Kobayashi to the C.I.R.P., Nottingham (1968).

[4] Prokhorchik, S. M., Productivity of a process for

grinding glass using an annular diamond tool,

Soviet Journal of Optical Technology, 34,

Number 3 (1967).

Discussion

Westwood: What cutting environments did

you use?

Allgeyer: Our environment was primarily

one of a circular wheel on a small surface grinder using a commercial coolant material.

Rice: I think Dr. Westwood was asking what

sort of coolant material ; was it a water soluble

one?

Allgeyer: It was a commercially available

water soluble material.

Martis: What type of single point cutter was

used in obtaining the unsatisfactory results on

the glass ceramic material?

Allgeyer: It was a commercial grade of one

of the tungsten carbides.

Martis: Was a coolant used?

Allgeyer: Yes and no. Both cases existed

here and I can't tell you now whether this was

one with or without coolant. We used a spray

mist form of coolant when we did use one. Martis: Is this a big application of single

point turning versus centerless grinding in this

type of machining?

Allgeyer: Actually, no. It was more of a

curiosity on my part. I had to explore it to see what the aspects of it really were. This type

of operation has been done on a production

basis in other applications to produce contoured

surfaces. Aspheric lens generation, for instance, at one time had been single-point turned.

national bureau of standards SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

57

The Principles of Grinding

Robert S. Hahn and Richard P. Lindsay ^

Cincinnati Milacron Corporation, Heald Division,

Worcester, Massachusetts 01606

Three distinct actions are shown to take place in the grinding process, namely

rubbing, ploughing, and cutting. A clue to the relative importance of each of these can be obtained by plotting the volumetric rate of stock removal against the wheelwork interface force for unit width (force intensity). Such plots show for some

materials a threshold force intensity below which no stock removal occurs. Equations

relating the stock removal rate and the interface force intensity are developed and

a "material removal parameter" is defined. Similar equations for the wheel wear

and a "wheel wear parameter" are also developed. It is found for many materials

that the stock removal rate is a linear function of the interface force intensity.

However the abrasive removal rate or dimensional wheel wear is found to vary

approximately as the square of the interface force intensity. Increasing wheel surface

speed is generally found to raise the material removal parameter.

The essential parameters in dressing with single pointed diamonds are introduced

and their effect on controlling the sharpness of a grinding wheel is discussed. Semi

empirical equations for predicting wheelwear are given in terms of force intensity,

wheel hardness and wheel-workpiece conformity. The factors influencing surface finish are described including the dressing lead, the depth of cut of the dresser diamond

and the interface force intensity. Thermal damage and surface integrity are dis- cussed. The existance of a threshold force intensity (or feed rate) below which no

thermal cracking takes place is shown. Raising the workspeed tends to raise this

threshold. Wheel sharpness is also found to be an important variable and means

for quantitatively measuring wheel sharpness are given.

Key words: Abrasive wear parameter; grinding; material removal parameter; principles of grinding; surface finish in grinding; surface integrity in ground surfaces;

wheel dressing.

1. Introduction

As manufactured components are used more

and more in critical areas where reliability is

important, a knowledge of the technological

factors in the manufacturing process and their

effect on the workpiece becomes essential. To

achieve greater productivity and lower costs

the proper setting of machining and grinding

parameters is also very important. In order to

accomplish the above it is necessary to develop

quantitative relationships between the important variables in the grinding process. Some of

these relationships are developed below.

Although the majority of the work to be

reported here deals with the grinding of metals,

the principles and the approach can be equally

well applied to the grinding of ceramic materials, sintered carbides etc. In the grinding

of metals Hahn [1]^ has shown that three distinct processes take place between the abrasive

grain and the work piece namely (a) rubbing

where the grain rubs on the work causing elastic and/or plastic deformation in the work

material with essentially no material removal,

(b) ploughing—where the grain causes plastic

flow of the work material in the direction of

sliding as well as transverse to the direction

' Manager of Research and Senior Research Engineer,

respectively.

' Figrares in brackets indicate the literature references at the end of this paper.

of sliding, extruded material being thrown up

and broken off along the sides of the groove

resulting in low rates of stock removal, (c)

cutting—where a fracture takes place in the

plastically stressed zone just ahead of the rubbing grain causing the formation of a chip

and resulting in fairly rapid stock removal

rates. These three processes can be observed

as shown in figure 21 where the "wheel depth

of cut" i.e., the radial advance of the wheel

per revolution of the workpiece in cylindrical

plunge grinding, is plotted against normal force

intensity (normal force per unit width of wheelwork contact). Rubbing takes place in figure

21 up to about 7 N/cm, ploughing up to about

26 N/cm and cutting thereafter. In this case the transition from rubbing to ploughing and

from ploughing to cutting is fairly distinct.

Recent grinding studies by Busch [2] on

glass, cemented carbides and several steels have

also shown evidence of rubbing, ploughing and

cutting. Under certain conditions he found

continuous chips in grinding glass and other

evidence of plastic deformation. The grinding

of sintered aluminum oxide and a group of

cemented carbides with diamomnd wheels was

studied recently by Buttner [3]. He observed

a migration of the diamond particles in the

resin bond as well as a charring of the bond

due to heat. Pahlitzsch [4] investigated the

grinding of sintered aluminum oxide and gives

59
resultant forces and wheel wear. Ida, Arai, and

Inomori [5] carried out surface grinding tests

on nine ceramic materials, observing grinding

forces, surface finish, and microstructure.

2. Material Removal Equations

In studying the grinding process it has been

the custom to report results in terms of down

feed (on surface grinders) or infeed rate (on

cylindrical grinders). Under these conditions

the grinding performance is dependent on the

rigidity of the machine tool. For example a down feed of 0.05 mm on a very stiff machine

will induce a much larger wheel-work contact

force than the same downfeed on a more compliant machine and accordingly the grinding

performance (surface finish, wheel wear, sur- face integrity) will differ even though the

same wheel and work material are used. Consequently, in order to become independent of

the machine tool rigidity it is convenient, in

plunge grinding operations to relate the grinding performance to the wheel-work interface

contact force instead of downfeed or feed rate.

Consider the grinding configuration shown

in figure 1, the cross slide is advancing at a feed

rate, Vf. Simultaneously the work radius is re- ceeding at a rate v,,-, while the wheel is wearing

at a rate, v.,. Obviously, in the steady state:

Vf = V,„ + (1)

During the steady state process, a force is

induced between the wheel and the workpiece.

The normal force per unit width, F„', or force

intensity is a convenient parameter to use in

describing grinding behavior. The metal re- moval rate per unit width can be written.

(2)

Likewise, the wheel removal rate or wheel

wear per unit width is:

(3)

Plotting data for external, internal and roFlGURE 1. O.D. center type grinder showing

system springs and nomenclature.

tary surface grinding [6, 7, 8] as shov^m in

figure 2, one notes that a linear relationship

exists between the material removal rate and

the force intensity, thus: — Aw Fn (4)

where A,„ is a constant of proportionality [9],

having dimensions of "volume of material re- moved per unit time per unit of wheelwork interface force and called the metal removal

parameter."

Using eq (2), eq (4) can be written in the

following two ways

Vn, = (5a)

(5b)

The first equation applies to controlled force

grinding where the force intensity is prescribed. The second equation applies to feed

rate grinding. In practice the wheel wear velocity v., in eq (1) is usually small compared to

v,n so that Vf can be substituted for t^,,, in eq

(5b) giving

Fn' =

TD,„Vf

A,,

(5c)

Equation (5c) gives the induced force due

to a feed rate Vf. It will be seen that the induced force is directly proportional to workpiece diameter and inversely proportional to

the Metal Removal Parameter.

Figure 2 is representative of "easy-to-grind"

materials such as normal tool steels, AISI

52100, AISI 4150 etc. In these materials a chip

is readily formed and material removal is predominately accomplished by the cutting process

with little or no rubbing or ploughing. The

Metal Removal Rate Z',„ plots as a straight line

going through the origin. On the other hand, for

"difiicult-to-grind" materials where excessive

rubbing and ploughing take place the plot of

Z',r vs Fn often exhibits a non linear behavior

and may have an intercept on the force intensity axis as illustrated in figure 3.

Since the slope of the Z',„ versus force intensity curve is A,,,, Metal Removal Parameters in

the range of 0.1 mm"* s-^N"^ are obtained for

these difficult-to-grind alloys, whereas from

figure 2, A,„ was 0.5 to 0.7 mm^ S"^ N-'^, without

any threshold force intensity.

Analogously, one may write the wheel re- moval rate or volumetric wear rate per unit

width as

Z', = nD.Vs (6)

From figure 4 it may be seen that a similar

relationship exists for wheelwear:

Z's = AgFn (7)

60
>•o

/

PAHllTiSCH EXTERNAL

80J7, Vs = 11820 sIm (59 M/Sl

12400 sIm (62M/S)

GRINDING

= 12500 sIm (62 M/SI

SURFACE GRINDING

i 60M8, q = 25, "Vs = 11500 sin * 70Ke, q = 25, (58 4 M/S)

500 (LB/INCH)

30 40 50 60 70 80 90 (N/mm)

FORCE INTENSITY,

ORENE 41 INCONEL Z

»M-2

500 (LB/INCH)

60 70 80 90 (N/mm)

FORCE INTENSITY, fJ

Figure 2. Metal removal rate versus force intensity

(after Pahlitzsch [6] J

External

Wheel:

Wheel Speed:

Q:

Wheel diameter:

Work diameter:

Work:

EK80J7VX and EK60L7VX

59 m/s and 62 m/s

56 and 113 and 59 and 118

28.3 in (720 mm)

4.0 in (102 mm)

RIV RINGS

Internal grinding

Wheel

Wheel Speed:

Wheel diameter

Work diameter:

Dressing lead:

Coolant:

Work material:

(Roller bearing cup 3920)

Rotary surface grinding

XA80M6VFMD2

63 m/s

15.6

3.25 in (82.5 mm)

3.75 in (95.2 mm)

0.004 in/rev (0.1 mm/rev)

Water-based soluble

AISI 4610, Rc = 62

Wheel

Wheelspeed

Wheel diameter:

Workpiece

Dressing lead:

Coolant

Work material:

32A70K8VBE

58.4 m/s

23

3.25 in (82.5 mm)

3 in O.D. X 2 in I.D. (76 mm O.D. X 50.8 mm I.D.)

0.004 in/rev (0.1 mm/rev)

Florex at 110 N/m'

AISI 52100 at Rock C 60

However if the wheelwear rates for the difficult-to-grind materials are plotted, one notes

that at high force intensities the Wheel Removal Rate Z's is no longer linear with F„'.

Figure 5 illustrates this behavior.

2.1. Factors Affecting A^

Since A,o is the "gain" of the material re- moval process, knowledge of the factors which

Figure 3. Metal removal rate versus force Intensity.

Wheel: 32A70M6VBE

Wheel speed: 33 m/s

q: 6.8 Wheel diameter: 1.73 in (44mm)

Work diameter: 1.87 in (47.5 mm)

Work material : M-2; T-15; Rene 41, Inconel X.

Coolant: Water based soluble

Internal grinding

z

.0005 ! = 0 004 INCH/SEV (0.1 MM/REV)

A J

= 0.250 ^= 0.125

LB/INCH)

10 15 20 25(N/MMI

FORCE INTENSITY Fn

Figure 4. Wheel removal rate versus force intensity.

Test conditions same as internal grinding for

figure 2. Additionally:

Dressing Ratio -y= 0.25, 0.125, 0.05

61
H4

500ILB/INCHI

90 (N/MM)

Figure 5. Wheel removal rate versus force intensity.

Test conditions same as figure 3

increase a„, would be important in reducing

grinding time and increasing production.

a. Wheel Speed

The effect of wheelspeed upon grinding is

fairly well known. Giihring [10] externallyground relatively soft steel and obtained decreasing force intensities at a given feed rate.

Replotting his data against F„' as in figure 6

and noting A„, to be the slope of the Z\„ vs. force intensity curve, one obtains A„, values

from 0.27 to 2.75 mm^* sec N as the wheelspeed is increased from 20 m/s to 90 m/s.

Internal grinding of roller bearing cups, as

illustrated in figure 7, also shows increasing

values for larger wheelspeeds.

If one plots A,„ versus wheelspeed, as in figure

8, one sees essentially a linear relationship.

b. Area in Contact and Dressing

It has been previously shown [11] that A„,

depends upon the real stress existing in the contact interface between the wheel and work.

Measuring the area of the wear flats existing

on the wheel by planimetering a photo micrograph and considering the length of the wheelwork contact, a plot of A„, versus normal contact stress (normal force divided by real area

of contact) can be obtained as shown in figure

9. Thus to maximize A,„, the real area of contact

O V s = 3920 ifm (20 M/S) Vs = 7840 (40) O Vi = 11,800 (60) V Vs = 15.800 (80) Vi = 18,000 (90)

Figure 6. Metal removal rate versus force intensity.

Wheel

Wheel speed:

Wheel diameter:

Work diameter:

Dressing lead

Coolant

Work material

EK80L7VX

3920, 7840, 11,800, 15,000,

18,000 sfm (20, 40, 60, 80,

90 m/s)

60

19.6 in (500 mm)

1.37 in (35 mm)

0.004 in/rev (0.1 mm/rev)

Cutting oil CK45N (AISI 1045)

at Rc = 23 (Vickers 250)

External grinding (after Konrad Guhring [10]

should be made as small as possible. (In figure

9, larger real areas resulted from flats wearing

on the grains through extended usage of the

wheel without dressing. As the area increases,

the real stress falls (at constant normal force)

and A,o is reduced).

Since dressing with a diamond is a way to

sharpen the wheel, one can define the dressing

parameters as in figure 10. The depth of diamond penetration while dressing is denoted as c/2 since compensation in the trade is generally

referred to diameter. The dress lead is the distance that the diamond moves along the wheel

axis per wheel revolution. The fact that dressing can actually machine a "thread" into the

wheel is shown in figure 11 which is a trace

(Bendix Profilometer) of a ground workpiece.

The periodic waves noted most easily in the

middle trace correspond to the dressing lead

used in dressing the wheel prior to a 3 second

plunge grind.

If the diamond point is assumed sharp, the

62
JllB/INCHl

Figure 7. Metal removal rate versus force intensity.

Wheel: 97A80L6VFMD2

Wheel speed: 6000, 8000, 10,000, 14,000 sfm

(30, 40, 50, 70 m/s)

400 sfm (2 m/s)

2.37 in (60.3 mm)

2.74 in (69.6 mm)

0.0035 in/rev (0.09 mm/rev)

Work speed:

Wheel diameter

Work diameter:

Dressing lead:

Climb work

rotation

Coolant

Work material

(Roller bearing cup 25520)

Internal grinding

Water soluble oil

AISI 4620 at Rc = 60

meaning of c/l, the dressing ratio, can be noted

from figure 12. As c/l approaches 1.0 the

"thread" on the grains becomes sharp, whereas

if c/l approaches zero, the dressed grain appears dull (figure 12 is schematic, but descriptive of what occurs in dressing). Thus increasing the dressing ratio c/l, at constant lead

produces a sharper wheel. Internal grinding of

roller bearing cups produced the data shown in

figure 13. As c/l is increased from 0.05 to 1.0,

A,c increases from 0.24 to 0.68 mm^ S"^ N-% a

factor of 2.85.

Also, at constant c/l = 1.0, with a lead of 0.1

mm/rev, = 0.68 mm^ s-^ whereas with

a lead of 0.025 mm/rev, A,„ = 0.40 mm S"^ N-\

as seen in figure 14.

Thus larger leads and larger c/l values in- crease A,o, as seen in figure 15. Both these ac- tions reduce the real area on the grinding wheel

thereby increasing A,,, by increasing the real

stress at the interface.

2.2. Wheelwear

As has been previously suggested, the wheelwear rate, Z'^ versus F„' curve is nonlinear.

Figure 16 shows stock removal and wheelwear

measurements for rotary surface grinding of

52100 steel (Rockwell C 60-62) and points out

very clearly the increasing wheelwear rate as

force intensity is increased.

To better define the wearing action, consider

figure 17 which shows schematically the elastic

action of the wheel in contacting a workpiece.

In production grinding at workspeeds above 100

fpm (0.5 m/s) the wheel depth of cut (advance

per revolution of work) is generally small compared to the elastic flattening of the grinding

wheel (Youngs modulus for a Vitreous bonded

wheel is approximately 1/4 that of steel) and

may be neglected in calculating the length of

the interference zone between wheel and work.

For low workspeeds as in reciprocating surface

grinding neglecting the wheel depth of cut may

not be permissible.

Neglecting the influence of the wheel depth

of cut on the length of contact, the length of

elastic contact k has been found by Lindsay

[12] to be:

Ir

r inches ]

[ meters J r DsD,o (

10-^

1.335 xlO-«

(1.33 + 2.2S - 8)

)]" -

where (lb and in) or {N and meters) are to be

used. The values of H and S are given below.

In plunge grinding the average normal wheelwork contact pressure P is defined as P = F ' (9)

or

lb in- 1000

0.75 xlO« IV D,,D, )

d?'^ (l.33iy + 2.2S-8)''' {Fn'Y" (10)

where: /^'""^^ V^" is used for internal

V

D.cDs I grinding

(D +D \^^^ ^ J-) J

is used for external grinding

is used for surface grinding

(1.33H -I- 2.2 S -8) is the percent of bond in

the wheel.

It can be found by:

Hardness of wheel: H, I, J, K, L, M etc. 1 "H" factor is : 0, 1, 2, 3, 4, 5 etc. J and

63
500

-6 «I0

/GUHRING EXTERNAL

PAHIITZSCH EXTERNAL

INTERNAL

10.000 (sfm)

I U _l L_ (M/SECI WHEELSPEED V«

Figure 8. The metal removal parameter versus wheelspeed.

Test conditions: , r, c c Guhring [10] external: Same as figure 6

Pahlitzsch external [6] Same as figure 2

Internal: Same as figure 7

<

-6

10

O 60 LB/INCH 00.4 N/MM)

(S 31 LB/INCH (5.36 N/MM)

• 120 LB/INCH (20.8 N/MM)

50,000 (LB/INCH*) '00,000 jlB/IN' 140,000

_l_ (N/MM'') REAL CONTACT STRESS

Figure 9. The metal removal parameter versus the

real contact stress. A60L5V

7200 sfm (36 m/s)

6.1 1.87 in (47.5 mm)

2.50 in (63.5 mm)

AIS152100, Rock C60

Texaco soluble D

Wheel

Wheel speed:

Q:

Wheel diameter:

Work diameter

Work material

Coolant

Climb work rotation

Dress lead: 0.003 in/rev (0.075 mm/rev)

L IS THE DRESSING LEAD.

IT IS RELATED TO THE WHEEL

ROTARY SPEED, AND THE

DRESSING VELOCITY.Vd.AS:

. _ vd_ inch /second

,L. N revs./second

N_iWORK ROTARY SPEED, RPS

r

DEPTH OF

DRESS: C/2

Vd.VELOCITY OF

DIAMOND ALONG

WHEEL AXIS,

INCH/SECOND

Figure 10. Dressing nomenclature.

S is the structure number of the wheel:

S = 4,6,S etc. Thus an 80M6 wheel has H= 5, S=6, and

so (1.33jtf + 2.25-8) = 1.33x5 + 2.2x6-

8) = 11.85

One may plot the pressure P existing for any

F„' and geometry.

Z'

If the factor —- is plotted versus P as in

figure 18, an empirical equation for wheelwear

may be written as: Z\{ inch ^ meter

Since P is a function of force intensity and

since A. is defined as the derivative of Z'^ with

3.0x10-1^ 1 ^0.22x10-^1

64
1—1 m [DR1 -EA

1 m n D s

1«)

lO

^- hh K> —j^'V m /If A

1

r

1

Figure 11. Proficorder traces of some workpieces.

Test conditions: Same as internal grinding from figure 2.

DRESS LEAD

i

^DEPTH OF DRESS = 0.05 L

d, AVE. GRAIN DIAMETER

i

DEPTH OF

DRESS =.25 L

DEPTH OF DRESS = .Sl

Figure 12. Dressing ratio and its significance.

respect to F„', one can insert eq (10) into eq

(11) and take the derivative to obtain

A., <^

in'

sec. lb

sec. N

6.0x10-8

0.186x10-=*

d^ (1.33H + 2.25-8) (12)

From eq (12) it v^^ill be seen that the wheel

wear parameter As is linear with force intensity. Using eq (12) the wheel wear eq (7) can

now be written as:

4 = 1

Figure 13. Metal removal rate ver- sus force intensity.

Test conditions: Same as internal

grinding from

figure 2.

Z's< sec

• = <

sec.

r 6.0x10-8

0.186 XlO-\

(1.33/^+2.25-8) Vs {Fn'Y (13)

It will be seen from eq (13) that the wheel wear

rate is a quadratic function of force intensity as was observed experimentally in figure 16.

65
^ i= 0.004 (0.1 MM/REV)

300 (LB/INCH)

Figure 14. Metal removal rate versus force intensity.

Test conditions: Same as internal grinding from

figure 2.

_ s

... S

G

RATIO

300 (IB/INCH) 500

300

1.0

0.5

OL

DRESSING RATIO C/L

Figure 15. The metal removal parameter versus the

dressing ratio.

Test conditions: Same as internal grinding from

figure 2.

Figure 16. Metal removal rate and wheelwear rate

and G ratio versus force intensity.

Test conditions: Same as rotary surface grinding from figure 2,

UNDEFORMED WHEEL SHAPE

Figure 17. Grinding wheel deformation

in an internal workpiece showing no- menclature (schematic).

3. Surface Finish

External grinding date published by Guhring

[10] and plotted in figure 19 indicates that the

surface finish is improved by increasing wheelspeed at a given stock removal rate. Hov^ever,

one may obtain the force intensity induced

during these grinds and cross-plot the surface

finish obtained versus the force intensity as

shov^^n in figure 20. This causes the external

grinding data of Guhring [10] to fall essentially on one curve.

Also shov^^n in figure 20 is internal grinding

data and the same relationship to force intensity is seen to exist: the surface finish is proportional to (F„')*^^ Moreover it has been experimentally observed that:

/ C

\3/10 / \ 1/3 surface finish - i ( ^

) fF„') (14)

Thus, low dress leads and dressing ratios as

well as low force intensities produce low surface

finishes (fig 21); wheelspeed only affects the

66
Hil!>

o 7

-

-

1

°i O © vir

o O InconelX

Rani 41 V 1-15 o 52100 a Rent 80 A Ren« 80 ^—

A Rene 80

-

o

o /

r Pa

la 7

L M ')p 3.0

1 1 1 500 1000 (LB/INCH2) 5000 10,000

' ' ' ' I I I I ' ' ' '

' 50,000

10 (N/MM2) 50 100 GROSS PRESSURE

Figure 18. Wheel removal factor versus gross contact

pressure.

Data from: A. Rotary surface grinding (test

conditions: See figure 2) B. Internal grinding

(Test conditions: See figure 3)

(Test conditions: See figure 2)

^7 ra

50

IVVs = 3,920 sfm (20 M/S)

tXTERHAl < V s = 11,800 sfm (60 M/SI

( A V s = I8«00 sfm (90 M/S)

INTERNAL .Vs = 12,500 sfm, c/j 1.0

100 ILB/INCH) 200

I 20 (N/MM)

Figure 20. Surface finish versus force intensity.

Test conditions: External grinding: Same

figure 6. Internal grinding: Same

figure 2.

Figure 19. Surface finish versus metal removal rate. Test conditions : Same as figure 6.

67
FORCE INTENSITY

Figure 21. Wheel depth of cut (radial advance per

work rev.) versus force intensity.

Wheel

Wheelspeed

Work speed

Wheel diameter:

Work diameter:

Dress lead:

Conventional work

Coolant:

Work material

38A801L8VBE

566 RPS

23 RPS

0.87 in (22 mm)

1.50 in (38 mm)

0.0001 in/rev (2.5 /j-m/rev)

rotation sense Flow Rex 100

AISI 4150 53-55 Rc

finish in so far as it reduces the induced wheelwork force during feed rate grinding.

4. Surface Integrity

Various authors [3, 4, 5, 6, 7, 8, 9] have discussed the stresses present in the surface layer

of ground workpieces. Generally to prevent

damage in the workpiece the admonitions are

to use a "gentle grind" consisting of low wheel

speeds, a lubricating rather than a cooling fluid

(i.e., oil rather than water), sharp wheels and

low down-feed rates. Unfortunately most of the

reported data was obtained on reciprocating

surface grinders where work speed is generally

less than 50 sfm (0.25 m/s). Obviously, the

shorter the wheel-work contact time the smaller

the heat input to the work. Work speed is a very important parameter in regard to surface

integrity and has been nearly completely ignored by most authors. In fact, in most of the

test work, the work speed has been at the worst

possible value: extremely low.

In order to investigate the surface integrity

problem, it is important to recognize that work

speed, interface force intensity and wheel sharpness are important variables relating to surface

damage. The most elusive variable among these

is the wheel sharpness. The wheel sharpness can

be defined in terms of the metal removal parameter.

The slope of the Z',„ vs. force intensity curve

has already been seen to be the metal removal

parameter, A„. It has also been shown [11] that

if a grinding wheel is dressed and then used

continuously under precision grinding conditions, wear flats will develop on the abrasive

grains. If gross wheelwear does not occur, (i.e.,

the wheelwear is attritions) these flat areas will

500

CUMULATIVE GRINDING TIME (SECONDS)

Figure 22. Metal removal parameter versus cumulative

grinding time.

60 lb/in (10500. N/m) Wheel A60L5V

A

A

30 lb/in ( 5250

120 lb/in (21000

60 Ib/in (10500.

20 lb/in ( 3500

100 lb/in (17500

Wheelspeed:

Work speed

Wheel diameter:

Work diameter:

Dress lead:

Climb work rotation

Coolant:

Work material

Internal grinding:

N/m) Wheel A60L5V

N/m) Wheel A60L5V

N/m) Wheel 2A90P6

N/m) Wheel 2A90P6

N/m) Wheel 2A90P6

7600 FPM (38 m/s)

1100 FPM (5.5 m/s)

1.87 in (47.5 mm)

2.37 in (60.2 mm)

0.003 in/rev (0.076

mm/rev)

Flow rex 100

AISI 52100 Rc 60

68
continue to grow to create some stable

characteristic real area of contact. If in addition, these tests be performed under conditions

where the interface normal grinding force between the wheel and work is held constant, it will be found that the developed grinding rate

v,o (rate of change of work radius) will diminish with time. Thus the wheel is losing it's ability to remove material through the dulling

action. Now since the growth of the real area

causes v„-, at some force intensity, to diminish,

then the metal removal parameter must be decreasing with time.

Figure 22 shows the relationship of A^, with

time under controlled-force, precision internal

grinding conditions. A,(, then is an indicator of

the sharpness of the wheel as measured by it's

ability to remove stock: When a,„ is large the

wheel is acting sharp, when dulling occurs A,;, diminishes.

Force Intensity: 60 lb /inch Force Intensity: 80 lb/inch

10500 N/m 14000 N/m

Since A„, is then a measure of wheel-work

grinding capability, we have a quantitative

means of describing wheel sharpness. It has

been shown that variations in the dressing ratio

c/1 can affect A,,,. Thus by dressing a wheel differently, its sharpness can be changed, but as long as A,„ can be measured, one can determine

changes in sharpness due to changes in dressing methods. Wheel sharpness will be seen to be

an important parameter in determining the sur- face integrity of ground surfaces.

Workpiece residual tensile stresses are a function of the amount of heat injected by the

grinding process. Various methods are available to determine the magnitude of such stresses

including x-ray diffraction and deformation

analysis.

Some steels, including AISI-E52100, are sus- ceptible to hydrogen embrittlement when immersed in hot hydrochloric acid. The absorbed

Force Intensity: 100 lb/inch Force Intensity: 120 lb/inch

17500 N/m 21000 N/m

Force Intensity: 140 lb /inch

24500 N/m

Force Intensity: 160 lb/inch

28000 N/m

Force Intensity: 180 lb/inch

31500 N/m

Force Intensity: 200 lb /inch

35000 N/m

Workpieces ground under various force intensities for constant wheel sharpness.

97A1001L6VFM

7700 FPM (38.5 m/s)

180 FPM (0.9 m/s)

2.0 in (51 mm)

2.31 in (59 mm)

0.004 in/rev (0.10 mm/rev)

Figure 23

Wheel

Wheelspeed

Work speed:

Wheel diameter:

Work diameter:

Dress lead:

Climb work rotation

Coolant

Work material

Internal grinding:

Hydrochloric acid etch after each grind—10 min

at 150 °F (65 °C)

Cimperial 20:1, 110 psi (7.7 Kg/cm")

AISI 52100 Rc60

439-921 O - 72 - 69
hydrogen causes the surface layer of the work

to become brittle and if residual tensile stresses

exist, the brittle layer will crack. Thus cracking of the workpiece after hot hydrochloric

acid etch, indicates that the grinding operation

had produced tensile stresses.

In order to establish threshold limits for force

intensities beyond which thermal damage will

occur, grinding tests were made on AISIE52100 steel rings, at Rockwell C 60.

The test procedure was to increase the interface force at each work speed and to examine

the workpiece after hydrochloric acid immersion. A representative run is shown in figure

23 which presents photographs of the ground

and etched parts taken at 15 x magnification

through a microscope. As the force intensity

is increased, cracking begins at the edges of

the workpiece and spreads across the piece,

until, at high force intensities, cross cracking

(i.e., parallel to the direction of the grinding

grit path) occurs. The limiting force intensity

is defined as that which causes the first appearance of cracking.

Figure 24 shows this critical force intensity

plotted against work speed for two sharpness

values: A,,, equal to 88 x 10-« inVs, lb, (0.323

mm^N"^ s"^) a sharp wheel, and equal to 25 X

10-« inVs, lb, (0.0918 mm^N-^S"^), a dull wheel.

As workspeed is increased, the tolerable force

intensity rises, for each wheel sharpness. The

importance of knowing the degree of wheel

sharpness, or A„„ may be illustrated by a hypothetical grinding operation, using skip dressing

(avoidance of dressing for some time). With

work speed at 500 sfm, (2.5 m/s) a force intensity of 180 lb/in (31 N/mm) is permissible

(see fig. 24) and as long as the wheel remains

sharp, thermal tensile stressing of the work

will not occur. However if the wheel dulls (as it will with continual usage) so that A,„ is reduced

to say 25 x 10-« inVs, lb, (0.0918 mm^'-

N-^S-i) the force intensity of 180 lb/in (31

N/mm) is above the permissible level for that

A,,, sharpness value, and the work will have

thermally induced tensile stresses. To be con- servative and set a force intensity below that

permissible for a dulled wheel, here say 120

lb/in, (20.6 N/mm) will cause grinding time to

be lost since the penetration rate, v„„ will be

slow, even with a sharp wheel.

In view of the above the maximum permissible grinding rate can be achieved by grinding

at controlled force intensities while continuously monitoring wheel sharpness ^ thereby maintaining surface integrity.

5. Conclusion

The development of the technology of grinding through the use of material and abrasive

'Patent App. No. 717853

«/t 30

o

o

1 ' > .

/

17.7 All-—A— 6 ( 0.1

WORK SURFACE SPEED

Figure 24. Threshold force intensity to produce

thermal cracking.

m/sec

Wheel:

Wheelspeed

Workspeed

Wheel diameter:

Work diameter

Dress lead

Climb work rotation

Coolant:

97A1001L6VFM

7700 FPM (38.5 m/s)

180 FPM (0.9 m/s)

2.0 in (51 mm)

2.31 in (59 mm)

0.004 in/rev (0.10 mm/rev)

0.0001 in/rev (0.025 mm/rev)

dull

Cimperial 20:1, 110 N/m (7.7

Kg/cm ^) Work material AISI 52100 Rc60

Internal grinding:

Hydrochloric acid etch after each grind—10 min at 150 °F (65 °C)

removal parameters permits the prediction of

grinding performance. The prediction of stock

removal rates, or induced force intensities, of

surface finish, of local and general wheelwear,

and of surface damage effects leads to the ability to rationally design grinding cycles and to

improve those grinding operations currently in

production. The use of interface force intensity as the independent variable allows the grinding

process to be described in terms which are independent of the machine tool rigidity. The use

of contact pressure as an independent variable

permits the evaluation of wheel wear for any

wheelwork geometry or conformity. The development of tabulor data on the material removal

and abrasive removal parameters for various

work materials and wheels will substantially

aid industry in the development of efficient,

economic grinding operations.

70
6. Nomenclature

b = Width of grind

Dk = Diameter of workpiece

D, — Diameter of wheel

Fy = Normal force intensity existing at the _ interface.

Vu, = Radial velocity of the recession of the workpiece surface due to stock removal.

Va = Radial velocity of the recession of the

abrasive wheel due to wheel wear.

Vf = Feed rate of cross slide

y, = Surface speed of grinding wheel

Vu> = Surface speed of workpiece

q = y./y,.

A„ = The Metal Removal Parameter

As = The Wheelwear Removal Parameter

Ic = Length of contact measured along the work

surface.

I = Dress lead: inches of axial motion that the

single point diamond moves per wheel

revolution.

-g" = Half-compensation, or radial depth of dress

d = Diameter of average grain in the wheel

P = Average grinding pressure

Fn = Normal force per unit width—(force

intensity)

Z,' = Volumetric rate of wheel wear per unit

width. = Volumetric rate of stock removal per unit

width.

7. References

[1] Hahn, R. S., On the Nature of the Grinding

Process, Proc. 3rd International MTDR Conference, (University of Birmingham), pp. 129-

154, September 1962 (Pergamon Press).

[2] Busch, Dieter M., Ritz und Verschleissuntersuchungen an sproden Werkstoffen mit einzelkornbestuckten Hartstoffwerkzeugen. Doctoral Thesis, Technische Hochschule Hannover,

Hannover West Germany (1968).

[3] Buttner, Axel, Das Schleifen Sprodharter Werkstoffe mit Diamant-Topfscheiben unter besonderer berucksichtigung des Tiefschleifens.

Doctoral Thesis, Technischen Universitat

Hannover, Hannover, West Germany. (1968).

[4] Pahlitzsch, G., Schleifen oxydkeramischer Werkstoffe. Annals of the C.I.R.P. XXXIII, (Pergamon Press, 1966), p. 331-338.

[5] Ida, Aral, Inamori, Grinding Characteristics of

Ceramics for Circuit Components with Metal

Bonded Diamond Wheels, Annals C.I.R.P. XVII

(Pergamon Press, 1969), pp. 259-268.

[6] Pahlitzsch, G., Ing. Thormahlen, C.I.R.P. Arbeitsgruppe G: Gemeinschaftliches Versuchsprogramm (C.I.R.P. Annual Meeting, Geneva,

Switzerland) ; Technischen Universitat, Braunschweig, August (1969).

[7] Lindsay, R. P., Dressing and its effect on grinding performance, Society of Manufacturing

Engineers, Paper No. MR69-568, S.M.E., Dearborn, Michigan (1969).

[8] Lindsay, R. P., Hahn, R. S., The effect of parameter variations in precision grinding; Paper

No. 70-PROD-l, American Society of Mechanical Engineers (ASME), New York (1970).

[9] Hahn, R. S., Lindsay, R. P., The Influence of

Process Variables on Material Removal, Surface Integrity, Surface Finish and Vibration

in Grinding, 10th Annual M.T.D.R. Conference, University of Manchester, Great Britain,

September (1969).

[10] Guhring, K., Hochleistungsschleifen, Dissertation,

T. H., Aachen, Germany, (1967).

[11] Hahn, R. S., Lindsay, R. P., On the Effects of

Real Area of Contact and N'^rmal Stress in

Grinding, Annals of C.I.R.P., XV, pp. 197-204,

Great Britain, (1967).

[12] Lindsay, R. P., Stock Removal Rates in Internal

Grinding: A Model of the Process, MSc Thesis,

M.I.T., Cambridge Mass. (1966).

[13] Littman, W. E., The influence of grinding on workpiece technology, ASTME paper MR67-

593, (1967).

[14] Colwell, L. v., Sinnott, M. T., Tobin, J. C, The

determination of residual stresses in hardened

ground steel. Trans. ASME, 77, 1955 1099-

1105 (1955).

[15] Letner, H. R., Residual grinding stresses in hardened steel. Trans. ASME, 77, 1089-1098

(1955).

[16] Letner, H. R., Influence of grinding fluids upon

residual stresses in hardened steel, Trans.

ASME, 79, 149-153 (1957).

[17] Halverstadt, R. D., Analysis of residual stress

in ground surface of high-temperature alloys,

Trans. ASME, 80, 929-939 (1958).

[18] Tarasov, L. P., and Lundberg, C. 0., Detection, causes and prevention of injury to ground

surfaces. Trans. ASME 36, 389 (1946).

[19] Hahn, R. S., The relation between grinding conditions and thermal damage in the workpiece,

Trans, ASME, 78, 807-812 (1956),

Discussion

Gielisse: I wanted to ask Dr, Hahn why do

some of the plots of your metal material re- moval parameters not go through the origin?

Hahn: Well, mainly because of the rubbing

and an inability of the grit to get a fracture

going in the surface.

Brown: When I saw your equation up there

I thought to myself in anticipation, well, that's

great, this is going to be different for different

metals. Then I saw the plot where you plotted

Z versus the F for several metals and they all

seemed to follow along the same plot. I thought

that this was quite amazing. Later you show

the bar chart of A„, versus various metals and I

see that a,,, does vary. Now I don't know what

to use,

Hahn: Well, it does vary with metals. The

bar chart was kind of a rough thing. If you

take metals that are classified "easy to grind

steels", the a„, won't be much different. In other

words, there is not much difference in An, in

4150 steels. There are small differences, but

if you go to different materials like M2 or M50

you can see large differences. There is a tremendous difference in A„- in cast iron and M4,

Brown: So then you'd actually get a differ- ence in slope on the Z versus F chart?

Hahn: Yes,

71
Smith, K. M.: Evidently then, you do have a

hammering action on these wheels as a high

spot comes around. Isn't that true?

Hahn: What do you mean a hammering?

Smith: In essence hammering means that

high spot is hitting ahead of anything else.

Hahn: Well, most of these grinding v^^heels are trued with a diamond. The surface of a

precision grinding wheel runs true within perhaps 5 millioniths of an inch, so there is not

much run out of the grinding wheel.

Smith: Does it stay that way very long?

Hahn: As the wheel wears, the soft areas

in the grinding wheel wear more than the little

hard knots. For example, like a log of wood with

knots in it, the knotty part doesn't wear as fast as the soft adjacent part.

Smith: In other words, you're setting up

strains. As soon as the wheel begins to wear it

will hit the soft spots.

Hahn: That is true.

72
A Basic Study of the Diamond Grinding of Alumina

D. M. Busch and J. F. Prins

De Beers Diamond Research Laboratory

P. O. Box 916, Johannesburg, South Africa

Experiments were performed using single diamond points to cut Coors AD96

alumina, under conditions simulating the grinding of this material. These studies

involved different diamond particles and the forces were recorded during the cutting

action.

Parallel to the single diamond experiments, grinding tests were performed on the same ceramic material using the same types of diamonds with which the single point

studies were done. The grinding experiments and the single diamond experiments were

then compared to get a better understanding of the actual cutting mechanism involved

when grinding aluminium oxide.

The wear of the diamond particles and the damage inflicted on the ceramic

material were studied with the aid of electron-microscopy. A relationship was found

between the forces measured, the deterioration of the bonding material, and the wear

of the diamonds.

Special attention was given to diamonds of the MDA-S type, a blocky, sharp

edged diamond grit, and diamond spheres. These two types of diamonds were compared

in grinding, and showed similar results due to the production of flat areas on the

spheres, and the movement and wear of the blocky particles in the bonding material

to present flat faces to the workpieces.

The grinding experiments also entailed a study of machine vibrations, especially

chatter vibrations and the grinding noise associated with it.

Key words: Abrasion; chatter vibrations; forces in grinding; fracture of alumina

grains; grinding noise; grinding of alumina; scratching of glass; single diamond

cutting of alumina.

1. Introduction

A perusal of the relevant technical literature

on the grinding of ceramic materials reveals

considerable variations in the suggested grinding parameters to be used. For example, if a high stock removal rate is required in the sur- face grinding of ceramics, then the recommended downfeeds vary between 10 jam and 250

jum, depending on the properties of the ceramic

material, and grinding parameters, such as table traverse speed, vv^heel speed, etc. [1, 2, 3, 4,

5].^ The present study is not an investigation

into the optimum grinding conditions for machining a ceramic, but deals more with some of

the basic problems encountered during the

grinding of such a material.

For the purpose of this investigation, experiments were carried out using single diamond

points to cut a high density alumina, (specifically 96% Coors aluminum oxide), under conditions that simulated the grinding of this material. These studies were executed using different types of diamond particles, and the surface

damage inflicted on the ceramic material with

these diamonds was studied with the aid of

electron microscopy. Microscopic investigations

were also carried out to study the wear of the

diamond particles and the effect on the bonding

material in which the diamond is embedded.

* Figures in brackets indicate the literature references at the end of this paper.

Parallel to these experiments, actual grinding

tests were performed on the same material

using wheels containing the same types of grits

with which the single point studies were done.

These grinding tests and the single grit experiments were then compared in order to obtain a better understanding of the mechanics of the

cutting process. In performing the grinding ex- periments due attention was given to machine

vibrations, especially chatter vibrations.

Special attention was given to the De Beers

synthetic diamond type MDA-S, a blocky, relatively sharp edged diamond grit, and to diamond spheres [6] shaped from the MDA-S

(fig. 1). It was reasoned that the behaviour

of these two different types of grit should give

an insight into the mechanism by which the

ceramic material is removed by the diamond

particles ; i.e., whether the single diamond particle removes the workpiece material by a cutting process or, by crushing the ceramic grains

to smaller sizes, and then scraping these smaller

particles from the surface of the material. Thus

the mechanism of the removal process should

influence the wear of the single diamond grit

particle: If the diamond particle removes material mainly by impacting and crushing the

ceramic grains, then one would expect less wear

of the spherical diamonds because these particles possess the greater compressive strength

[7, 8].

73
500/um

I 1

DIAMOND MDA-S DIAMOND SPHERES

60/80 SHAPED FROM MDA-S

60/70

Figure 1. Diamond particles used in grinding and

single diamond experiments.

2. Test Equipment and Conditions

The grinding experiments were executed on

a small surface grinding machine with a horizontal spindle, using peripheral diamond wheels.

Table 1 is a summary of the applied grinding

parameters, the types of grinding wheels and

the work-piece specifications.

The scratch tests, which were done in con- junction with the grinding experiments, were

executed on the Single Grit Testing Machine

which has been described previously. [9]. A

single diamond particle was mounted on the

periphery of a wheel, such that the effective

radius was 63.5 mm. While the wheel rotates

at the selected speed, the ceramic material is

brought to the correct position within a single

revolution, and a cut is then made on the material. During the next revolution the material

is removed from the wheel to ensure that only

a single cut is made.

When the single diamond is in contact with

the ceramic surface the normal and tangential

forces associated with the cut are sensed by

means of pressure sensitive transistors, the

signals from which are displayed on a double

beam oscilloscope. Figure 2 is an example of

such a recording, in which the time basis is

100 /iS. per cm. The average force is estimated

by dividing the area under the curve by the total

time of cutting. As the force curves follow approximately a Gaussian distribution, the aver- age force is taken to be about 0.4 times the

maximum force as indicated in figure 3.

For convenience, all the force values obtained

from the scratch tests reported in this paper

are given in terms of the maximum forces,

Fmax* From the cutting time r (fig. 3) and the

known wheel speed, it is possible to calculate

the length of cut, Lt, and the maximum depth

of cut, hr. The maximum depth of cut, hr, or the

theoretical depth of cut was taken as equal to

the infeed, and force measurements were re- corded as a function of this theoretical infeed.

Figure 2. A recording of the normal (top) and tan- gential (bottom) forces on the single grit machine.

Time scale 100 /xs/cm.

Table 1. Test conditions (grinding)

Machine parameters Diamond wheels Workpiece

Surface grinding machine Wheel type DlAl Coors alumina AD96

(Horizontal grinding s spindle drive motor

pindle) with 3 hp Wheel size 127mm X 4.7mm X 31mm

Bond depth 3.1 mm

Specific gravity typical 3.72

Spindle revolution 3600 rpm Diamond types : De Beers MDA-S- Minimum 3.67

Wheel speed 24 m/s (MC) and diamond spheres shaped Crystal size range : 2-20 fim

Downfeed 50 ytim; 125 /j.m. from MDA-S

Crossfeed 1.25 mm. per table

reversal

Diamond mesh size 60/80 and 60/70

Diamond concentration

Modulus of elasticity:

3X10' kg-force/mm

Table traverse speed 16 m/min 100 (=4.4ct/cm') Compressive strength

Stock removal rate 1cm Vmin ; 2.5 Vmin

Bond Bronze 80/20; Resin R-grade typical 25 °C 210

kg-force/mm

74
r"i* 1 Revol Grinding Wheel

Trued Diomond Wheel

( MDA-S)

Figure 3. The relationship between the

average force and the maximum force on a gaussian force curve.

The conditions under which the single grit ex- periments were performed are given in table 2. In the grinding experiments, forces were

measured with a strain gauge wired dynamometer [10]. Displacement and acceleration

transducers and a frequency analyzer were also

used to determine the chatter vibrations during

grinding. The grinding noise was recorded on

a tape recorder.

3. Experiments and Results

3.1. Preparation of the Grinding Wheels

and Single Diamond Tools

Before the diamond wheels were used in

grinding, they were dynamically balanced and

then trued while rotating on the spindle of the

grinding machine. The latter operation was

achieved with the aid of a diamond block (with

a diamond concentration of 200) and a silicon

carbide block dresser. After trueing, the eccentricity of the wheels was checked using an

accurate displacement transducer on the

periphery of the wheel [14] . The profiles of the

circumferences of the metal bonded wheels after

trueing are shown in figure 4. After the trueing process, the number of

- 1 Revol. Grinding Wheel

1 —

Trued Diomond Wheel

A /

-

Di imond Spheres

e'

Oi

Figure 4. The circumferential waviness of the

bronze bonded diamond wheels after trueing.

diamonds visible in the bond was counted on ten

different areas on the wheel circumference [11]

The counting was carried out with the aid of a

stereoscopic microscope containing a grid in

one of the eyepieces. The average number of

diamonds in the square field (2 mm. x 2 mm.)

was determined.

In table 3, the average number of diamonds

per 2 mm square on the wheel circumference

is given for the two different types of diamond

wheels after trueing, after a stock removal of

925 cm^ at a downfeed of 50 fjm, and after a further stock removal of 575 cm^ at a downfeed of 125 fxm. From this table it can be seen

that the bronze-bonded diamond wheel containing the MDA-S type crystals possessed 20 to

24 percent more diamond particles than the

wheel containing the diamond spheres. This percentage difference did not alter appreciably

during grinding and, as will be seen later, this

Table 2. Test conditions (single diamonds)

Machine parameters Diamond types used Workpiece

Single grit testing machine Diamonds mounted in hollow screw heads

using resin.

Same as in table 1.

Spindle revolution

S600 rpm Screw with diamond then mounted on a Wheel Speed 24 m/s mild steel wheel of diameter 127 mm.

Downfeed

Increasing in steps from Diamond types

about 10 ix.m to about 90

flTXV.

De Beers

MDA-S and diamond

spheres shaped from

MDA-S

No crossfeed: Table stationary during

cutting process.

Diamond size: —'220 ixvc\ diam.

[

75
Table 3. Number of diamonds per 2 mm X 2 mm field on the bronze bonded diamond

wheels

Time of counting MDA

After trueing 25

After a stock removal of 925 cm ' at 50

/im downfeed 24

After a further stock removal of 575 cm ' . at 125 fim downfeed 25

difference in number of diamonds is a necessary

prerequisite to explain some of the grinding

results.

The single diamond points were mounted in a little cup using resin as the bonding material.

The cup was machined to give it a threaded

stem which screwed into the periphery of a mild steel wheel. The distance from the centre

of the wheel to the diamond tip was kept at

63.5 mm. Figure 5 is a scanning electron micrograph of the cup with the diamond protruding

from it. After the diamonds had been mounted,

the excess of bonding material surrounding

the single diamond was cautiously removed

with the aid of a small file and steel tweezers,

to allow the diamond to protrude above the sur- face of the bond. Scanning electron micrographs

of a diamond sphere and a blocky MDA-S particle, which are mounted and ready to cut, are shown in figure 6.

3.2. Wear Behaviour of the Bronze-bonded

Diamond Wheels, and the Single

Diamond Points

A relatively high downfeed was applied during the grinding experiments (50 ^am and 125

ju.m) in order to subject the diamond particles

in the grinding wheels and the workpiece materials to high cutting loads. (Stock removal

rates were 1 cmViT^in and 2.5 cmVmin respectively).

Figure 5. Single diamond tool as seen from the

side with the diamond protruding.

-S Diamond spheres Difference

percent

19 24

19 21

20 20

SPHERICAL

DIAMOND

POINT

MDA-S

DIAMOND

POINT

Figure 6. Scanning electronmicrographs of two

single diamonds mounted and ready for

scratching tests.

The wheel containing MDA-S diamonds gave

G ratios of the order of 17,000 to 20,000 which

seemed to be almost independent of the downfeed (50 fxm and 125 fxm). After a stock re- moval of 1500 cm% a wear of about 40 ;u,m was

observed on the grinding wheel.

The grinding wheel containing spherical diamonds gave a grinding ratio of 15,000 which is

about 22 percent lower than the G ratio obtained with the wheel containing MDA-S. In

view of the difference in the number of diamonds on the periphery of the two types of

wheels, which is given in table 3 as approximately 20 percent, the two types of diamond

wheels gave very similar grinding ratios when

interpreted in terms of diamonds retained in

the matrix.

In order to observe the possible wear of diamonds in the grinding wheels specific areas,

and diamonds, on the wheels were replicated

successively after different amounts of stock

76
removal, and observed in the scanning electron

microscope [15]. The results are shov^^n in

figures 7 and 8 for the MDA-S and spherical

diamonds respectively.

For comparison, the single diamond tools

were also observed in the scanning electron

microscope after scratching. The results are shown in figure 9. It has to be remembered that

the scratch tests were done without using a coolant on the surface while the actual grinding

experiments were done using a coolant.

For both the grinding experiments and the

single grit tests, the diamonds quickly developed wear flats. No pronounced break-out of the

edges of the diamond particles was observed.

Although cracks were observed in some of the

diamonds, wear of the diamond particles took

place mainly by abrasion. The diamonds that

did crack during grinding had already been

cracked during the wheel trueing process.

The wear was slightly more pronounced on

the spherical diamonds than the blocky ones,

due to the fact that the blocky diamonds turned

in the bond and thus presented new surfaces

towards the workpiece material, (compare fig- ures 7 and 8). The turning of the MDA-S

£

3., O

o

un

AFTER TRUEING

STOCK REMOVAL 620 cm^

STOCK REMOVAL 920cm^

STOCK REMOVAL ISOOcm^

77
AFTER TRUEING

Figure 8. Wear behaviour of the bronze bonded diamond wheel containing diamond

spheres {shaped from MDAS). Workpiece: Coors AD96.

crystals was even more pronounced in the single

grit tools, and in these tools the blocky particles

were usually lost before a sufficient number of

scratches could be obtained to produce pronounced wear. In contrast, the spherical diamonds stayed in the scratching tools and developed marked wear, with pitting, after a

large number of scratches (see fig. 9). This

pitting was seldom found on the diamonds in

the grinding wheel, and might be due, in the

single grit tools, to an absence of coolant on

the workpiece surface. The pitting had some

effect on the forces measured during cutting

(see para. 3.3).

It is further clear from figures 7 and 8 that

the first diamond layer on the grinding wheels

did not wear away by the time 1500 cm^ of

stock has been removed. This is to be expected

because of the low wheel wear (about 40 /xm),

and because of the grit size of the diamonds

(177—250 /xm). On the other hand, the wear

of the bronze bond was high enough to make

redressing unnecessary.

78
Table 4. Normal and tangential forces when grinding with bronze bonded diamond

wheels

Downfeed Diamond type Normal Fn

(kg Force)

Tangential Fx

(kg-Force)

Difference between the

two wheels

For Fn For F,

r MDA-S 11.8 1.2

1 125 fim J. 18% 17%

I Spheres 9.7 1.0 J

r MDA-S 5.0 0.46

50 < 24%

[_ Spheres 3.8 Not measurable

Figure 9. Wear of single diamond tools during scratching of Coors

AD9 6 alumina.

DOWNFEED 50/jm

Figure 10. Measurement of the normal forces

exterted by the bronze bonded diamond wheels.

DOWNFEED 125Mm DOWNFEED SO/um

Figure 11. Measurement of the tangential forces

exerted by the bronze bonded diamond wheels.

79
3.3. The Forces Involved During Grinding

and Scratching

The normal and tangential forces for the

grinding wheels and the single point tools were

measured. For the grinding wheels these forces

were measured after a stock removal 1500 cm^.

The normal and tangential forces recorded

for the two types of wheels used are shown

in figures 10 and 11 respectively [10]. In order

to measure the relatively low tangential forces,

the sensitivity of the carrier frequency amplifier had to be increased. This caused the influence of the table-reverse on the dynanometer

to be recorded with the force recordings. To

separate the two signals, an output force was

recorded while the table moved without the

wheel contacting the workpiece surface (see

fig. 11). When the downfeed was changed from

50 /xm to 125 ju,m, the forces increased cor- respondingly.

The average normal and tangential forces,

listed in table 4, show that the forces exerted

by the grinding wheel containing MDA-S are

17 to 24 percent greater than those developed

by the wheel containing the spheres. This result

can again be explained in terms of the different

numbers of diamonds in the two wheels as shown in table 3. The magnitude of the forces

is small compared to the forces measured when

tungsten carbide is ground under similar grinding conditions. Another surprising feature is

the ratio of approximately 10:1 measured for

the normal forces relative to the tangential

forces.

For the single diamond points, the forces

were measured as a function of the theoretical

depth of cut, hr. In figure 12 the curves are given for the normal force iF„), tangential

force (Ft) and the ratio of the normal force

to the tangential force for the spherical diamond and a blocky MDA-S diamond. These

curves are quite constant from one diamond

to another of the same type, but are different

for the spherical and blocky diamonds. The

blocky MDA-S crystals, however, turned in

the bond during cutting, causing the force

curves and ratio of the forces to be more uncertain than the measurements on spherical

diamonds. From the curve of the ratio Fn/Ft

for the spherical diamonds it is clear that the

ratio of the forces is of the correct magnitude

when compared to the force measurements

on the diamond grinding wheels. Although the

ratio increases with higher infeeds, to values

over 20:1, for the single diamond, the increase

would not be so marked when higher downfeeds are used with a grinding wheel. If the downfeed with a grinding wheel is 125 /xm, not all

the diamonds making contact with the ceramic

surface will have depths of cut of 125 /i.m.

Maximum Norrral Force as a function of

tfie theoretical depth of cut

( Kg -force

)

Diamond Sphere

MDA-5 Diamond

Infeed m

20 AO 60 80 100

Maximum Tangential Force as a ftnction of the theoretical depth of cut &8

06

( Kg -force)

OA

0-2

MDA-S Diamond

Diamond Sphere

Infeed ^m

20 AO 60 80 100

Thie Ratio of the Normal Force to the

Tangential Force

Diamond Sphere

MDA-S Diamond

Infeed um

20 40 60 80 XX)

Figure 12. Normal and tangential forces, and their

ratio as measured when using single diamonds on Coors AD96 alumina.

There will be some diamonds with lower infeeds.

The curve for the MDA-S diamond tools is

low due to the fact that the diamond turns

during cutting. This causes the diamond to be

more easily pressed into the bond. The bond

is then effectively more elastic, causing the

actual depth of cut to be less than the theoretical depth of cut. A lower normal force is thus

measured.

If the forces are compared for the fresh

spherical diamond surface, shown in figure 6,

with the worn pitted spherical diamond surface, shown in figure 9, the results are widely

different as shown in figure 13. The normal

80
Maximum Normal Force as a function of

the Itieoretical deptti of cut

8

20 ^0 60 80 100

Maximum Tangential Force as a function of

the theoretical depth of cut

0-8

20 AO 60 80 100

The Ratio of the Normal Force

to the Tangential Force

Infeed ^m

20 «) 60 80 100

Figure 13. Normal and tangential forces, and their

ratio when a freshly mounted diamond is cutting

compared to these forces when the same diamond is worn.

force is lower for the pitted diamond than for

the fresh diamond. This is probably due to the

fact that the pitted diamond sits very loosely

in the bond as can be seen in figure 9, causing

the bond to be more "elastic" and having the

same result as discussed above. On the other

hand the tangential force is higher for the

pitted, vi^orn diamond sphere than for the fresh

diamond sphere. This result is expected owing

to the roughness of the pitted diamond. The

ratio between the normal force and the tangential force is thus considerably smaller for the

worn, loose diamond than for the freshly

mounted diamond.

The last result has some significance on the

results found for the grinding wheels. Some

diamonds are more worn and have loosened

to a greater extent in their bond than others,

thus lowering the ratio of the normal force

to the tangential force compared to the better

bonded, less worn, diamonds. The average ratio

of the normal to the tangential force will thus

be somewhere between the two ratio curves in

figure 13, and a ratio of 10:1 as obtained is

thus understandable. For a wheel that is less

worn than the wheels used for the force measurements in this study even higher ratios could

be expected for high dovmfeeds.

3.4. Development and Determination of

Chatter Vibration When Grinding a Ceramic Material

During the grinding experiments with the

bronze-bonded diamond wheels, the wheel profile of each wheel was measured after different

amounts of material had been removed. The

increase in waviness of the wheels as a function of stock removal could thus be determined

and is represented in figure 14.

The waviness increases with the amount of

stock removed, and thus, with the time of

grinding. From this it follows that even when

grinding ceramics, chatter vibrations develop.

The origin and development of chatter vibrations have been investigated previously, when

diamond wheels were used on cemented tungsten carbide [12,13,14]. Vibrations frequently occur on the machine tool when grinding, and

lead to a relative motion between the grinding

wheel and the workpiece. These vibrations

which cause visible, as well as macroscopically

invisible, chatter marks on the ground workpiece surface are called chatter vibrations.

Chatter vibrations are responsible for deterioration of the workpiece surface, increased

grinding wheel wear, excessive machine tool

0 100 200 300 iOO 500 600 700 800 900 1000

STOCK REMOVAL [cm") ^

Figure 14. The development of waviness on the cir- cumference of a bronze bonded diamond wheel con- taining MDA-S—diamonds.

81
Rcvol. Grind ng Wh

t0 01 sec

CHATTER VIBRATION AMPLITUDE DISPLACEMENT

GRINDING WHEEL WAVINESS

10

Frequency Spectrum

Grinding Noise

1

Frequency [Hz ]

Figure 15. Curves showing the amplitude of

chatter vibrations, the waviness of the wheel

and the frequency spectrum of the grinding

noise when grinding with a resin bonded

MDA-S diamond wheel at a downfeed of

50 fim.

wear and a reduced stock removal rate.

These vibrations were recorded when using

a resin bonded wheel containing metal clad

MDA-S diamonds of size 60/80 mesh. Figure

15 shows an example of the recorded chatter

vibrations, from which the vibration amplitude

can be derived, the development of the waviness

on the circumference of the wheel, and a frequency spectrum obtained from a recording

of the grinding noise.

The amplitude of the chatter vibrations was

found to increase with the grinding time. With

the development of the chatter vibration amplitude a waviness occurs on the grinding wheel

circumference, the depth of which also increases

with the grinding times.

From previous investigations [12,14] it was

established that the chatter vibrations were

caused by a self-excited vibration, the frequency

of which depends on the natural frequencies,

the rigidity, the damping coefficients and the

coupling stiffness of the grinding machine

grinding wheel—workpiece system. Chatter

vibrations can also be reinforced by forced

vibrations which stem, for example, from imbalances of the grinding wheel or other rotating

machine parts, damaged and misaligned bearings, etc. When comparing chatter vibrations on alumina and cemented tungsten carbide a retarded

development of these vibrations was found for

the alumina. This slower development of the

chatter vibrations is the result of the lower

cutting forces necessary to grind alumina.

3.5. Influence of the Grinding Wheels and

Single Diamond Tools on the Surface

of Coors AD96 Alumina

After grinding and scratching of the alumina,

replicas were made of the ground and scratched

surfaces and studied in the transmission electron-microscope. The ground surfaces could

then be compared with the as-sintered un- ground surfaces.

In the case of the scratching experiments

the preparation of the surface to be scratched

presented some difficulties. The ceramic blocks

required for mounting on the single grit machine must have the dimensions of 10 mm X

8 mm X 4 mm. In order to obtain these size

blocks, the ceramic material had to be sawn

and ground to the correct size. The ceramic

blocks thus had an as-ground surface and not

an as-sintered one. Polishing of the surfaces

with a diamond powder did not improve matters at all, because during polishing, the glassy

binder material melts and is smeared over the

surface rendering a totally different surface

than the one seen by the diamond during grinding (see fig. 16). If the polished surface is

etched in molten sodium hydroxide to expose

Figure 16. A polished alumina surface.

82
Figure 17. An etched, -polished alumina surface which was then scratched.

Figure 18. A scanning electronmicrograph of cuts made in alumina.

the underlying grains, the situation is not improved. Figure 17 shows an etched polished

surface on which a single scratch was made.

It is clear that the polishing process causes profuse fracture of the surface grains with

subsequent reduction in grain size. A scratch

on such a material thus only removes the fractured top layer of material and exposes the

larger underlying grains (fig, 17). It was thus

decided to use an as-ground surface obtained

by using a wheel containing DXDA-MC grits

of size 140/170 in an R-grade bond and to compare the cuts on these surfaces with the as- ground surface.

Viewing the cuts in the scanning electron

microscope before etching reveals that some of

the glassy binding material was extruded during cutting and smeared over the surface. This

effect was more pronounced for the spherical

diamond than for the MDA-S diamond. Figure

18 shows two cuts of theoretical depth about

60 /xm, one made with a diamond sphere and

the other with a blocky MDA-S particle. More

extrusion and smearing of the glassy material

is found for the diamond sphere than for the

blocky particle, and the cut obtained with the

sphere is obviously deeper than the one obtained

with the blocky particle. This r'^sult is expected

when compared with the fore ; measurements.

When replicas of the scratched and etched

surfaces are viewed in the electron microscope

the difference between the scratched part and

the as-ground part is not very pronounced. This

is to be expected because the grains in the

ETCHED

AS GROUND SURFACE

ETCHED

SURFACE IN CUT

Figure 19. Grains in a cut and on an as-ground

surface using a spherical diamond.

Figure 20. Abrasion marks on aluminium oxide grains.
DIAMOND MDA-S (60/80) DIAMOND SPHERES (60/70)

UNGROUND SURFACE

GROUND SURFACE WITH DOWNFEED 125/um

GROUND SURFACE WITH DOWNFEED 50/um

WORKPIECE AD 96

Figure 21. Replica's of etched aluminium oxide surfaces comparing the

ground surfaces with the unground surfaces.

scratch should fracture in the same way as the

grains on the as-ground surface. No difference

could be found between the surface and scratches on the specimens which were scratched with

the blocky diamonds. However, in the scratches

made with the spheres, the grains in the

scratches are a bit more fractured than the

grains on the as-ground surface (see fig. 19).

This result is also expected from the force

measurements. Abrasion of the aluminum oxide

grains can be seen clearly on some of the

larger grains. An example is given in figure 20.

A comparison of the ground and etched sur- faces produced with the bronze-bonded diamond wheels, and the as-grown surface of the

aluminum oxide is given in figure 21. On all

the ground surfaces fracture of the grains

is evident, but the fractures did not alter appreciably the size distribution of the aluminum

oxide grains. Some of the grains have large

fractures which nearly cleave them, but these

large fractures were not usually accompanied

by the break-out of one part of the grain. Most

of the exposed fracture surfaces show that

small splinters of the ceramic grains were re- moved. These fracture surfaces are usually

confined to the tops of the ceramic grains, and

rarely extended further than the depth of a

single layer of the grains. On all surfaces abrasion marks could be detected, especially on the

larger grains. It is interesting to note that the

grains exposed after grinding with a downfeed

84
of 125 jxm, are more closely packed, with the

wheel containing diamond spheres, and have

a flatter appearance, thus giving a smoother

surface.

Both wheels remove material without changing the size distribution of the grains on the

ceramic surface to any marked extent, and the

surfaces are very similar to the as-grown surface, but do have a slightly better texture. Consequently, these wheels are not suitable for a

very fine surface finish, which, according to the

polishing experiments mentioned (figs. 16 and

17), should have finer grains than the as-grown

grains. Both wheels are, however, suitable for

stock removal without inducing too much sur- face damage.

4. Conclusions

Both the spherical diamonds and the MDA-S

diamonds gave comparable grinding results

and remove material from the aluminum oxide

by fracturing and abrasion. The fracturing

occurs as a consequence of the large ratio of

the normal forces to the small tangential force,

and the extent of fracture should thus be a

function of the size of the diamond.

It was also established that the normal force

decreases when the bond deteriorates.

The spherical diamonds and the MDA-S diamonds developed wear flats very quickly when

used in the grinding and scratching experiments. Some of the MDA-S diamonds especially

those with sharp corners protruding, turned in

the bond thus presenting a flat surface to the

workpiece. Thus within a relatively short time

both the spherical diamonds and the MDA-S

diamonds present flat areas to the workpiece

material. It is thus not surprising to find similar grinding ratios for both types of diamonds.

5. Scratching Tests on Glass Surfaces

The comparative study between sharp-edged

blocky diamonds and diamond spheres was derived from a previous study where diamonds

were used to scratch glass surfaces. In those

experiments it was found that, at high loads,

the glass fractured, and, if a sharp-edged diamond was used for cutting, the diamond sometimes also broke at the same time. This was not

as frequently observed for spherical diamonds

because of their higher compressive strength.

It was then thought that if the grinding of a ceramic material entailed fracturing of the

ceramic grains and subsequent removal of these

grains then, according to the tests done on

glass, the spherical diamonds should wear less

than the sharp-edged ones. The previous tests [7] on scratching glass

surfaces were conducted at room temperature,

20° C, and a relative humidity of 65 percent.

CAMERA

MICROSCOPE

MOTION -LEVER

Figure 22. Apparatus used in glass scratching

experiments.

The scratching speeds were 20 ^^m/s and 40

fxm/& respectively. The tests were carried out

with a scratch device similar to the Martens

Scratch—Hardness—Tester (fig. 22). The

spherical polished tools, which were fastened

to a steel lever were pressed against the glass

surface with a known force. The movement to

produce the scratch was executed by the workpiece. For the observations and photographing

of the scratching process an optical microscope

and a 16mm—Ariflex camera were used, the

latter being attached to the top of the microscope. The sequence occurring during the

scratching process on the glass surface was

filmed at a speed of 24 frames per second.

The film of scratching rough spectacle glass

B260 (Deutsche Spiegelglass A.G. Griinenplan)

depict the elastic and plastic behaviour of glass

and the development of fractures [16, 17, 18,

19, 20].

For many years these phenomena which

occurred on the surface during the polishing of

glass could not be explained. It was thought

that there was no essential difference between

the grinding and polishing process.

The elastic behaviour of the glass, when

scratched by the spherical tipped tools, manifests itself in the film by a wave motion on

the glass surface which forms behind the tool

and which disappears again with increasing

distance away from the tool.

At a certain surface pressure of about 350

kp/mm- there is an increasing wave motion

of the glass, which cannot elastically reform

itself when the force with which the tool is

pressed against the glass surface is increased.

Chipping and glass fractures are thereby not

observed so that the plastic behaviour of the

439-921 0-72-7 85
glass by this process demonstrates itself in

microscopic ranges. Figure 23 illustrates the

scratch traces produced during the tests using

a diamond cone having a spherical tip. Walllike mounds can be observed at each side of

the scratch groove and these consist of plastically displaced material.

It can be seen from figure 24 and from the

film recordings that the fractures of the glass

material to form the scratch groove take place

in front of the tool. It seems possible that at

Figure 23. Cross-section of scratches on a

glass surface showing plastic deformation

(wall-like mounds)

I 1 100 jum

Figure 24. Fracture formation in front of

the tool when scratching a glass surface.

A

Figure 25. Flow-chip formation of glass

when cutting with a diamond tool.

certain places on the scratch groove and at

some distance behind the scratching tool, a

break formation also could occur because of

elastic resiliences of the glass material.

The plastic behaviour of glass in microscopic

ranges can be explicitly proved by figure 25

and by the film recordings of the flow-chip

formation of the glass when it is scratched

with a sharp and pointed edged diamond tool

having a positive rake angle. During these

cutting tests a glass chip approximately 8mm

long was produced and did not tear off at the

cutting point.

The chip surface was approximately 2 ^^m

thick and 7 fim wide.

The grinding process can, in principle, be

distinguished from the polishing process on

the basis of the evidence of plastic behaviour

of glass in microscopic ranges. Glass is rapidly

abraded by cracking and breaking out of chips

during grinding. The surface layer thereby

remains continuously broken up for a thickness

of several lengths of light waves. The abrasion

of the material is caused by side cracks, and

the spreading of these cracks and interaction

results in the chip-like pieces of glass peeling

off. The polishing process, in contrast to the

grinding process, is based on plastic deformation on the glass surface in microscopic ranges

which represent a genuine flowing process.

The cracks and indentations caused by the

grinding process are evened out to a fraction

of a wave length.

It is a pleasure to acknowledge the cooperation of the following persons: V. Conradi for

manufacturing the grinding wheels, D. R. Harding for help given in running the single grit

testing machine, J. Fourie for help on the grinding experiments, J. Eckert who made the re- plicas for electron microscopy, and J. Grassmann for the force measurements on the grinding wheels.

6. References

[1] Luce, E.G., Why use diamond wheels when maching ceramics?

[2] Ida, I., Arai, Y., Fukuda, M., Inamori, K., and

Sugiura, M., Relations between material properties and grinding characteristics when using

diamond wheels on ceramics for circuit components, Reprinted from the Review of the

Electrical Communication Laboratory. 17, No.

9, (1969).

[3] Primak, L.P., and Shripko, G.F., The contact

temperature in the surface grinding of high

alumina ceramics, Industrial Diamond Review

30, No. 357, 318-320 (1970).

[4] Sawluk, W., Flachschleifen von Oxydkeramischen

Werkstoffen mit Diamand-Topfscheiben, Dr.-

Ing. Dissertation, Technische Hochschule

Braunschweig (1964).

[5] Gielisse, P.J., Mathewson, W.F., Martis, J.A., and

Rattermann, E., Keramikbearbeitung mit

86
Diamenten, Sonderdruck aus Fachberichte fur

Oberflachentechnik 7, Heft 1/2, 3/4, 5/6 (1969).

[61 Fluid Energy Mills, private communication.

[7] Busch, D.M., Ritz-und Verschleiss-unter=;uchune:en an sproden Werkstoffen mit enzelkornbestiickten

Hartstoffwerkzeugen. Dr.-Ing. Dissertation,

Technische Universitat Hannover (1968).

[8] Kobayashi, A., Recent development in diamond

wheel grinding.

[9] Prins, J.F., and Harding D.R., An apparatus for

testing single diamond particles under simulated

working conditions, Diamond Research 1970, A

supplement to Industrial Diamond Review (Jan.

1970).

[10] Levitt, C, and Grassmann, J., Diamond Research

Laboratory dynamometer uncover grinding

force secrets, Industrial Diamond Review 29,

No. 347, 401-405 (1969).

[11] Biittner, A., Das Schleifen sprodharter Werkstoffe

mit Diamant-Topfscheiben unter bes^nderer

Beriicksichtigung des Tiefschleifens, Dr.-Ing.

Dissertation Technische Universitat Hannover

(1968).

[12] Busch, D.M., Machine vibration and its effect to the diamond wheel. Diamond Information L20.,

De Beers Industrial Diamond Division.

[13] Busch, D.M., and Biittner, A., Effect of machine

vibrations on the grinding result, D.R.L. Report

No. 12/70(0.

[14] Busch, D.M., The development of chatter vibrations when grinding with diamond grinding

wheels. D.R.L. Report No. 16/70 (C).

[15] Eckert, J.D., and Caveney, R.J., A replica technique for conventional and scanning electron

microscopy. Journal of Scientific Instruments 3, No. 5, 413-414 (May 1970).

[16] Busch, D.M., Verformen von Glas mit Ritzwerkzeugen, Film B984 des Inst. Wiss. Film,

Giittingen (1967).

[17] Busch, D.M., Verformen von Glas mit Ritzwerkzeugen—Mikroelastisches Verhalten—Film E

1328 des Inst. Wiss. Film, Gottingen (1967).

[18] Busch, D.M., Verformen von Glas mit Ritzwerkzeugen—Mikroplastisches Verhalten—Film E

1329 des Inst. Wiss. Film, Gottingen (1969).

[19] Busch, D.M., Verformen von Glas mit Ritzwerkzeugen—Bruch—Und Schollenausbildung—Film

E 1330 des Inst. Wiss. Film, Gottingen (1967).

[20] Busch, D.M., Zerspanen von Glas mit Diamantschneidwerkzeug—Spanbildung beim Ritzen

Film E 1331 des Inst. Wiss. Film, Gottingen

(1967).

Discussion

Wiederhorn: Tin curious about two tilings.

One, in the movie there were some waves on

the surface that appeared to be damping out.

What did you attribute these waves to?

Busch: I'm not quite sure yet. It could be

an effect of vibrations on the system. That

could be possible but it might be something

else because we have to expect, let's say, a deformation in front of the tool like a bow

wave.

Wiederhorn: Would you say it was relaxation of a deformation ? Busch: You mean elastic deformation? Yes,

I would say so, but there are some effects which

can't be explained at present.

Wiederhorn: The other question was, what

was the environment? Was it air?

BUCSH : It was in air, 20 °C and humidity 65

percent. No coolant was used.

Greenfield: In the last movie sequence, you

showed a very large helical chip. Did you make

any diffraction patterns of the chip? In other

words was this chip at all crystalline?

Busch: We could not obtain the chip for

examination.

Lange: It appeared to me that the movie

showed that the contact area was much larger,

but yet, the crack or cracks only started at one

point on the perimeter of the tool. Was there

any reason ? Busch: The geometry of the tool was always

the same. It was a cone with a radius of curvature of about 50 yam.

Lange : In other words, what we saw was

actually the point of the tool cutting the glass

and not what appeared to be a spherical contact

area. Busch: Sorry, there were two types of tools;

spherical tools and in this last movie sequence

where chips were produced, it was actually a

cutting tool with special angles. In this sequence

you saw the point of the tool cutting the glass.

In the earlier sequences you saw the projection

of the spherical tool through the glass.

Gielisse: Have any experiments been made

on the effect of change on what one refers to as the rake angle cutting point during the

generation of the glass chips?

Busch: No.

Gielisse: Do you expect this to behave as you showed in the film as the rake angle grows?

Busch: Yes. In these glass investigations a diamond cutting tool was used with a positive

rake angle.

national bureau of STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Sympwsium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (IssueS May 1972).

87

On the Strength of Ceramics as a Function of

Microstructure, Grinding Parameters, Surface

Finish, and Environmental Conditions

Rudolf Sedlacek, F. A. Halden, and P. J. Jorgensen ^

Stanford Research Institute, Menlo Park, California 94025

The tensile strength of four aluminas is described in terms of their microstructure

and composition. The effect of microstructure on the attainable surface finish is discussed, and the results of profilometric evaluation of ground surfaces are presented.

Various postgrind treatments are shown to have differing effects on microstructure.

Griffith's flaw theory is applied to aluminas of different grain sizes, and calculated

strengths are found to agree with experimental data. The phenomena of delayed

fracture and stress corrosion are discussed; the strength of alumina is found to be

higher under vacuimi than in air and is independent of stress rate. Key words: Alumina; delayed fracture; grinding of ceramics; microstructure; stress corrosion; vacuum strength.

1. Introduction

Ceramic manufacturing is a complex process

consisting of a number of separate steps, each

of which can have a pronounced influence on

the physical and mechanical properties of the

final product. Each individual step requires

quality control that cannot always be applied

effectively, and, as a result, the properties of

ceramics show a certain amount of variation.

Considering the high potential of ceramics

for load-bearing applications at elevated temperatures, an understanding of the factors

influencing mechanical properties is most desirable. Particularly, the tensile strength of

these materials is of interest, because it is

relatively low and often shows a disturbing

lack of reproducibility. In recent years, new

tensile testing methods have been developed

which have proved that much of the data scatter found in older literature can be attributed

to the use of faulty testing techniques. These

new techniques now appear to provide data

that can be utilized with much more confidence.

In order to work systematically toward an

improvement of strength of ceramics, it is

necessary first to gain a better understanding

of the factors presently limiting their strength.

These factors cannot all be studied simultaneously. Instead, it appears more practical to

concentrate on those factors that are most

obvious and that can be most easily isolated.

Since the formulation of Griffith's flaw theory, it has been widely accepted that the key

to the strength of ceramics lies in the perfection

of their surfaces. Whenever a ceramic part

has to meet close dimensional tolerances, it is

' Senior Research Ceramist, Associate Director Materials Laboratory, Manager Ceramics Group, respectively.

^ Figures in brackets indicate the literature references at the end of this paper.

shaped in the final step of the manufacturing

process by grinding with bonded diamond. This

operation then determines not only the geometry of the piece, but also the quality of the

surface finish. Obviously, grinding is an important step in the manufacturing process, and it is surprising how little is known about it. It has been shown [1] - that by changing the

grinding conditions, the strength of alumina

can be varied about 10 percent, while the attendant surface finish changes by a factor of

four. It is also known that, for a given grinding

treatment, different alumina ceramics attain

different surface finishes. This means that the

results of the grinding operation are linked to

the microstructural properties of the ceramic

materials. This relationship is discussed in the

following article. Also discussed are the dependence of the strength of alumina on various

post-grind treatments, on stress rate, and on

test conditions such as vacuum.

2. Experimental Procedure

2.1 Test Materials

The alumina bodies used in this study were

Al-300, Al-600, Al-995 (all manufactured by

the Western Gold and Platinum Company, Belmont, California) and AD-995 (product of the

Coors Porcelain Company, Golden, Colorado).

These materials were chosen in the hope that

the correlation of different chemical compositions and microstructures (see fig. 1) with the

grinding procedure, surface finish, and various

post-grinding treatments would indicate—in

terms of strength—the overall problems of

diamond finishing of ceramics.

The composition of the four bodies, determined by semiquantitative emission spectrochemical analyses are shown in table 1, and

89
Table 1. Emission spectragraphic analyses of alumina bodies

Amount in each body, wt percent

FllpTYiPnt AD-995 Al-300 Al-600 Al-995

Al Principal constituent in each sample

Si 0.25 1.25 2. 0.4 Mg 0.25 0.02 0.75 0.5 Ca 0.04 0.6 0.3 0.08

Na <0.02 <0.02 <0.02(?)

Cr 0.2 0.001 0.002

Fe 0.08 0.07 0.12 0.12

Ga 0.02 0.015 0.01 0.015

Ti 0.01 0.025 0.02 0.025

Zr 0.02 0.006 0.015 0.005

B <0.01 0.15 <0.01

Mn <0.001 <0.001 0.001 0.001

Pb <0.002 <0.002

Cu 0.0005 <o.6o65 <0.0005 o.'ool Ba 0.002 0.001 0.001 0.002

Ag less than 0.0005% in each

Sr less than 0.005% in each

V less than 0.005% in each

Table 2. Physical properties of the test materials

Material

Density

kg/m ' X 10 ' Average grainsize m X 10 -° Color

Al-300 3.75 32 Bluish white

Al-995 3.80 18 White

AD-995 3.75 13 Pink

Al-600 3.71 10 White

some of the physical properties of the four

test materials are presented in table 2. The

aluminas designated with the prefix Al were

white, although each type was of a slightly

different shade of white. The pink color of

AD-995 alumina varied considerably from piece

to piece, but no correlation between color and

strength was observed.

2.2. Grinding Procedure

The general grinding procedures used in the

study have been described in detail elsewhere

[1], and the grinding variables were held constant during this investigation. The faces and

0. D. of all specimens were ground with the

320 grit wheel (D-320-N-100 Ml/16) and the

1. D. was ground with the 120 grit wheel (D120-N100M1/16). The same grinding conditions were used in both operations:

Surface speed

Infeed

Table travel

Spark-out

Figure 1. Microstructure of test materials.

30 m/s (5700 sfpm)

6.4 X lO " m/pass

(0.00025 in./pass)

5.1 X lO-^* m/s (1 ft/min)

180 s All specimens, except the Al-600 group, were

ground to the same dimensions:

Inside diameter 5.08 x 10 ^ m (2.000 in)

Wall thickness 2.54 x 10^^ m (0.100 in)

Height 7.6 x 10"^ m (0.300 in)

The Al-600 blanks were ground to 2.0 x

10-3 m (0.080 in) wall thickness, 9.5 x 10-^ m

90
(0.375 in) height, and 5.08 x 10-^ m (2.000 in)

internal diameter. After grinding, the specimens were degreased in trichlorethylene vapor,

thoroughly washed in detergent, and rinsed

with hot water. Finally, they were dried for 1 h

in an oven at 150 °C, sized, and stored for

testing or for postgrinding treatments.

2.3. Testing Method and Apparatus

The ultimate tensile strength of all specimens

was determined by the expanded ring test

method developed at Stanford Research Institute. A description of this method can be found

in the open literature [2] and will not be re- peated here. All specimens were loaded to fracture at the rate of 2.1 x 10« N/m^s (3000

psi/s)

2.4. Postgrinding Treatments

The purpose of this work was to determine

whether or not a particular set of conditions

to which alumina ceramics are exposed in storage and/or in testing may be responsible for

a decrease in strength or some abnormal trend

in test results. In this context, the influence of

water on the strength of alumina is of particular interest. There is general agreement among

workers in the field of strength properties of

ceramics that the presence of water is detrimental to strength. However, opinions vary

as to the mechanism by which water affects

the strength. Some authors [3,4] believe that

hydrates of alumina are formed on long ex- posure to water, (liquid or vapor). The majority of workers prefer the explanation that

in the presence of water a stress-enhanced re- action takes place at the tip of the flaws which

results in high localized stresses. The following

postgrinding treatments were carried out in

an attempt to clarify the part which water plays

in influencing the strength of ceramics.

a. Treatment 1 Specimens were heated under a vacuum of

13.3 X 10-* N/m^ (1 X 10 * torr) and the temperature was increased linearly in 4 hours

to 1000 °C where it was held for 1 h. Then the

temperature was lowered again linearly for 4

hours, until it reached approximately 300 °C;

below this temperature cooling proceeded more

slowly. While the specimens were still warm,

50 to 60 °C, the vacuum was broken, and the

specimens were removed from the furnace and

tested in air as rapidly as possible. The vacuum

was expected to decompose alumina hydrates,

if present, and drive off physically absorbed

water. The short time for which the specimens

were exposed to atmosphere during testing

should not permit the formation of new hydrates and the absorption of water on the sur- face may not go to completion.

b. Treatment 2

Vacuum-baked specimens were placed for 2

weeks in desiccators over a saturated solution

of potassium carbonate. At the room temperature of 23 °C the relative humidity in the desiccators was 43 percent. This treatment was

designed to show whether prolonged exposure

to water vapor would cause alumina to revert

to the prebakeout conditions.

c. Treatment 3

Vacuum-baked specimens were immersed for

1 week in distilled water. Before testing, the

specimens were thoroughly dried with paper

towels and left to equilibrate for 1 h under

ambient conditions. Only Al-300 and Al-995

aluminas were given this treatment. The objective here was the same as that of treatment

2, except that liquid water was substituted for

vapor.

d. Treatment 4

Ground specimens were stored in water for

1 week, then they were dried and left to equilibrate for 1 h under room conditions. The results

of exposure to water of as-ground and vacuumbaked specimens were compared.

e. Treatment 5

As-ground specimens were placed—one set

at a time^—in a closed jar provided with a

thermometer and cooled in an ice bath slightly

below the dew point. After 15 min below the

dew point, the specimens were taken out of the

jar, one by one, and tested as rapidly as possible.

f. Treatment 6

As-ground specimens were immersed in distilled water for a few minutes, then they were

taken out, the excess water was shaken off,

and the specimens were tested while they were

still covered with a visible film of water. In

treatments 5 and 6 the effect of liquid water

on the surface of alumina under stress was

appraised.

3. Results and Discussion

3.1. Evaluation of Ground Surfaces

The texture of the ground surfaces was evaluated using a Clevite Surfanalyzer 150 System.

All measurements were made with the cutoff

width of 7.6 X 10"* m (0.030 in) and stylus

speed of 2.54 x 10 * m/s (0.01 in/s). The re- sults are presented as roughness (arithmetic

average, AA) and profile (defined by Standard

B46. 1-1962 of the American Standards Association).

Profile traces of the four materials ground

identically are shovni in figure 2. On these

91
graphs each vertical division represents 5.08

X 10'^ m (20 |U.in), and each horizontal division

equals 2.54 X lO"'^ m (1000 jain). The subscripts

under the profile traces show the corresponding

AA roughness values. It can be seen that there

is little difference between the profile traces

and the AA roughness values for the Al-300

and Al-995 aluminas, notwithstanding the sizable difference in grain size (32 x lO *^ m and

18 X 10^^ m, respectively). However, the sur- face of the AD-995 material is considerably

smoother, and Al-600 has the best finish of all. In fact, the Al-600 finish is even slightly better

than the finish obtained on Al-995 alumina,

when ground with 1200 grit diamond [1].

Since all four materials have the same grinding

history, it appears that the surface finish obtained in a given grinding operation is largely

controlled by some inherent materials property,

e.g., grain size and/or grain boundary strength.

3.2. Strength Measurements

Because of the relative nonuniformity of tensile strength data generally obtained in testing

ceramics, any attempt to compare results must

include a careful examination of all factors

that can influence the strength. Moreover, these

factors are usually interdependent, so that it is

difficult, if not impossible, to isolate the effect

on strength of a single factor. The factors

studied in the work described were: microstructure (typified by grain size) ; to a lesser

extent chemical composition of the test materials; and the various postgrinding treatments described in the preceding section. The

test results are summarized in table 3. The

data lend themselves to comparison of microstructure effects in several ways. The easiest

comparison can be made between groups having

no postgrinding treatment. From table 3 it is seen that the average tensile strength values

(groups A-1, B-1, C-1, and D-1) cover a broad

range:

Al-300-19.2 X lO^N/m^ (27.8 x 10^ psi)

Al-995-21.6 X 10'N/m= (31.4 x 10^ psi)

AD-995-24.3 x lO^N/m^ (35.3 x 10^ psi)

Al-600-27.3 X lO'N/m^ (39.6 x 10^ psi)

The individual group strength values are

distinctly different. There is no overlap of

standard deviations, and the results of the

statistical analyses indicate that these strength

values are significantly different on the 1 percent confidence level. The increase in strength

is accompanied by a decrease in grain size (see

table 2) and improvement of surface finish

(fig. 2).

The clear-cut difference between the strengths

of the four test materials prompted a cursory

comparison of the above experimental results

with theoretical strength values derived from

Griffith's theory of failure of brittle materials.

Griffith's criterion for failure enables one to

predict the strength of polycrystalline materials on the basis of known or estimated parameters, i.e., E (Young's modulus), y (fracture

Table 3. The effect of post-grinding treatments on the tensile strength of alumina

Group Material Treatment

Avg. tensile strength Std. deviation

N/m= X 10' psi X 10' ±N/m' X 10" ±psi X 10' A-1 Al-300 None 19.2 27.8 7.6 1.1 B-1 Al-995 None 21.7 31.4 12.6 1.8 C-1 AD-995 None 24.3 35.3 12.4 1.8 D-1 Al-600 None 27.3 39.6 14.7 2.1

A-2 Al-300 1 20.3 29.4 8.5 1.2 B-2 Al-995 1 23.0 33.4 7.2 1.00

C-2 AD-995 1 23.2 33.6 7.2 1.0 A-3 Al-300 2 20.3 29.5 11.4 1.7 B-3 Al-995 2 22.5 32.7 9.5 1.4 C-3 AD-995 2 23.2 33.6 9.6 1.4

A-4 Al-300 3 19.6 28.4 8.3 1.2 B-4 Al-995 3 22.6 32.8 8.1 1.2 A-5 Al-300 4 19.1 27.7 5.2 0.8 B-5 Al-995 4 21.2 30.8 7.5 1.1 C-5 AD-995 4 23.8 34.5 8.8 1.3

A-6 Al-300 5 17.4 25.2 6.6 1.0 B-6 Al-995 5 17.8 25.8 4.2 0.6 C-6 AD-995 5 22.8 33.0 11.4 1.7 A-7 Al-300 6 14.6 21.2 5.0 0.7 B-7 Al-995 6 17.0 24.6 5.1 0.7 C-7 AD-995 6 19.7 28.5 11.9 1.7

92
^}:|t|#it1±[t| ! iH ' : ; M

I i-!mtlii|Hi-Hi+H-|+

m^mmmmmmM ROUGHNESS (AA) 16-20

Figure 2. Profile traces of surfaces

of various aluminas ground with

120 grit diamond.

surface energy), and L (maximum length of

crack).

The predicted strength; S, is given by the

following relationship:

S

1/2 / 2 A E\ -

\ )

(1)

For the sake of expediency the following

assumptions were made:

£7 = 37.9 X 10^" N/m- (55 X 10« psi) for all

four materials.

V = 7.0 J/m^ (see ref [5])

L = twice the average grain size when steady

state has been achieved following nor- mal grain growth [6]

Substituting these values into eq (1) for each

of the four aluminas, we obtain the comparison

of experimental and calculated strength data

shown in table 4. The good agreement obtained

tends to support the contention that grain size

rather than surface finish—is the important

factor controlling strength. If one assumes that

in a polycrystalline body the depth and length

of a Grifl[ith-type surface flaw is proportional

to the grain size, then it is not surprising to

find that changing grinding variables over a

broad range (see ref [1]) does not change

strength in proportion to the observed changes

in surface finish. It appears that reasonable

grinding conditions do not alter the density

and severity of Griffith-type flaws, and conse- quently strength variations are relatively small.

Similarly, harsh grinding conditions (deep infeed) lower strength without impairing noticeably the surface finish because the forces ex- perienced by the work piece in grinding are

sufficient to drive existing flaws beyond the

one grain size limit.

3.3. Effect of Various Treatments on the

Tensile Strength of Alumina

a. The Effect of Treatment 1 on the Strength of

Alumina

The maximum temperature of the vacuum

bakeout and the firing schedule were fixed arbitrarily, since no information was found in

the literature as to the optimum conditions cov- ering all four test materials. The results in

table 3 show that the effect of vacuum bakeout

on strength varies widely. The strength differ- ences between the heat-treated specimens and

those tested as ground can be considered to be

the result of the vacuum bakeout on the 1 to 5

percent confidence level.

In Al-300, the bakeout produced a 5.8 percent increase in strength. In Al-995, the in- crease in strength was 6.4 percent. The effect of

the bakeout is reversed for the AD-995 material which has a smaller grain size. In AD995, the bakeout produced a decrease in

Table 4. Comparison of calculated and experimental strength values for

different aluminas

Material

Strength calculated Strength experimental

N/m= X 10' psi X 10' N/m' X 10' psi X 10'

Al-300 17.9 26.0 19.2 27.8

Al-995 21.4 31.1 21.7 31.4

AD-995 25.9 37.5 25.2 36.3

Al-600 28.3 41.1 27.3 39.6

93
strength of 3.4 percent. Fragments of the

bakedout specimens were polished and etched,

and a new grain count was performed. The

bakedout specimens had a significantly larger

grain size, particularly the AD-995 alumina,

(17.3 X 10-« m as compared to 11.7 x 10-« m).

In Al-995 no grain growth was observed. An

increase in grain size from 32 x lO"" m to

34 X 10 *' m was observed in Al-300.

b. The Effect of Treatment 2 on the Strength of

Alumina

In applying treatment 2 it was reasoned

that since the attack of water must start from

the surface, test specimens having different

surface conditions might react differently to

the same atmospheric moisture conditions.

Also, if the action of water was essentially

chemical in nature, materials having different

compositions might show different responses.

Both of these premises were found to be wrong.

Taking the data from table 3 and comparing

groups of specimens given treatment 1 with

those that received treatment 2, we find that the

differences in strength are either negligible or

non-existent. We therefore conclude that ex- posure to atmospheric moisture under the conditions used in this program does not cause any

deterioration of strength.

c. The Effect of Treatment 3 on the Strength of

Alumina

After the vacuum bakeout, a group of Al-300

and a group of Al-995 specimens were immersed in distilled water for 1 week. The aver- age tensile strength of these groups was statistically indistinguishable from that of the

groups of Al-300 and Al-995 aluminas which

were given treatments 1 and 2. Because of the

limited number of specimens available, treatment 3 was not extended to the other materials.

d. The Effect of Treatment 4 on the Strength of

Alumina

As-ground specimens were also immersed

in distilled water for 1 week. Before testing,

the specimens were again thoroughly dried as described previously. The Al-300 alumina

(group A-5) had a slightly lower strength than

groups A-3 and A-4, but the difference is negligible and it is not statistically justifiable to attribute it to treatment 4. The same is true for

Al-995, where there is no meaningful difference

between group B-5 and groups B-3 and B-4.

In the case of AD-995, the difference is again

insignificant. The strength of group C-5 is

somewhat higher than the strength of other

groups in the AD-995 series (C-2, C-3, and

C-4) that were given a vacuum bakeout. This

observation appears to corroborate the finding

that the bakeout under conditions used in the

work described is detrimental to the strength

of this material.

e. The Effect of Treatment 5 on the Strength of

Alumina

The presence of condensed water on the sur- face of specimens during testing caused a marked lowering of strength of the Al-300 and

Al-995 aluminas (groups A-6 and B-6, respectively). The strength of AD-995 alumina

(group C-6) did not seem to be affected. The

reason for this is not clear, but it is possible

that, because of poor control, the temperature

of these specimens might have risen above the

dew point.

f. The Effect of Treatment 6 on the Strength of

Alumina

The strength data for groups A-7, B-7, and

C-7 show conclusively that a film of water on

the surface of the specimens drastically lowers

the strength. An interesting fact is that the

extent of strength deterioration grows larger

with increasing grain size. Comparison of the

strength of groups C-1 and C-7, B-1 and B-7,

and A-1 and A-7 shows that the decrease in

strength is 19.3 percent in AD-995, 21.7 percent in Al-995, and 24.5 percent in Al-300.

The mechanism of the attack by liquid water

can possibly be found in the nature of the test

method used. It can be visualized that the rubber bulb transmitting the hydrostatic pressure

to the specimen forces the liquid into cracks

and pores on the inner wall where the water

can act as a hydraulic wedge.

3.4. Vacuum Strength of Alumina

To complete the picture of the effect on

strength of alumina of various post-grinding

treatments—all of which involve moisture—it seems appropriate to consider the opposite

situation, i.e., the strength behavior of alumina

under vacuum.

It has long been suspected that atmospheric

moisture is detrimental to the strength of ce- ramics, but this phenomenon was not properly

recognized until Roberts and Watt observed that

alumina specimens failed in static tension after

varying lengths of time under load [7] . Pearson

observed that delayed fracture can be practically eliminated when the test specimens are given a vacuum bakeout and are tested under

vacuum [8].

Sometimes the deleterious effect of the moisture on the strength can assume more subtle

aspects and may not be recognized as such. A

good example is the effect of stress rate on

strength, which has been studied previously at

SRI [9]. The results of this study are shown

as the S-shaped curve in figure 3 which is a

plot of strength of alumina versus stress rate.

When the original curve was constructed, the

reasons for its existence and shape were ex- plained on a simple mechanistic basis. It was

postulated that at high stress rates the propa94
STRESS RATE — psi/sec

10^ 10^ 10" 10^

STRESS RATE — N/m^/sec

•Data obtained undef Contract No. N00019-69-C-0229 TA-4372-2R3

Figure 3. The effect of stress rate and

test conditions on the tensile strength

of Al-995 alumina.

gating crack is forced to seek the shortest route

even if it entails going across grains which re- quire a higher driving force. On the other

hand, at slow stress rates, the crack is allowed

to find the path of least resistance, which leads

to lower strength values. However, this ex- planation was never considered very satisfactory and other interpretations were sought.

Static fatigue studies of other workers [7, 8]

provided the clue, and stress corrosion was considered. Accordingly, the following experiments

were performed. Al-995 alumina specimens

ground in the conventional way were tested at

two stress rates, 2.1 x 10^ N/m^ s and 2.1 x 10"

N/m^ s (3 X 10 psi and 3 x 10^ psi/s) first in

air, and then under vacuum. The results are shown superimposed on the curve in figure 3 and

the comparison of air and vacuum strengths is

presented in table 5. The strength values obtained in air (16.8 x 10^ and 21.2 x 10^ N/m^

-24.3 X 10^ and 30.8 x 10^ psi) agree closely

with the old data generated on a different lot of

Al-995 alumina [9] . The vacuum strengths are much higher and very similar, i.e., 23.4 and

24.7 X 10^ N/m^ (33.9 x 10^ and 35.8 x 10^

psi) and strongly indicate that under vacuum

the effect of stress rate on strength is practically nil. The difference (5.3%) between vac- uum strengths is small when compared to the

difference between air strength values obtained

at the two stress rates (21.1%). Even the 5.3

percent difference can be plausibly explained as a finite permeability of rubber separating the

pressurized water from the specimen.

In the light of these results the relationship

between stress rate and strength becomes quite

understandable. It is safe now to presume that

there is a stress-activated chemical reaction

(stress corrosion) between alumina and water

which takes place at the crack tip. The kinetics

of this reaction are controlled by the diffusion

rate. Therefore, at slow stress rates, there is a

sufficient supply of water at the crack tip,

stress corrosion takes place, and low strengths

result. As the stress rate increases, the rate of

water transport begins to fall behind the rate

of force application. Stress corrosion at this

point becomes diffusion or transport limited,

and a plateau in strength versus stress rate is

observed. Finally, at higher stress rates, stress

corrosion becomes less and less active and the

strengths increase.

There is experimental evidence that the water

responsible for the lowering of strength of ce- ramics is physically adsorbed on the surface

and can be easily driven off. It was not possible

to obtain the same vacuum each time and ex- periments were carried out in the pressure

range of 6.7 x 10"^ N/m- to 2.7 x 10-« N/m^'

(5 X lO-'^ to 2 X 10-' torr). The variations in

pressure produced no noticeable change in

strength. Also the difference in lengths of time

for which the specimens were kept under vac- uum could not be correlated with strength.

4. Conclusions

Microstructure, typified by grain size, is the

most important factor controlling the strength

of ceramics. Under reasonably careful grinding

conditions, surface flaws of the Griffith type

penetrate only one grain depth into the body.

The relationship of grain size to strength can

be demonstrated by the application of Griffith's

failure criterion for brittle materials. The type

of surface finish obtained in grinding aluminas

is not solely the result of the grinding technique,

but is strongly dependent on the grain size of

the ceramic. Exposure to water (liquid or va- por) does not impair the strength of alumina

since hydrates were not formed. The strong deTable 5. Average tensile strength of Al-995 alumina in air and under vacuum

Conditions

Stress rate Strength

Coeff. of var., N/m^'s psi/s N/m= X 10' psi X 10' percent

Air 2.1 X 10' 3.0 X 10 16.8 24.3 2.3 Vacuum 2.1 X 10* 3.0 X 10 23.4 83.9 1.5

Air 2.1 X 10« 3.0 X 10' 21.2 30.8 3.2 Vacuum 2.1 X 10« 3.0 X 10' 24.7 35.8 4.4

95
pendence of the strength of alumina on stress

rate is shown to be caused by stress corrosion.

The strength of alumina is higher under vac- uum than in air and is independent of stress

rate.

The work described was sponsored by the

Naval Air Systems Command under Contract

Nos. NOOO19-68-C-0388 and N00019-69-C0229.

5. References

[1] Sedlacek, R., Halden, F. A., and Jorgensen, P. J., The effect of grinding variables on the strength

and surface finish of alumina, these proceedings.

[2] Sedlacek, R., and Halden, F. A., Method of tensile

testing of brittle materials. Rev. Sci. Instr. 33,

298-300 (March 1962).

[3] Frish, B., Die Hydratation von a Aluminiumoxid,

Ber. Deutsch. Keram. Ges. 42, 149-160 (1965).

[4] Charles, R. J., and Shaw, R. R., Delayed failure of

polycrystalline and single-crystal alumina. General Electric Research Laboratory Report No.

62-RL-3081M, (July 1962).

[5] Wiederhorn, S. M., Fracture of sapphire, J. Am.

Ceram. Soc. 52 [9], 485-491 (1969).

[6] Hillert, M., On the theory of normal and abnormal grain growth, Acta Metallurgica 13,

227-238 (1965).

[7] Roberts, J. P., and Watt, W., Mechanical properties of sintered alumina, Ceram. glass 10. 53 (1952).

[8] Pearson, S., Delayed fracture of sintered alumina,

Proc. Phys. Soc. (London) 69 (ffB), 1293-1296

(1956).

[9] Sedlacek, R., Tensile Strength of Brittle Materials, Technical Report AFML-TR-65-129,

(August 1965).

Discussion

Firestone: I notice that you did some ex- periments where your specimens were immersed in water. You said that there was no

significant decrease in strength, and yet when

the water condensed on the surface there is. Sedlacek: You have misunderstood. When

the specimens are immersed in water and tested

wet there is a quite pronounced decrease in

strength. The decrease follows the grain size;

the greater the grain size, the greater the de- crease in strength. However, if you immerse

the specimens in water after bakeout for instance, and then take them out, wipe them, and

leave them at room temperature to eauilibrate

for about one hour, their strength will be exactly the same as if they had not had any

treatment. The same goes for variable humidity as I mentioned.

Cameron: Would the presence of a film of

moisture between the rubber membrane and

your specimen act as lubricating layer leading

to a more efficient load transfer?

Sedlacek: I don't think there is a need for

any lubrication because the rubber expands

linearly with the expansion of the specimen.

There is no friction between the two.

Wiederhorn: About three years ago, we did

some experiments on crack propagation in single crystal sapphire using the double cantilever

technique.^ Crack propagation in this case was

on the (1012) twin plane. We duplicate your

results if crack velocity versus the applied load

is plotted.

Sedlacek: I have seen your data. You plotted

the velocity of crack propagation as a function

of relative humidity and you ended up with a

series of S-shaped curves. Wiederhorn: We had S-shaped curves and

we believe that we had plateaus in the curves.

' S M. Wiederhorn, Int. J. Fract. Mech., 4, 171 (1968)

Sedlacek: Do you have an explanation for

the plateau ? Wiederhorn: Yes. When cracks were moving fast enough the rate of crack proDagation

was controlled by the rate at which the water

diffused to the crack tip. Water diffusion in

gas is not load sensitive, therefore, the crack

velocity was not sensitive to applied load in the

plateau region.

Westwood: What was the mechanism?

Sedlacek: I would say that the reduction

of surface energy holds for the case of the

specimens being covered with a liquid film.

The concept of stress corrosion is more important in this case. Westwood: May I suggest another possibility involving the Rebinder effect? Consider a crack with water in it. The water "softens" the

reeion ahead of the crack tip, allowing dislocations to be generated and moved more easily.

Subcritical crack growth then occurs by the

mechanisms discussed by Bob Stokes, Freddy

Clarke, and others. In other words, slow crack

growth occurs because of the moisture—but not

because of corrosion or surface energy reduction phenomean. You should see such effects at

slow rates of loading.

Heuer: I think there are other explanations,

not involving dislocation phenomena, which allow a slow rate of crack growth in the presence

of a film of moisture.

Westwood: Corrosion is one. Sedlacek: There is one thing that I would

like to add if I can have the last slide back

again. The curve, as I said, undergoes inflection.

I must emphasize that this curve is typical only

for Al 995 alumina. It may not assume exactly

the same shape for other aluminas. Presently,

we are working with PZT materials and they.

96
for instance, don't show any effect of different

stress rates from 5000 to 150,000 psi per second.

Westwood: That is an important point, be- cause one of the things which has come out of

our studies of the Rebinder effect is the fact

that in ionic sohds it is very sensitive to impurities. In fact, one can get adsorption-induced

hardening or softening effects depending on the

concentration and state of ionization of the impurities. I would imagine that if you could find

the right impurity concentration (and chemistry) to give you an environment-sensitive hardening effect, you would not see delayed failure.

On the other hand, corrosion or surface energy

dependent processes should not be so sensitive

to impurity concentration.

Heuer: No, that's not true. Impurities can

also affect the surface energy. The point I was

making is that you can get slow crack growth

by an adsorption mechanism. The difficulty with

a conventional Grifl^th-type approach, involving

the reduction of surface energy, is that the effect is too small. On the other hand, it may be

that a slow crack growth can take place at subcritical Griffith stresses if adsorption occurs.

Rice: Following along with what Bert Westwood said, you essentially explained the lower

value of strength at the lower stress rate as

possibly being due to an environmental problem, but it might also be an indication of, for

example, a dislocation or a twin type effect

since you can expect a lower strength for lower

stress rates.

Sedlacek: Well, as soon as I removed the

moisture, I don't get a decrease.

Rice: Your chart showed somewhat lower

strengths at lower stress rates.

Sedlacek: Because the membrane is not completely impermeable.

Rice: Right, but that's one hypothesis, but

I'm saying another. Dislocations or twins are important, and these would also cause lower

strengths.

Rhodes: We have noticed similar effects in

hot pressed alumina. I think for the record I would like to ask whether or not your alumina

is free of glassy phase.

Sedlacek: No, it is not. It is nominally 99.5

percent alumina and" it has about 0.3 percent

silica. So it has a glassy phase.

national bureau of standards special pub. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

97

Grinding Alumina with Diamond Abrasives

R. J. Caveney and N. W. Thiel

De Beers Diamond Research LaboratoryJohannesburg, South Africa

This paper is the result of a preliminary investigation into the grinding of

high density alumina with diamond abrasive grinding wheels.

Following extensive work on the application of diamond for tungsten carbide

and steel grinding, the De Beers Diamond Research Laboratory recently embarked

on a research programme to study the various aspects of ceramic grinding. In the

Diamond Abrasive Technology Centre's Grinding Section, ceramics with three different alumina contents were ground with a variety of diamond types both natural and

synthetic to establish what effect changes in wheel speed have when surface grinding

these materials. The mode and extent of damage of the ground surfaces of the

specimens were investigated using conventional electron microscopy. In addition to these investigations, measurements were made of the quality of the surface finishes

that can be obtained under different grinding conditions. Variations in grinding effi- ciency, machining costs and surface damage are discussed for the three alumina

grades tested. The mechanism of wear of the different diamond types was studied

and is demonstrated by means of scanning electron micrographs.

Also reported in this paper, are the initial results obtained from tests on cylinders of the same three alumina grades mentioned above. These initial tests were

aimed at finding the most suitable combination, from the point of view of wheel life,

of wheel and workpiece speeds, and incorporated tests with both natural and synthetic

grits.

Key words: Alumina; cylindrical grinding; diamond grinding wheel; electron microscopy; removal mechanism; surface damage; surface finish; surface grinding; wear

mechanism.

1. Introduction

Little has been published on the science and

technology of grinding ceramics with diamond

abrasive wheels. Many articles have, however,

demonstrated the feasibility, utility and econ- omy of diamond wheels in this application. It

is further known that diamond wheels are used

extensively in the ceramics industry. Due to

the lack of published data, it is not known

whether this use has been optimised using

scientific and economic principles. For these

reasons, the De Beers Diamond Research Laboratory recently embarked on a research programme to establish the techniques and fundamentals of the diamond grinding of ceramics.

The present paper and an associated paper [1]^

report the preliminary results of this study.

The underlying principle on which this work

has been based is that the maximum removal

rate consistent with surface damage that is not

excessive should be employed. In order to evaluate surface damage, it was necessary to undertake an electron microscope study of the workpieces after grinding. This study also enabled

an assessment of the mode of material removal

to be made. Besides stock removal and surface

damage, surface finish or roughness must also

be determined and the latter minimised.

As was noted earlier, this work is still in its

* Figures in brackets indicate the literature references at the end (if this paper.

initial stages. Only three grades of alumina

have been studied, namely 85 percent, 94 percent and 96 percent alumina. The material

ground was in the form of large blocks. At a

later stage it is anticipated that ceramics with

alumina contents greater than 99 percent will

be studied, and the effect of the grinding parameters on the chipping of small samples will

be evaluated.

1.1. General Test Conditions

It is well known that the condition of the

surface of a diamond abrasive grinding wheel

can have a considerable effect on the results

obtained by the wheel. Because of this fact all the wheels used in this test programme were

given a conditioning run under the test conditions before the actual tests were started.

A coolant flow rate of 3 litres per min (40

gallons per hour), measured by rotameters, was

used throughout the tests. Coolant tanks were

replenished with premixed coolant at the start

of each test and completely renewed after five

tests. Each machine is equipped with a centrifuge to clarify the used coolant and the bowl

of the centrifuge, containing the grinding

sludge, was cleaned at the end of each test. Since the tests were initially aimed at achieving a high stock removal rate, no attempt was

made to obtain good surface finishes. At the

end of each of these stock removal rate tests.

99
Figure 1. Measurement of specimen block for amount

of material removed.

a spark-out, without downfeed, of two full

passes across the specimen was given. A further investigation was carried out at the conclusion of the stock removal programme to

determine the surface finish that could be obtained. Measurements of the surface finishes

obtained were carried out with Hommel-Tester

Type T equipment, the figures quoted being AA

or middle roughness values.

The amount of material removed from a

specimen during a test was measured, in the

case of surface grinding, with a dial depth

indicator (see fig. 1) at each corner of the

ground block before and after the test, and

with micrometers in the case of the cylindrical

specimens. In surface grinding the average

amount of material removed in one test was 345

cm^ (21 in^), while, for cylindrical grinding, the

average removal was 800 cm^ (49 in^).

After each conditioning run or test, the

wheels were placed in a temperature controlled

measuring room (held at 20° C ±l^ deg) for

at least 24 hours before measurements were

made. This is the standard procedure used in

the Diamond Abrasive Technology Centre's

Grinding Section, since it allows the wheel to

overcome the effects of centrifugal and thermal

expansion.

'"'^he amount worn from the wheels during

a test was measured on special jigs (see fig. 2

and 3).

The wheel efficiency, G ratio, is the ratio of

volumes worn from the specimen and the wheel

during a test. ^ _ Volume material removed

Volume wheel worn away

Calculations of cost have been made at a rate

of $12.00 per hour for machine, operator and

overhead costs, and $25 per cubic centimeter

Figure 2. Special jig for measuring wear of wheels for

surface grinding tests.

Figure 3. Special jig for measuring wear of

wheels for cylindrical grinding tests.

100
of grit-bearing bond for 100 concentration

wheels.

2. Surface Grinding

2.1. Introduction

This preliminary investigation of the effect

of alumina content and wheel speed on grindability, with different diamond and bond types,

was conducted on horizontal spindle, reciprocating table surface grinders. This type of

grinder had to be used because the Diamond

Abrasive Technology Centre (D.A.T.C.) does

not, at present, have machines of the vertical

spindle, rotary table type, on which the greater

proportion of surface grinding of ceramics is

done. Further tests on the latter type of machine are planned for the future.

Three different alumina types were ground

with ten different wheels. Of the ten, six were

metal bond wheels containing three grit types

in each of two sizes, while the remaining four

wheels contained two different resin bond grits

also in two sizes. Initially the machines were set

at maximum crossfeed and table speed (1.75

mm (0.070 in) and 20 m/min (65 ft/min) re- spectively), and the downfeed was increased

from 0.025mm (0.001 in) to 0.0625 mm (0.0025

in) per pass. When an attempt was made to use a downfeed of 0.125 mm (0.005 in), the

machines stalled and both crossfeed and table

speed had to be reduced.

It was found that by reducing the crossfeed

and table speed to the D.A.T.C.'s standard val- ues used for tungsten carbide grinding, viz 1.5

mm (0.060 in) and 16 m/min (53 ft/min) re- spectively, and the 0.125 mm (0.005 in) downfeed, the maximum stock removal rate was

achieved. Because of the relatively small sur- face damaged caused by this heavy downfeed,

all subsequent tests, with the exception of sur- face finish tests, were conducted under these

conditions. The wheels were then run at three

different peripheral speeds on each ceramic

type to ascertain the effect on the grinding effi- ciency.

The ceramic workpieces, before and after

etching, were examined in the electron microscope, and the depth of damage assessed. The

modes of diamond wear and wheel wear were

studied using optical and scanning electron

microscopy.

2.2. Test Conditions

Machine

Wheel type

Two Jones and Shipman

540 Tool Room Surface

Grinders with 1.5 kw

(2 hp) spindle drive

motors.

DlAl

127 X 6.3 X 31.8 mm

(5 X 1/4, X VA in)

3000, 4000, 5000 r.p.m.

19.9, 26.6, 33.3 m/sec

(3930, 5240, 6650 ft/

min)

De Beers RDAR55N,

DXDA-MC (resin

bond), DXDA, MDA-S,

SND-MB (metal bond)

80/100 and 140/170 U.S.

mesh.

100

3.1 mm ( Vs in)

0.125 mm per pass

(0.005 in)

22.5 mm (approx % in)

1.5 mm per table reversal

(0.060 in)

16 m/min (53 ft/min)

Water plus Bryto 5

(100:1)

3 1/min (40 gallons per

hour)

Coors AD 85, 94, 96

152.4 X 101.6 mm—155

cms^ (6x4 in—24 in^)

Two spark out passes,

without downfeed at

the end of each test. The initial surface damage tests were carried out with a table traverse speed of 20 m/min

(65 ft/min), a crossfeed of 1.75 mm per table

reversal (0.070 in) and 4000 rpm wheel speed.

2.3. Test Results

The results of the wheel efficiency tests un- der the above conditions are given in figures

4, 5 and 6. The results of examinations of the

wheels and workpieces are discussed below.

2.4. Discussion

In the ensuing discussion, frequent mention

is made of the various grit types.

RDA55N is a metal clad synthetic diamond

abrasive designed for use in resin bond wheels.

The underlying particle is characterised by its

friability and irregular shape. The metal cladding, 55 per cent by weight, results in better

retention of the diamond particles in the wheel

bond.

DXDA consists of strong, fracture resistent

diamond particles which are blocky with a controlled proportion of cubo-octahedra. This grit

is used in metal bond wheels. The same grit,

after metal cladding, is known as DXDA-MC

and is used in resin bond wheels.

MDA-S consists of very strong synthetic

Wheel size

Wheel speeds

Wheel peripheral

speeds

Diamond types

Diamond mesh

sizes

Diamond con- centration

Bond depth

Downfeed

Total downfeed

per test

Crossfeed

Table traverse

speed

Coolant

Coolant flow rate

Specimen types

Specimen face

size

Sparkout

439-921 O - 72 - 101
Figure 4. Changes in grinding ratio with

wheel speed for 85 percent alumina.

Resin bond wheels: 1. DXDA-MC

80/100, 2. DXDA-MC 140/170, 3. RDA55N 80/100, 4. RDA55N 140/170.

Metal bond wheels: 5. MDA-S 80/100,

6. MDA-S 140/170, 7. SND-MB

80/100, 8. SND-MB 140/170, 9. DXDA

80/100, 10. DXDA 140/170.

diamond crystals, mainly cubo-octahedra, for use in metal bond wheels.

SND-MB is a natural diamond abrasive designed for use in metal bond wheels. This grit

consists of strong, blocky particles which are prepared from larger diamonds using a milling

process.

a. Examination of the Workpieces

Two sets of observations were made. The unetched workpieces of all three grades of alumina were studied after grinding tests at 4000

rpm and infeeds of 0.025, 0.0625 and 0.125 mm

(0.001, 0.0025 and 0.005 in). Replicas, prepared using a water soluble plastic [2, 3], of

the ground surface and the face perpendicular

to the ground surface (and parallel to the

grinding direction) were studied in a Philips

EM200 electron microscope. The second set of

observations was made on etched (one minute

in boiling NaOH) ground surfaces of the three

ceramics after grinding tests at 4000 rpm and

0.125 mm (0.005 in) downfeed.

(1) Unetched Specimens. Damage of the

Figure 5. Changes in grinding ratio

with speed for 9^ percent alumina.

(Curves marked as in fig. 4).

grains in the ground surface was noted in almost all cases. The surface structure consisted

of fracture surfaces of grains, and freshly ex- posed grain surfaces, some of which exhibited

traces of cracks. The degree of damage varied

depending on the wheel with which the material was ground and the grinding parameters.

In the case of the 85 percent alumina, some

evidence of extrusion of the binder material

was noted. This is illustrated in figure 7 for

the 85 percent alumina ceramic ground with a

resin bond wheel containing the 140/170

RDA55N grit.

Figure 7 also shows the multiple grain fracture characteristic of the surfaces ground with

resin bond wheels containing 140/170 RDA55N.

This is further illustrated for the 96 percent

alumina in figure 8. The same behaviour, but to a lesser degree, was noted for the resin bond

wheels containing 80/100 RDA55N and 140/

170 DXDA-MC, and the metal bond wheel

containing 140/170 SND-MB. These observations are similar to those made on tungsten car- bide-cobalt composites after grinding with resin bond wheels [4]. In this case, cobalt is

extruded to the free surfaces and the carbide

grains are extensively fractured.

The ground surface of the ceramics machined

with metal bond wheels containing MDA-S and

102
0

1

1 ' 1 20 25 30 35

I WHEEL SF^EED m/sec. , 3000 4000 5000 WHEEL SPEED rpm.

Figure 6. Changes in grinding ratio

with wheel speed for 96 percent

alumina. (Curves marked as in fig. 4).

DXDA grits exhibited fairly extensive fracture

of the grains (see figure 9). The degree of fracture and extrusion was less than for the ce- ramics ground with resin bond wheels as described earlier. This behaviour is different from

that of tungsten carbide-cobalt composites when

grinding with metal bond wheels [4], where

extensive cobalt extrusion and minimal fracture

of the grains were noted. The differences are most likely due to the different flow characteristics of the binder materials and the different strengths of the tungsten carbide and

aluminum oxide grains.

Less grain damage was observed for the

ceramics ground with the resin bond wheel con- taining 80/100 DXDA-MC and the metal bond

wheel containing 80/100 SND-MB. This is il- lustrated in figure 10. Another interesting observation is that there appears, when grinding

with these wheels, to be a tendency for the

damage to decrease on increasing the severity

of the grinding above a critical level.

On studying the face perpendicular to the

ground surface, it was noted that the damage

was, for the most part, confined to the surface

layer of grains. The depth of damage did not

increase appreciably on increasing the infeed

from 0.0625 mm (0.0025 in) to 0.125 mm

(0.005 in) ; in the latter few grains in

the second layer below the surface exhibited

Figure 7. Electron micrograph of a replica

of the surface of 85 percent alumina after

grinding with a resin bond wheel containing 80/100 RDA55N. Extruded binder material (G) is observed.

\

Figure 8. Electron micrograph of a replica of the surface of 96 percent alumina

after grinding with a resin bond wheel

containing HO/1 70 RDA55N.

fracture (see figure 11). Elimination of "sparkout" does not appear to affect the extent or depth of damage.

(2) Etched Specimens. It was difficult to assess the damage because of the fact that both

the fracture and grain surfaces have a similar

103
Figure 9. Electron micrograph of a replica

of the surface of 96 percent alumina after

grinding with a metal bond wheel con- taining 80/100 MDA-S.

Figure 10. Electron micrograph of a replica of the surface of 96 percent alumina

after grinding with a resin bond wheel

containing 80/100 DXDA-MC.

appearance after etching. However, the grain

boundaries are more readily observed, and the

fracture cracks readily observed as shov^n in

figure 12. The assessment of damage gave the

same results as described for the unetched

specimens.

b. Study of the Grinding Wheels

The wheels, after grinding the 96 percent

alumina at 4000 rpm and 0.125 mm (0.005 in)

infeed, were studied by means of optical scan- ning electron microscopy. A plastic replica

technique [3] was used to study the macroscopic specimens in the scanning electron microscope.

The degree of retention of the grit in the

bond, and the mode of diamond wear, are summarized in table 1. For the resin bond wheels, the RDA55N

grits are better held in the bond than the

DXDA-MC grits. This is particularly noticeable in the 80/100 size range. This effect is

caused by the increased roughness of the RDA

particles, and hence the coated grit surface,

compared to the DXDA particle. The 140/170

grits are better retained in the bond than

Figure 11. Electron micrographs of replicas showing the ground surfaces (G), the edge (E) and the face perpendicular to the ground surface (P) of (a.) 85 percent alumina, (6) 9^. percent alumina and (c) 96 percent alumina

after grinding with a metal bond wheel containing 80/100 MDA-S (0.125 mm (.005 in) downfeed).

104
Figure 12. Electron micrograph of a replica of an etched 96 percent alumina sur- face after grinding with a metal bond

wheel containing H.0/170 SND-MB.

the 80/100 grits. The percentage of grits re- tained in the metal bond wheel is of the order

of 80 percent, regardless of size. The SND-MB

grit retention is superior to that of the MDAS and DXDA grits due to the serrated nature

of the surface of the SND-MB grit particle.

The DXDA grit appears to be held better than

MDA-S grit in the 80/100 size, and vice-versa

for the 140/170 size.

From table 1 it is apparent that the mode of

diamond v^ear is different for the different types

Figure 13. Scanning electron micrograph

of a replica of the resin bond wheel con- taining 80/100 RDA55N.

of grit. As observed after grinding tungsten

carbide-cobalt composites, the RDA grit particles are extensively fractured (figure 13).

The RDA grit is engineered to be of a friable

nature and suffers repetitive multiple fracture

so that sharp fracture points are continuously

presented to the v^^orkpiece. The 140/170

DXDA-MC grit in a resin bond v^heel has a

similar fractured appearance to the RDA grit,

except that a significant amount of abrasion

wear is present.

Table 1. Summary of wheel observations

Percentage Percentage

Grit Size Bond Percentage Percentage of exhibiting exhibiting

retention grit exhibiting predominantly predominantly

in bond no wear abrasion wear fracture wear DXDA-MC 80/100 R

DXDA-MC 140/170 R

RDA55N 80/100 R

RDA55N 140/170 R

MDA-S 80/100 M

MDA-S 140/170 M

SND-MB 80/100 M

SND-MB 140/170 M

DXDA 80/100 M

DXDA 140/170 M

50 12 55 33

72 13 32 55

73 Predominantly fracture, small abrasion

marks on approximately 60 percent of grit

particles.

81 Very little abrasion observed.

78 2 55 43

84 12 55 33

86 8 36 56

86 9 36 55

82 11 57 32

75 4 67 29

105
Figure 14. Scanning electron micrograph

of a replica of the metal bond wheel con- taining 80/100 DXDA.

Abrasion wear flats are dominant in the

MDA-S and DXDA grits in the metal bond

wheels (fig. 14). Most of the SND-MB grits

in the metal bond wheels have a predominantly

fractured appearance; some abrasion wear is,

however, present (fig. 15). The behaviour of

the SND-MB grits is caused by the fracture

centres introduced in the milling process, as opposed to the as-grown strong MDA-S and

slightly weaker DXDA grits.

Of the grit particles present in the surface

of the 80/100 DXDA-MC wheel, 55 percent

exhibit predominantly abrasion wear flats.

However, only 50 percent of the grit particles are retained in the bond, and hence the role

of abrasion wear is not as dominant as is the case with the MDA-S and DXDA grits in the

metal bond wheels (for example: 80/100 MDAS, 55 percent of diamond grit particles exhibit

predominantly abrasion wear flats, but 78 percent of the grit particles are retained in the

wheel). In fact the balance between abrasion

and fracture of the grit in the 80/100 DXDAMC resin bond wheel is similar to that of the

80/100 SND-MB metal bond wheel. Grit pullout, however, starts at an earlier stage in the

working history of a grit particle for the

DXDA-MC grit than for the SND-MB grit.

The balance between the abrasion and fracture

modes of diamond wear is dependant on the

effective strength of the grit particles and the

bond in which the grit is incorporated.

c. Grinding Efficiencies

(1) 85 percent Alumina. As is shown in

Figure 15. Scanning electron micrograph

of a replica of the metal bond wheel con- taining lhO/170 SND-MB.

figure 4, with the exception of two of the wheels,

the highest efficiencies on this ceramic are achieved at a wheel speed of 26.6 m/s (5240

ft/min). The exceptions are the metal bond

wheels with 140/170 SND-MB grit, which

shows an increase in efficiency as the wheel

speed increases, and with 80/100 DXDA, which

shows a decrease in efficiency with an increase

in wheel speed. At this stage, there can be no

definite explanation of why these two grits

should behave in this way, but it is thought

that a change in mode of grit wear could be

the cause of the diff'erent behaviour of the

SND-MB wheel.

Clearly, the resin bond wheels containing

RDA55N and DXDA-MC do not perform as

well as the metal bond wheels. It is interesting

to note that the weaker resin bond grits

(RDA55N) give higher efficiencies at all speeds

and in both grit sizes.

With the exception of the anomalous behaviour of the metal bond wheel containing

80/100 DXDA (highest G ratio at 3,000 rpm),

the wheel containing 80/100 MDA-S appears

to give the best grinding efficiencies.

(2) 94 percent Alumina. From figure 5

it is immediately apparent that this material

grinds in an entirely different way from the

85 percent alumina. With the 94 percent alumina, there are only two wheels that achieve the

highest G ratio at 26.6 m/s, these being the

resin bond 80/100 DXDA-MC and RDA55N

wheels. The same grit types in the 140/170 size

show an increase in efficiency as the wheel

speed increases. On the other hand, the SND106
MB wheels decrease in efficiency as speed increases, the G ratio at 33.3 m/s (6550 ft/min)

being only 50 percent of that obtained at 19.9

m/s (3930 ft/min) for the 80/100 grit size

and 46 percent for the 140/170 grit size.

The other metal bond wheels give their lowest

G ratios at 26.6 m/s (5240 ft/min). Before

any conclusions can be drawn from these

trends it is obvious that further tests are neces- sary at higher and lower speeds.

One other fact that is surprising is that the

G ratios when grinding 94 percent alumina

are generally higher than those obtained on 85

or 96 percent.

(3) 96 Alumina. It can be seen from figure

6 that 96 percent alumina grinds in a similar

manner to 85 percent. The resin bond wheels

show a general peak in efficiency at about 27

m/s (5300 ft/min). The G ratios at this speed

can be correlated with the appearance of the

wheel surface i.e., grit pull-out and grit wear

mechanisms. The RDA55N wheels have higher

efficiencies and higher grit retention than the

DXDA-MC wheels. This is shown in table 2.

The SND-MB grit has the best retention in

the metal bond wheels. However, grit wear, in

this case, is mainly by fracture, and hence more

rapid than when abrasion wear is dominant.

For this reason, the wheels containing SNDMB have lower grinding efficiencies than those

containing MDA-S and DXDA grits when

grinding at 26.6 m/s (5240 ft/min). The MDAS and DXDA grits have the same basic wear

pattern, and hence grinding efficiencies are, to a large degree, dependant on grit retention.

This is illustrated in table 3 and figure 6. The patterns of behaviour of the DXDA and

140/170 MDA-S wheels on this material are again similar to those on the 85 percent alumina. Here, however, the 80/100 MDA-S wheels

also decrease in efficiency as the speed increases,

and the SND-MB wheels behave in a similar

way. In fact, the trends of the latter wheels

are very similar on 94 and 96 percent alumina.

It should be noted, however, that the diamond grit wear mechanism may change with

grinding conditions. This is especially true for

the 80/100 DXDA-MC grit in a resin bond

wheel and the 80/100 SND-MB grit in a metal

25000 -

85 94 95 % ALUMINA

Figure 16. Changes in grinding ratio

with alumina content for 3000 rpm.

Resin bond wheels: 1. DXDA-MC

80/100, 2. DXDA-MC 140/170, 3. RDA55N 80/100, 4. RDA55N 140/170.

Metal bond wheels: 5. MDA-S 80/100,

6. MDA-S 140/170, 7. SND-MB

80/100, 8. SND-MB 140/170, 9. DXDA

80/100, 10. DXDA 140/170.

bond wheel. Under light grinding conditions,

abrasion wear will play a dominant role, but as

the conditions become more severe, the balance

will swing to fracture wear with a corresponding large decrease in grinding efficiencies.

It is felt that these concepts apply to all the

workpieces and grinding conditions.

(4) Effect of Alumina Percentage. The

ways in which alumina percentage affects

grinding efficiencies for the three wheel speeds

used are shown in figures 16, 17 and 18.

At 19.9 m/s (3930 ft/min) all wheels, ex- cept the resin bond wheels containing RDA55N

grits, achieve their highest efficiencies on 94

percent alumina, with 85 percent being the

most difficult to grind in all cases. This grade

is also the hardest to grind with the RDA55N

Table 2. Comparison of grinding ratio and grit retention for the resin bond wheels, (96% alumina, 0.125

mm (0.005 in) downfeed and UOOO rpm)

Percentage

Grit Size grit Grinding

retention ratio

Table 3. Comparison of the grinding ratio and grit

retention for metal bond wheels containing MDA-S

and DXDA (96% alumina, 0.125 mm (0.005 in.)

downfeed and AOOO rpm)

Grit Size

Percentage

grit retention

Grinding

ratio

RDA55N 80/100 73 4660 MDA-S 80/100 78 15,260

DXDA-MC 80/100 50 2720 DXDA 80/100 82 15,920

RDA55N 140/170 81 1990 MDA-S 140/170 84 10,700

DXDA-MC 140/170 72 1720 DXDA 140/170 75 10,190

107
20000

iOOO r.p.m. 20000

5000 rp.m.

2 0

I 1 1 L. 85 94 96 % ALUMINA

Figure 17. Changes in grinding ratio

with alumina content for 4000 rpm

(Curves marked is in fig. 16).

wheels. These latter wheels, however, achieve

their highest G ratios on 96 percent alumina.

At 26.6 m/s (5240 ft/min) the pattern

changes and this can be related to the appearance of the ground surfaces. With the exception of the surfaces ground with the resin bond

wheel containing 80/100 DXDA-MC and the

metal bond wheel containing 80/100 SND-MB

the damage on the ground surface generally in- creases with increasing alumina content. The

grinding efficiencies also increase. For the 80/

100 DXDA-MC resin bond wheel and the 80/

100 SND-MB wheel, the amount of damage

after grinding the 96 percent alumina was less

than that on the 85 percent and 94 percent

aluminas. The grinding efficiencies also corre- spondingly decreased. It appears, therefore,

that the efficiency and surface damage are related. This is interpreted as a change of diamond grit wear from a dominant abrasion

wear mechanism to a dominant fracture mechanism on increasing the alumina content of the

workpiece.

With a wheel speed at 33.3 m/s (6550 ft/

min), only the RDA55N 80/100 and 140/170

SND-MB wheels are exceptions to the same

trend as found at 19.9 m/s (3930 ft/min) i.e.,

highest G ratios on 94 percent and lowest on

85 percent alumina. The 140/170 SND-MB

wheel exhibits an unusual decrease in efficiency

with increase in alumina content.

d. Material Removal Mechanism

The alumina is removed by two basic processes, fracture of the grains and subsequent

removal of the fracture pieces, and removal of

15000 -

< 10000 -

^ 5000 -

0

I 1 1 1 85 94 96 % ALUMINA

Figure 18. Changes in grinding ratio

with alumina content for 5000 rpm

(Curves marked as in fig. 16).

the individual grains. This is analogous to the

removal of material when grinding tungsten

carbide-cobalt composites [4]. The balance between these two modes depends upon the nature

of the grit surface presented to the workpiece.

When many fracture points are presented to

the workpiece, as is the case when grinding

with the RDA type grits, then the principal

mode of material removal is by grain fracture

and removal of the fractured pieces. When the

grit has many abrasion wear flats (MDA-S and

DXDA in metal bond wheels), then a certain

amount of crushing of the grains occurs, and

some grains are removed individually. The

fracture of the 80/100 DXDA-MC and SNDMB grits appears to favour removal of individual alumina grains. The diff'erence between

the behaviour of the two forms of fractured

grit, i.e., the RDA type and the DXDA-MC/

SND-MB type, can be explained as follows. In

the RDA type grit, the fracture is governed

by the metal inclusions and other such bulk defects in the diamond particles; hence a large

number of very sharp points are presented to

the workpiece. For the stronger blocky grits,

however, the fracture is controlled by surface

energies, i.e., usually takes place along cleavage plains and hence the fracture points are

not generally as sharp as those found for the

RDA type grits. These "blunter" fracture

points have the ability to remove individual

grains.

e. Surface Finish

As indicated earlier, two spark-out passes,

without downfeed, were given at the end of

each test. The surface finishes obtained in this
the results of this preliminary investigation,

these complaints are not really justified. It must, however, be borne in mind that the grinding conditions used here are extremely severe,

when compared with conditions normally used

in the ceramic industry. Despite this fact, the

wheel costs are very low, the highest wheel cost

throughout this investigation being only 6.8

cents (U.S.) per cubic centimeter of material

removed, compared with a wheel cost of 17.6

cents (U.S.) per cubic centimeter of material

removed with a good wheel of the same size

grinding tungsten carbide at a downfeed of

only 0.025 mm (0.001 in). [6]. The lowest

wheel cost achieved when grinding alumina in

this test series was 0.11 cents (U.S.) per cubic

centimeter of material removed.

Table 5 shows some of the costs involved in

removing a cubic centimeter of the different

aluminas. Included with these figures are costs

for grinding tungsten carbide and titanium

under the same machine conditions but with a reduced downfeed, [6, 7] . Clearly, as the actual

removal rate decreases, time costs increase

substantially.

2.5. Conclusions

From this preliminary investigation it is difficult to draw very definite conclusions. However, one can deduce that 94 percent alumina

is generally easier to grind than either 85 percent or 96 percent. A wheel peripheral speed of

about 26 to 27 m/s (5100 to 5300 ft/min)

gives the highest efficiency with most grit and

bond types on both 85 percent and 96 percent

alumina. Metal bond wheels are more efficient

than resin bond wheels at all speeds tested.

However, it is possible that resin bond wheels

may be better in cases where edge chipping

presents a problem. With the three grades of

Table 4. Comparison of surface finishes achieved with different diamond types.

(96% alumina, 0.125 mm (0.005 in) downfeed and JfOOO rpm)

Grit Size

Surface roughness AA /j.m

After testing After finishing

DXDA-MC 80/100 3.00 1.68

DXDA-MC 140/170 3.38 2.68

RDA55N 80/100 3.35 1.65 RDA55N 140/170 2.75 1.75

MDA-S 80/100 2.62 1.75

MDA-S 140/170 3.42 1.82

SND-MB 80/100 3.68 1.81

SND-MB 140/170 3.80 1.85

DXDA 80/100 2.12 1.47

DXDA 140/170 2.37 1.25

way for the tests at a wheel speed of 26.6 m/s

(5240 ft/min) on 96 percent alumina are given

in table 4. At the conclusion of the G ratio tests, the

wheels were redressed to get a flat profile and

then given the following runs on the material

mentioned above and at the same speed :-

5 passes at 0.025 mm (0.001 in) downfeed.

2 passes at 0.01 mm (0.0004 in) downfeed.

2 passes at 0.005 mm (0.0002 in) downfeed.

2 spark-out passes.

The surface finishes obtained in this way are

also given in table 4. It is evident that the surface finishes obtained at the end of the G ratio tests are

far from good. However, one point shows up

clearly. Most of the metal bond wheels used on

ceramics at the present time contain natural

diamond and it is frequently stated that better

surface finishes can only be obtained with resin

bond wheels. The figures, however, show that

with DXDA in the metal bond, finer finishes

can be obtained than with resin bond wheels.

This is true even when the finishing technique

given above is used.

Although the quoted figures obtained using

the above technique are not yet considered good

for ceramics, one point must not be overlooked.

The Standards of the Alumina Ceramic Manufacturers Association [5] are based on a measuring instrument stylus having a 0.013 mm

(0.0005 in) radius. The figures given in table

3 were, however, measured with a stylus having a radius of 0.005 mm (0.0002 in) i.e.,

smaller than the average grain size of 0.011 mm

(0.0004 in) of the 96 percent alumina tested.

f. Grinding Costs

Complaints about the high costs of grinding

ceramics are very common but in the light of

109
Table 5. Comparison of grinding costs on aluminas, and titanium and tungsten carbide

Per cubic centimeter removed

VJX 1 u Size

Material

frronnd \ 4 A- C4f vXv Time

(min)

Time cost U.S. cents Wheel cost U.S. cents

Total cost U.S. cents DXDA-MC 140/170 AD85 368 0.83 16.6 6.8 23.4

RDA55N 80/100 AD96 4660 0.82 16.4 0.54 16.94

SND-MB 80/100 AD94 17170 0.82 16.4 0.15 16.55

MDA-S 80/100 AD94 19160 0.82 16.4 0.13 16.53

DXDA 80/100 AD94 23000 0.82 16.4 0.11 16.51

DXDA-MC 80/100 Titanium * 50 5.32 106.4 50.0 156.4

RDA55N 80/100 P20' 142 4.12 82.4 17.6 100.0

' Downfeed 0.02 mm (0.0008 in)

' Downfeed 0.025 mm (0.001 in)

alumina used, wheel costs are negligible, and

where a part is rigid enough, severe conditions

can be used without causing undue surface

damage.

3. Cylindrical Grinding

3.1. Introduction

Cylindrical components make up a fairly

large percentage of all ceramic parts ground.

Diameters vary from 3 mm (Vs in) to 300 mm

(12 in) and larger. Despite the fact that so much cylindrical grinding is done, the feeds

and speeds recommended by different suppliers

vary a great deal. It was therefore considered

advisable to look at the effects of these

parameters.

In this preliminary evaluation, three different diamond grits were used in metal bond

wheels to grind 75 mm (3 in) diameter X 300

mm (12 in) long rolls of three different alumina

grades. Two work speeds and two wheel speeds

were used to determine which combinations

would give the longest wheel life.

Machine

Wheel type

Wheel size

Wheel speeds

Wheel peripheral

speeds

Diamond types

Diamond mesh size

Diamond concentration

Bond depth

Infeed

Tota 1 infeed per test

3.2. Test Conditions

Jones and Shipman 1310

Cylindrical Grinder

DlAl

254 X 12.7 mm (10 x

1/2 in)

2200 and 2700 rpm

29.2 and 35.9 m/s (5750

and 7070 ft/min)

De Beers DXDA, SNDMB, MDA-S

80/100 U.S. mesh

100

3.1 mm (1/8 in)

0.025 mm (0.001 in)

Approx. 30 mm (1 3/16

in)

Workpiece rotation

Traverse rate

Coolant

Coolant flow rate

Specimen types

Specimen size

Sparkout

145 and 415 rpm

300 mm/min (12 in/

min)

Water plus Bryto 5

(100:1)

3 1/min (40 gallons per

hour)

Coors AD85, 94, 96.

75 mm O.D. X 300 mm

long.

(3 in. O.D. X 12 in long)

Two spark out passes,

without infeed at the end

of each test.

3.3. Test Results

The results of this series of tests are shown

in table 6.

3.4. Discussion

a. Grinding Efficiencies

The main purpose of this investigation was

to determine which combinations of wheel and

work speeds would give the highest G ratio and

hence the longest wheel life. Considering all

three alumina grades together, preliminary

tests showed that a low wheel speed with a high work speed, or a high wheel speed with

a low work speed are poor combinations. However, tests at low speeds for both wheel and

workpiece, and high speeds for both, yielded

substantially better results. The G ratios obtained with these two combinations are the

ones shown in table 6. Considering the individual alumina types

separately, it can be seen that the highest G

ratio on 85 percent alumina is achieved at high

work and wheel speeds. The most suitable grit

for this material is DXDA which gives much

better results than the other grits tested. On the

94 percent alumina, high work and wheel

speeds again yield the best efficiency but here

the MDA-S grit is superior. Recalling how, in

110
Table 6. Grinding ratios for cylindrical grinding

Work and wheel speeds 145, 2200 rpm 415, 2700 rpm

Alumina content 85% 94% 96% 85% 94% 96% —

Grit Size Grinding ratio Grinding ratio

DXDA 80/100 3470 6110 6620 8025 3940 5225

SND-MB 80/100 3130 4860 4125 2470 5020 6150

MDA-S 80/100 4970 7090 2910 4540 7260 3510

surface grinding, the 94 percent alumina appeared to be the "odd man out", with 85 percent and 96 percent being similar, it is surprising to find that the 96 percent alumina gives a higher grinding efficiency at low wheel and

work speeds in cylindrical grinding. It should

be noted, however, that the G ratios obtained

at high work and wheel speeds are not much

lower than those obtained at the lower speeds.

For this material, DXDA is again the better

grit type.

b. Surface Finish

Table 7 shows the surface finish that can be

obtained when grinding 96 percent alumina at

both combinations of wheel and work speed

under discussion. It is interesting to note that

the finer surface finishes are obtained when

grinding with high wheel and workpiece

speeds. This, then, is another reason why the

higher speed combinations should be used. AlthouGfh the SND-MB wheel gives the best finish,

the DXDA wheel gives an only slightly rougher

surface.

c. Grinding Costs

Once again, as with the surface grinding

tests, the wheel costs are extremely low, varying from 0.31 to 1.01 cents (U.S.) per cm^ material removed. The time costs in this case are

fairly constant and are of the order of 12 to

14 cents (U.S.) per cm^ material removed.

When these are compared with a wheel cost of

55.5 cents (U.S.) and a time cost of 16.5 cents

(U.S.) per cm^ of material removed, obtained

when grinding tungsten carbide rolls of the

same diameter at the same feeds and speeds

[8], it is once again obvious that wheel costs

in particular are negligible when grinding ce- ramics of the types tested here.

3.5. Conclusions

The conclusions that can be drawn from this

investigation are that a high workpiece speed

and wheel speed are generally the most suitable for grinding 85, 94 and 96 percent aluminas. DXI)A appears to be the better grit type

to use. Wheel costs are so low that they do not

influence the total product cost very much.

4. General Conclusions

Although it has been stated several times

before, it must be stressed that this is a preliminary investigation. Many of the results

obtained have merely been presented as found,

with no explanation of the strange behaviour

being given. What, it is hoped, has been shovra

is that, as with other materials, there are variations and exceptions to the general rules and

it is advisable to study each particular operation carefully in order to get the best result

from diamond abrasive grinding wheels.

5. References

[1] Busch, D.M., and Prins, J.F., A basic study of the

diamond grinding of alumina, these Proceedings.

[2] Eckert, J.D., A replica technique for examining

crystals and for retracing particles in grinding

wheels, Industrial Diamond Review, 28, 75-77,

(Feb. 1968).

[3] Eckert, J.D., and Caveney, R.J., A replica technique

for conventional and scanning electron microscopy. Journal of Physics E : Scientific Instru- ments 3, 413-414 (1970).

[4] Caveney, R.J., and Eckert, J.D., Electron microscopic study of grinding workpieces, Proc. Diamond Conf., Reading, Sept. 24-27, 1969.

[5] Standards of the Alumina Ceramic Manufacturers

Association for High Alumina Ceramics, 3rd

Edition, (1969) published by the Alumina Ceramic Manufacturers Association.

[6] Hughes, F.H., Grinding with diamond abrasives.

Table 7. Comparison of surface finishes achieved with different wheel and workpiece

speeds. (96% alumina)

Grit Size

Surface roughness AA fjun Low sneeds High sneeds

(145, 2200 rpm) (415, 2700 rpm)

DXDA 80/100 2.48 2.15

SND-MB 80/100 2.03 1.93

MDA-S 80/100 2.88 2.73

111
part 2: Tungsten carbide, De Beers Industrial 1: Steel, De Beers Industrial Diamond Division's

Diamond Division's Diamond Information L23. Diamond Information L22.

[7] Dyer, H.B., Grinding with diamond abrasives, part [8] Diamond Abrasive Technology Centre test result.

Discussion

Gielisse: I noticed an increase in 50 percent

of your cost in grinding AD 85 as compared to

AD 94.

Caveney: I think this was due to a different

grit type. Referring to the table of costs (table

5) the G ratio was very low for the resin bond

wheel used in the cost evaluation of grinding

AD 85.

Gielisse: The reason I asked the question

was the danger in this type of comparison.

Caveney: Well, what we are pointing out

here is the different G ratios. This one was

particularly low and it did make a big difference in the cost. The others which were comparable, ranging from 5000 up to about 23,000,

made very little difference in the cost. I think

that high cost was due to a very low grinding

efficiency.

Donovan: You are convinced that different

aluminas grind quite differently?

Caveney: Well, from our evidence, yes. You

can see from those curves (fig. 4 to 6) that

two ceramics exhibited similar behavior, but

that one of them had an enormous dip in the

G ratio with increasing wheel speed.

Donovan: What about the resin versus the

metal bond on alumina?

Caveney: Well, the behavior was more similar in the resin bond case than it was in the

metal bond case, I agree with you there. We

have to obviously extend this work in the next

phase of grinding triple 9 aluminas and also

grinding small parts where we can't use these

large infeeds.

Smith, C. J.: Did you do all this work dry?

Caveney: This was all done wet. The details are given in the paper.

Smith, C. J.: My question is, have you ever done anything dry?

Caveney: Not as yet. Do you suggest that

we do some work dry?

Smith: No, I was interested in grinding

ratio comparison, wet and dry.

Caveney: I see.

national bureau of standards special pub. 348, The Science of Ceramic Machining and Surface Finishing, Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

112
Ceramic Substrate and Specimen Fabrication

H. C. Leistner ^ and W. A. Wilson ^

National Bureau of Standards

Boulder, Colorado 80302

Procedures have been developed to fabricate ceramic materials used as substrates

for precision, coaxial, thin film resistors. Excellent radio frequency characteristics

have been obtained from this procedure in developing radio frequency standards at the

National Bureau of Standards at power levels of 10 milliwatts to 100 watts and

frequencies from direct current to 4.0 gigahertz. Substrates fabricated for this purpose

require precise tolerances and surface finish to insure uniformity of thin films deposited on a ceramic substrate. In addition, procedures have been developed to fabricate specimens used in determining the dielectric and magnetic properties of ferrites, insulating ceramics, glass,

fused silica, and single crystal materials. A variety of shapes are I'equired including

spheres, rods, disks and ellipsoids. Precise tolerances are required to obtain qualitative

property measurements of the materials involved. In both procedures, special techniques are used with existing equipment in grinding

ultrasonic machining, lapping and polishing to obtain the required end product. This

paper deals with both the application of the required products and the methods used

to fabricate these products.

Keywords: Ceramics; finishing; grinding; specimens; subtrates; ultrasonic machining.

1 . Introduction

In recent years the use of ceramic materials

has become quite prominent in scientific re- search and development programs because of

their characteristics and adaptability to special

applications. A number of applications have

evolved at the Boulder Laboratories of the National Bureau of Standards v^^hich have required

special consideration of available fabricating

techniques and equipment to accomplish the

desired end product. Most requirements were

accomplished through the use of surface and

cylindrical grinders and an ultrasonic cavitron

machine.

It should be noted that the Instrument Shops

Division at the Boulder Laboratories primarily

supports research and development type programs and as a consequence, normally fabricates "one of a kind" pieces of equipment or components as opposed to a multiple, production-type operation. The result is that, because

of the lack of specialized equipment to perform

a specific task, improvision and special tooling

is necessary to fabricate the required product.

With the increasing use of ceramic materials,

even more improvision is required and oftentimes an experimental approach to fabricating

techniques is necessary to meet the specifications.

The intent of this paper is to recognize some

of the tasks that have been accomplished in

fabricating ceramic materials for the support

of research and development programs at the

' Instrument Shops Division, National Bureau of Standards Laboratories, Boulder, Colorado 80302.

2 Figures in brackets indicate the literature reference at the end of this paper.

Boulder Laboratories. The examples cited are by no means all-inclusive, but are representative of approaches taken to machine and finish

ceramic materials with available equipment.

The areas covered include radio frequency

power measurement, utilizing coaxial, thinfilm resistors, and electromagnetic measurements of the dielectric and magnetic properties

of materials. These areas are dealt with separately to relate in further detail the approaches

taken to accomplish the required objectives.

2. Radio Frequency Power

Measurement

To meet the increasing complexity and higher

performance characteristics of radio frequency

power, developments have been accomplished

at NBS to more accurately measure radio frequency quantities. Using the calorimetric principle, Crawford and Hudson [1]^ developed a dual-flow coaxial calorimeter power meter as a

reference standard to accurately measure radio

frequency power. Quoting Crawford and Hudson, "The measurement of electrical power

using this principle depends upon the complete

conversion of the electrical energy as delivered

by a generator into thermal energy in a resistive load." To extend the measurement range

and further improve the radio frequency ref- erence standard, Crawford [2] developed an

rf-dc substitution calorimeter incorporating

automatic controlled reference dc input power.

An important feature of the calorimeter used

to develop these reference standards is the load

resistor which transforms the electrical energy

into thermal energy. The configuration most

suitable to this application is a thin film re113
sistor of either a cylindrical or conical shape,

matched to the mounting structure such that

the characteristic impedance of the line is

equal to the resistance of the load resistor at

any point along the length of the load.

The desired thin film resistor was fabricated

by depositing a metallic film on a ceramic substrate of the proper configuration. Thin films

used for resistors are typically about 4 microinches thick. In providing substrates for these

thin film resistors, precise dimensions and

smooth surfaces are important to insure uniformity of the film deposited on the substrate.

Thin spots in the resistant film cause the film

to operate at higher temperatures in the thin

area, resulting in short life or complete failure.

Any surface defects with sharp edges are disastrous. Consequently, all surfaces must be

polished to a finish of 20 microinches or better

and scratches more than a few microinches

deep are intolerable.

Figure 1 shows the configuration of a more

complex substrate fabricated from aluminum

oxide. Even though a variety of modifications

were used to meet different substrate requirements, this example will serve to illustrate the

approach necessary to accomplish this fabrication task. This substrate was ground between

centers on a cylindrical grinder, using a resinbond diamond wheel and oil-free cutting fluid.

A male center was ground on the small end

of the cone while a brass adapter epoxied to the

large end of the cone provided the female center for that end. These provisions were neces- sary in order to maintain concentricity of the

conical and cylindrical portions of the ceramic

substrate within a tolerance of ±0.0002 in

(±5.08 micrometers). The table of the grinder

was tilted to the proper angle to grind the

conical portion.

The grooves on the large end of the substrate

were machined with a surface grinder. A brass

adapter was fabricated to match the taper of

the cone and waxed to the ceramic substrate.

Using this adapter as a holding fixture, the

grooves were cut in the aluminum oxide on a

1.0

Figure 1. Conical substrate for a load resistor

used as the termination of a coaxial trans- mission line.

surface grinder with a resin-bond diamond

wheel.

Since the surface of the substrate was relatively coarse after grinding, the cylindrical and

conical portions were lapped, using a 4-8 micron grade diamond lapping compound to obtain a 10 to 20 microinch finish surface. With

the substrate fabricated to the closest tolerance

and finish possible, the load resistor was completed by depositing an alloy of Au, Ni, and

Cr 4 microinches thick on the cone portion of

the substrate and firing silver on both cylindrical ends of the substrate.

3. Dielectric and Magnetic Properties

For a number of years investigations have

been underway at the National Bureau of

Standards to measure the dielectric and magnetic properties of materials relative to radio

and communications applications. These measurements were made in a variety of ways to es- tablish absolute values and determine the effects of various shapes and orientations of the

sample material examined. The purpose of the

investigations was to establish reliable values

of these parameters for use in circuit design

and wave transmission calculations.

Bussey [3] has made a survey of the methods

used to measure these parameters, citing the

correlation that developed and the conclusions

that were resolved after many years of experimental and theoretical investigations.

The electromagnetic measurements of the

dielectric and magnetic properties of materials

required precise fabrication of sample specimens. Among the specimens examined were

small diameter rods, thin disks, spheres and

ellipsoids fabricated from a number of materials

including ceramic ferrites, alumina, fused silica,

and glass.

A variety of sizes of round rods were fabricated from the above materials. A typical ex- ample was the fabrication of a rod 0.039 in

(0.9906 mm) in diameter and approximately

1-1/4 in (31.75 mm) long. Occasionally diameters of 0.016 in to 0.020 in (0.4064 to 0.508

mm) were required. We did not have a centerless grinder available when this work was done;

however, some experimentation revealed that

such a rod could be fabricated satisfactorily on

a cylindrical grinder. A larger piece of stock

material was placed on a chuck and oriented

such that the axis of the stock material was

parallel with the cross feed of the grinder and

the grinding wheel set at a slight angle from

the centerline of the stock material (fig. 2).

By starting at the unsupported end of the stock

material and plunge grinding in one continuous

operation, the sample specimen was fabricated

within the desired tolerance, but more important, with a uniform diameter throughout its

114
entire length. The precise diameter required

was established by trial operations on an ex- pendable glass rod before starting on the final

test piece.

An auxiliary, motorized gear drive was con- nected to the cross feed of the cylindrical grinder to provide a slow, uniform feed during this

plunge grinding operation. The cross feed rate

was approximately 0.010 in (0.254 mm) per

minute. A metal-bonded, diamond grinding

wheel was used to grind the aluminum oxide

material.

In contrast with this approach, when attempts were made to plunge grind the specimen

from the side in multiple steps with the axis

of the grinding wheel parallel to the axis of the

stock rod material, control was not precise

enough to avoid small steps between each plunge

position of the grinding wheel. Since the diameter of the specimen was not consistently uniform, it was not acceptable as a test specimen,

and the plunge from the end approach described

above was developed.

Small disk specimens were fabricated from

aluminum oxide using a conventional surface

grinder and a diamond grinding wheel. A typical size was 2 in (50.8 mm) in diameter by

0.100 in (2.54 mm) thick. The tolerance on the

thickness was ±0.0003 in (±7.62 /.m). This

fabrication requirement was accomplished by

either using a vacuum chuck or attaching the

sample material to a base plate with wax and

holding the base plate in position with a magnetic chuck. After surface grinding was completed on both sides of the specimen, the final

desired tolerances and surface finish were obtained by hand lapping with a diamond lapping

compound.

Spheres were fabricated primarily by the

tumbling method, using equipment constructed

specifically for this purpose. Bond [4] introduced an early tumbling method for making

small spheres. The tumbling machine shown

in figure 3 consisted of a hollow tube oriented

perpendicular to a high speed, rotating, abrasive wheel. The specimen material, initially in

the form of a cube, was retained inside the tube

and formed into a sphere through operation

of the rotating wheel. A further development

in tumbling equipment resulted in the air driven machine shown in figure 4. The inside of the

abrasive wheel was formed to the desired shape

and provided a track for the specimen to follow

as it was forced around the inside of the wheel

by high velocity air flow.

Using these two tumbling machines, spheres

were fabricated from ceramics, ferrites, garnets and silica in a range of sizes from 0.016 in

(0.4064 mm) diameter to 0.120 in (3.048 mm)

diameter. It was possible to obtain spheres

with a roundness tolerance of ±0.0005 in

(±12.7 ;Lim). The size of the sphere was controlled by time of machine operation and, with

-TUBE

ABRASIVE WHEEL

MOTOR

f 1

I J

Figure 3. Motor driven tumbling

'machine with guide tube and

rotating abrasive wheel.

INLET

Figure 2. Plunge grinding a cylindrical specimen; arrows show how diameter is set, and

direction of grinder travel.

OUTLET

FLAT GRINDING WHEEL

CUP GRINDING WHEEL

Figure 4. Air driven tumbling machine

showing air flow and shape of abrasive

wheel.

''/ ''''' vm////,

115
some experience, could be held quite close to

the desired size. To further control the size

and finish of the sphere, a finishing tool was

developed as shown in figure 5. This tool consisted of a motor-driven, brass rod with a

conical depression on the extended end to

receive the spherical specimen. A similar bakelite rod was manually held against the specimen,

which provided a means for lapping the specimen to a more precise size and finish. Using

this procedure it was possible to control the

roundness of the sphere within a tolerance of

± 0.0001 in (±2.54/xm).

Ellipsoidal specimens were fabricated by

first milling the desired size and shape of lapping plate in either cast iron or brass material.

The required elliptical groove was obtained

by tilting the milling cutter with relation to

the alignment of the stock material. Both oblate and prolate ellipsoid lapping plates were

formed as shown schematically in figure 6. The lapping plate was securely attached to the

work surface of a conventional milling machine

with adhesive wax. The lapping procedure and

quantity of material removed was controlled

by adjusting the work surface and the spindle

of the milling machine. Thus, through the vertical and rotary motion of the milling machine,

the specimen material was lapped into the desired shape. With the lapping plates as shown

in figure 6, half of the ellipsoid was formed in

one setup. The specimen was then reversed

180° and the remaining half of the ellipsoid

completed. A variety of sizes were fabricated

for use in measuring the spin wave magnetic

properties of the specimen materials. A typical

example was an oblate ellipsoid 0.080 in (2.032

mm) major diameter by 0.040 in (1.016 mm)

minor diameter. Ellipsoidal specimens were

fabricated from ferrites and garnets, using

BAKELITE ROD

SPECIMEN

BRASS ROD

y/////////////////A//////////////'

lapping compounds appropriate for the material

involved.

4. Cavitron Machining

In citing the above machining approaches,

it is evident that other machining capabilities are necessary to meet the many demands re- quired in machining ceramic materials. An

ultrasonic cavitron machine is used to machine

various sizes and shapes of holes, slots or grooves in ceramic materials. With proper tooling and appropriate fixtures, very good control

can be obtained in machining ceramic materials

with a cavitron machine. A boron carbide slurry

has performed quite well in penetrating silicon

carbide, fused quartz and aluminum oxide, as

well as other brittle and hard materials.

5. Summary

In summary, the above examples illustrate

some of the approaches taken to fabricate ce- ramic materials for special applications. They

further illustrate the variety of fabrication re- quirements that are necessary to support a research and development organization. As

noted in these examples, the basic approach is

to use grinding equipment with diamond wheels

and diamond lapping compound to provide the

desired machining and finishing capability for

ceramic materials. In addition, the ultrasonic

cavitron machine is used to machine various

sizes and shapes of holes, slots or grooves in

ceramic materials.

The authors wish to acknowledge M. L.

Crawford, P. E. Werner, and H. E. Bussey for

SPECIMEN

1/4 PROLATE ELLIPSOID

LAPPING PLATE

1/2 OBLATE ELLIPSOID

LAPPING PLATE

Figure 5. Lapping tool used to finish

small diameter sphere specimens.

Figure 6. Lapping plates used for forming

ellipsoidal specimens.

116
the information they supplied in the preparation of this paper.

6. References

[1] Crawford, M. L., and Hudson, P. A., A dual-load

flow calorimeter for RF power measurement to

4 GHz, J of Res., Nat. Bur. Standard. (U.S.),

71C (Eng. and Instr.), No. 2, 1-178 (April-June

1967).

[2] Crawford, M. L., A new rf-dc substitution calorimeter with automatically controlled reference

power, IEEE Trans., Instrument and Measurement, IM-17, 378-384 (Dec. 1968).

rsi Bussev H. E., Measurement of rf properties of

materials, a survey, Proc. IEEE, 55. 1046-1053

(June 1967).

r41 Bond, W. L., Making small spheres, review of

scientific instruments, 22, 344-345 (May 1951).

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gahhersburg, Md. (Issued May 1972).

439-921 O - 72 - 9 117

On the Shaping of Brittle Solids by Erosion and Ultrasonic Cutting

H. L. Oh, K. P. L. Oh, S. Vaidyanathan and I. Finnic ^

General Motors Research Laboratories, Warren, Michigan 48090 and

University of California, Berkeley, California 94720

In mechanical shaping processes for brittle solids, materials is removed by the

propagation and intersection of cracks. To control the resulting surface finish, the

extent of cracking is localized by loading the surface in very small regions, usually

with small abrasive grains. Thus, in analyzing such processes as emsion and ultrasonic

cutting, we have to examine the fracture patterns produced by small hard indenters.

As a starting point, it is shown that fracture loads for indenters subjected to normal force may be predicted from conventional strength tests by using WeibuU's

statistical treatment of brittle strength. It is then shown that WeibuU's procedures

may be extended to predict the distribution of fracture location in brittle solids. For

the case of a spherical indenter under normal and tangential forces, many cracks may

form as the load is increased. It is shown that the location of the outermost of these

cracks may be predicted. Estimates may also be made of the denth of cracking and

thus the extent of the cracking produced by a single contracting particle may be

described.

These results may be applied directly to study the process of erosive shaping. The

influence on volume removal of changes in the velocity and size of the impacting

particles is predicted and shown to be in accord with experiment. Many other features

of the erosion process may be described including a transition to "ductile" behavior

in the case of glass when the impacting particles are small enough. The analysis of

ultrasonic cutting presents greater difficulty because the distribution in size of the

abrasive grains is of great importance. However, predictions may be made for the

influence of load, grain size, and material properties on the rate of volume removal and

these are shown to be in general accord -with experiment.

Key words: Abrasion: brittle-ductile transition; erosion; fracture location; ring-cracking; shaping; size-effect; statistical nature of strength; ultrasonic cutting; Weibull

distribution.

1 . Introduction

The shaping processes that we shall consider are those in which material is removed from a body by mechanical action to convert it to a

useful shape. When one attempts to develop an

analysis for such material removal operations,

it is often convenient to consider the two ex- tremes of mechanical behavior—the ideally ductile and the ideally brittle. An ideally ductile

material undergoes large plastic strains before

separation occurs and such materials may be

shaped by the displacing or cutting action of a

tool. By contrast, in an ideally brittle material

only elastic deformation occurs prior to fracture. Thus, in mechanical shaping operations

on brittle solids, material will be removed from

the surface by the propagation and intersection

of cracks ahead of and around the loaded re- gion. The concepts of ideally ductile and ideally

brittle behavior while somewhat over-simplified

do describe to a close approximation the behavior of many real materials and allow analytical solutions to be obtained.

With the reader's indulgence we will discuss

very briefly a few of the historical aspects of

• The authors are respectively. Associate Senior Research Engineer, General Motors Corporation Warren, Michigan 48090; Research

Assistant. Post Graduate Research Engineer, Professor of Mechanical Engineering and Chairman of the Division of Mechanical

Design, University of California at Berkeley 94720.

the shaping of solids before turning to consider some of the mechanical shaping processes

which may be applied to brittle materials.

The classical problem is of course, the shaping of brittle solids since stone was used for

tools and weapons and buildings by ancient

civilizations. Techniques were developed, during

the Stone Age and later, for loading a block of

stone and splitting off pieces of a desired shape.

The methods may appear primitive but fracture

in a brittle solid is hard to control and such

techniques could only have been developed after a great deal of careful experiment and experience. Much later in the pre-Columbian Americas the art of shaping brittle solids became very

highly developed. The Incas in Peru matched

massive stone with incredible perfection over

large areas while Aztec craftsmen are said to

have been able to produce 100 knife blades of

obsidian in an hour. Unfortunately, one cannot

find any scientific articles from the Stone Age

or even from the Aztec or Incas and what we

know about their shaping techniques is largely

based on the finished objects and in a few cases on construction sites where work was interrupted and never resumed.

The first technical literature one finds on the

shaping of soUds is in the 1850's. A number of

machine tools appeared in their present form

119
between 1840 and 1850. Steel was first produced in large volume in the 1850's and shortly

afterwards, technical papers dealt with the

cutting of ductile metals began to appear. Based

on a study of the early literature [1]- it appears

that analytical studies kept pace with practical

developments until about 1900. By that time,

the analysis of cutting forces in machining

had become almost as well developed as it has

todav. But then, hiorh speed tool steels were

developed, the practical art of metal cutting

took an enormous step forward and the analytical aspects appear to have been left behind.

Despite the many studies of metal cutting

which have been made in recent years, the

equations one finds in text books to relate tool

geometry, friction coefficient and the angle at

which metal is sheared, are essentially those

presented by Zvorykin [1] in Czarist Russia

in 1893. Unfortunately, the eauations are far

from adequate for accurate predictions. It may

be reassuring to those who are attempting to

understand brittle behavior to know that despite the modern equipment which is used to

shape ductile metals we are still lacking a basic

relation between the material properties and

the various geometrical features involved in

the cutting process.

Now, in recent years we find the shaping of

brittle solids coming back into the picture as designers wish to employ such materials as

oxides, carbides, and graphite. A single point

cutting tool or a single impact is rarely suitable

for machining such materials for it produces

large and uncontrolled fractures, perhaps leaving the object in several pieces. To our knowledge it has not been possible to control the age

old techniques of "knapping" or flaking brittle

solids with sufficient accuracy to make this a

precise shaping process. So what is done, is

to control the extent of fracture and the resulting surface roughness by loading the brittle

solid in very small regions using small hard

particles. Various processes are used and differ

essentially in the manner by which the particles are brought into contact with the surface. Thus

we have:

Erosion in which a stream of particles,

usually entrained in a jet of air, strike the

surface with a given velocity and direction

Ultrasonic Cutting in which a vibrating tool

forces particles contained in a slurry into the

surface being machined

Abrasion or Grinding in which the particles are embedded in a flat sheet or a wheel and

are moved across the surface while under a

load which tends to push them into the

surface

Lapping in which loose particles are pressed

' Figures in brackets indicate the literature references at the end of this paper.

against the surface being machined by an- other surface, the lap.

The terminology may vary but it is clear

that all these processes involve fracture under

small hard particles loaded by various combinations of normal and tangential forces. We

have listed erosion first and will treat it in detail because it appears to be the easiest of the

processes to analyze. In erosion the interaction

of a typical impacting particle and the surface

may be considered without reference to the

particle size distribution. For the other processes the applied load is distributed over all

the contacting particles and thus the loads on

the individual particles will depend on particle

size distribution.

In contemplating an analysis of erosion or one of the other processes, one thinks first,

perhaps, of an energy approach in which the

volume removal is taken as proportional to the

kinetic energy of the impacting particles or to

the kinetic energy lost by the impacting particles. This type of approach is often followed

in rock drilling studies and enables crude predictions of volume removal to be made. However, the values reported for "energy required

per unit volume crushed," i.e., the specific energy, vary greatly for a given type of rock and

it seems most unlikely that we could make

detailed predictions about the role of particle

size, particle velocity, impact angle, etc., on this

basis. Crushing studies on brittle solids show

that the essential energy requirement, that to

produce fresh surface, is a very small and unpredictable fraction of the total energy ex- pended in crushing [2]. So unless fragment

size and surface area can be predicted in advance, it would appear to be most unrealistic

to use a model based on energy to study the

shaping of brittle solids.

In the fields of rock mechanics and coal mining, analyses have been made of the chipping

that occurs under an indenting tooth or wedge.

The solutions apply only when the included

angle of the wedge is small, so that chipping

rather than crushing occurs, and do not appear

to be applicable to the case of fracture under

polyhedral or spherical abrasive grains.

Since volume removal in brittle solids occurs by fracture, we decided to study shaping operations by looking at the cracking that occurs when an idealized abrasive mrticle of spherical

shape loads the surface. This route leads us

into a study of the fracture of brittle solids

by spherical indenters. Some of this work will

now be summarized and then applied to study

shaping operations.

2. Fracture Under Spherical Indenters

We consider the case of a sphere pressed

against a semi-infinite solid by a static force P

120
or striking the solid with velocity U. For the

present we consider only normal loading (perpendicular impact) but will show later that the

approach may be extended to combined normal

and tangential loads. The elastic solution due

to Hertz and Huber is well known (e.g., it has

been summarized by Hamilton and Goodman

[3] ) and shows that the surface stresses, prior

to fracture, are compressive within the contact

area while outside, on the free surface the

radial stress is tensile while the circumferential

stress is of equal magnitude but compressive.

Below the surface the tensile stress decreases

very rapidly. Figure 1 shows the variation of

the radial tensile stress on the free surface with

radius. It is known [4] that the impact problem

is pseudo-static up to very high velocities and

hence the maximum load P corresponding to a given impact velocity U may be shown to be

P = (f)

^3/5 (pj3/5f76/5^2

2 2/? 1 —

+

where pi is the mass density of the sphere and

V and E are the Poisson's ratio and Young's

Modulus of the sphere (1) and the surface (2).

Fracture, when it first occurs, appears as a ring on the surface, usually well outside the

contact area, and spreads out below the surface

to form the frustum of a cone. If loading is

continued, the initial crack spreads deeper into

the solid and additional concentric ring cracks

form at successively greater radii. ^ Eventually,

at a high enough load, or velocity, extensive

RADIAL STRESS

AT SURFACE

Figure 1. Coordinate system used in study of ring

cracking.

' We have observed this type of cracking during tests on glass and by sectioning aluminum oxide specimens after testing.

crushing occurs under the indenter. This behavior is illustrated in figure 2 where we have

used a relatively large spherical particle for

analytical convenience. It is interesting to note

that despite the extensive amount of cracking

which occurs, the load-deflection relation is

estimated quite well by the elastic solution for

an uncracked solid. Presumably this is because

the stress state under the indenter is primarily

compressive and the cracked material is capable of transmitting compressive load. Thus,

we base our study of the shaping processes on

the type of cracking shown in figure 1 and

assume that the Hertz equations may be used

to relate deflections to load.

In attempting to predict ring cracking, there

is little value in the classical concept that a

brittle solid fractures when the tensile stress

reaches a certain magnitude. The load required

to produce ring cracking in glass shows a great

deal of scatter. In addition there is a strong

size effect such that if the indenter size is decreased from, say 1 in (2.54 x lO 'w) to 0.01

in (2.54 X 10-^m), there is a threefold increase

in the value of o-a, the stress at the rim of the

contact area, computed from the mean fracture

load. These observations are not new for about

30 years ago they led Weibull [5] to propound

his statistical theory of brittle strength. Basically, Weibull's idea was that flaws are distributed at random in brittle solids. These flaws

are, in general, too small to be detectable, and

so one cannot relate flaw size to strength as in

the familiar "Linear Elastic Fracture Mechan10

<

o

I

4

CL^ 2

0

1 1 1

HERTZ

RELATION / /^COLLAPSE

*^ EXPERIMENTAL _

RELATION FOR

STEEL SPHERE

INITIAL

1

FRACTURE

1 1

100

o

75

<Q.

50

<

>

25 o

0

0 12 3

CENTER APPROACH DISTANCE , lO^m

Figure 2. Variation of approach of center of 1/8 in

(3.175 X 10 'm) diameter steel ball to undeformed

region of glass plate with load. The impact velocities

corresponding to the applied load were calculated for

steel sphere from the Hertz equations.

121
ics" approach. Rather one has to test a large

number of specimens and determine the probability of failure for a given stress level and

specimen size. As the size of the stressed region

decreases the probability of finding a severe flaw decreases and as a consequence the mean

stress at failure increases. Strangely enough,

although the ring crack was what led Weibull

to his statistical treatment of brittle strength,

he did not analyze this problem in any detail.

Recently two of the present writers [6,7] studied the ring-cracking of glass in some detail

and showed that the mean fracture load, the

standard deviation of fracture load and the

mean location of the initial ring-crack for indenters of various radii could be predicted from

bending test data. Here, to illustrate the procedures followed, and to obtain a result which

we will use later, we compute the distribution

of the stress o-a at which ring cracking occurs. Following Weibull [5] we take the probability F((t) that a specimen of volume V

fractures under a uniform tensile stress < a as

( o-o-„ = 1-exp-F (_)

o->o-„

where a„, o-o, »„ the parameters of the probability distribution, are assumed to be constants for

a given material. If these constants are estimated, e.g., from a series of tension tests on

specimens of a given size, then this equation

may be used to predict the probability of failure at a given stress level for other specimen

sizes and loadings.

When the tensile stress, o-, is not uniform but

varies over the volume, then the expression

integrated over the volume stressed

in tension.

F(ct)=1- exp-

= 0

a->o-i,

In this form the Weibull distribution may be

used, for example, to make predictions about

the distribution of bending strength.

Glass is a brittle solid which is convenient

for experimental work. However, unlike polycrystalline ceramics in which the strength impairing flaws are distributed throughout the

volume, glass inevitably fails due to surface

flaws. Thus, if preceding equations are applied

to glass, V must be taken as the surface area

stressed in tension.

For ball indentation on glass the Weibull

distribution may be shown to take the form

/.(0=l-exp- {/ {^J^Ip^ 2. ^

A,

2iTrpdp

= l-exp-^ 7r(ii:cTo) =

( 1 V

i = o r_J_/W ^ m!(3-H)! / _ c^Y'+^-l

= 0

Ca > fit

O-a < CTu

and we have expressed the radial tensile stress

o- in terms of the maximum tensile stress o-a and integrated over surface areas which are

or have been under tensile stress. From this

distribution we may calculate, for example, 5-a the average value of va at fracture:

o-a = J"

<TadF{(Ta)

^^i'^'V^^lU (3 + m)!

1

2 ^ <Tu\ I + (constant) R '"^^ I (x)

After obtaining the parameters, o-„, <To,m, in this case from bending tests [6], the computation

may be carried out and gives the results shown

in figure 3. The agreement with experiment is

reasonable over a wide range of indenter sizes.

Another quantity which is of interest in shaping operations is the location of the outermost

ring crack. The subject of fracture location has

seldom been considered in fracture studies and

has never formed part of the Weibull theory.

However, by extending the Weibull approach

and looking at the probability of failure of

difi'erential elements which compose the solid,

we may derive the distribution of fracture location.

For a given load in ball indentation, if we

divide the region under tension into concentric

122
mail: glass

G[j = 34.4 MN/m2

CTn = 24.1 MN/m2

= 3.2

0.0001 0.001 0.01 INDENTER RADIUS , R, m

Figure 3. Predicted and experimental values of ca as a function of indenter size.

differential annuli, we can formulate the distribution of the outermost ring crack location as:

, foutermost ring crack "1 frob. |_Qgcurs at a radius a* J

, rthe annulus at location"! = Prob. La* fails J

X Prob. [

all annuli exterior l

to this annulus survivej

This formulation has been carried out in detail

in a separate paper [7] and it is shown there

that aside from elastic constants of the materials, the distribution of the outermost ring

crack locations G(a*) is also a function of

Weibull parameters m, cr„, o-o, the indenter ra- dius R and the stress level in the solid, aa. That is

1.0 u

u.< °g 0.8

z

'^^r. 0.6

>o

<o

q:

0.4

0.2

MAIL: GLASS

(Tu =34.4 MN/m2

O-Q =24.1 MN/m2

2 4 6 ^8

RADIUS OF OUTERMOST RING CRACK x 10 m

Figure 4. Predicted and experimental values for the

cumulative distribution of the outer ring crack when

a glass plate is loaded by a 3/32 in. (2.38 X

radius indenter under a load of ISi lbs (68i N).

= f((Tn, (To, m, <Ta, R).

Figure 4 shows the calculated result for G(a*)

when a glass plate, m=3.2, o-„ = 5000 psi (34.4

MN/m-), ,70 = 3500 psi (24.1 MN/m^), is loaded

by a 3/32 in (2.38 x lO-^m) radius indenter

up to a load of 154 lbs (684 N) . The agreement

between prediction and experiment is quite

satisfactory, particularly in view of the complex

nature of the multiple type of fracture being

considered. Again, knowing G{a*) we may

calculate various moments of the distribution.

For example, we make use later of (a*^), the

average value of (a*)^ which is:

(a*') = S d*^ dG(a*) = fA^u, <ro, m,oa, R) (2)

3. Analysis of Erosion

3.1. Spherical Particles Making

Perpendicular Impact

We are now in a position to analyze the ero- sion of brittle solids and consider first the case

of perpendicular impact by spherical particles.

By examining fracture patterns such as shown

in figure 1, and noting that the radius a* is

considerably larger than a, it seems reasonable

to take the volume damaged by a single eroding

particle as

(3)

where F{<Ta) is the probability of the particle

ring cracking the surface and S is the depth of

the primary ring crack as shown in figure 1. Examination of ring cracks in glass indicates

that the depth S of the primary ring crack is

proportional to the depth X to which the sphere

indent the surface.

From the Hertz equations we find:

X cc R (Ta X OC Jl -1/3 ^ ^[74/5 (4)

So thaW„ OC C/^/s and cx RU^i"- a*^ F(tr<,).

Since o.*^ is a function of m, o-„, ao, R and

o-a, where <Ja in turn depends on U we now have

an expression for the volume damaged by each

particle which is a function only of R, U and the

Weibull parameters (o-o, o-„, m). We examine

this prediction, which requires numerical computation to evaluate the integral, by varying the

Weibull parameter o-», holding the other four

quantities (o-o, m, R, U) constant. Effectively o-„, the stress below which fracture does not occur,

may be changed by inducing biaxial compressive stress in the surface. This was done by

shrink-fitting steel rings onto glass disks. The

predicted and observed values for relative volume removal are shown in figure 5 and agree

very well. As another test of our formulation.

123
LO

<

>

O

QJ

CC

o

>

UJ>

UJ 0.2

1 1 I 1

— -

•

— -

MAIL: GLASS

cru = 34.4 MN/m^

CTq =24.1 MN/m2

1 1

m = 3.2

1 1 0 30 60 90 120 180

INDUCED COMPRESSIVE STRESS, MN/nn2

Figure 5. Influence of biaxial compressive stress on

erosion of glass by 0.0165 in (4.19 X IQ-'m) di- ameter steel shot at a velocity of 270 ft/sec (82.29

m/s) and perpendicular impact.

we took erosion data reported by Sheldon and

Finnie [8] for three brittle solids which were

eroded by spherical steel shot of various sizes at

different velocities. The observed volume re- moval per particle is plotted against RU^^^ {a*~)

Fi<xa) in figure 6, 7, 8 for glass, magnesia and

graphite respectively. For magnesia and graphite (a*"-) was computed by integrating over the

volume stressed in tension rather than using

area stressed in tension as in the case of glass.

If the volume removed per particle is linearly

e

Q. >

d 10

cr

CC

LU

CL<

>

o

UJ

CC

o

>

MAIL: GLASS

10 10°

R X U

10'

X F(a-n

10'

Figure 6. Experimentally observed values of volume

removed per particle (for a wide range of velocity

and particle size) as a function of the predicted

damaged volume. Material: pyrex glass.

E

>

o 10

12

CE

cr

UJ

a.

> 10-'^

UJ

CC

Z)o

>

MAIL: MAGNESIA

rr2 10 ' 10°

RxU^^^x (a^^) X F(a-a)

10'

Figure 7. Experimentally observed values of volume

removed per particle (for a wide range of velocity

and particle size) as a function of the predicted

damaged volume. Material: magnesia.

10

Q. >

o 10

cr<

Q.

trUJ

CL<>

o

UJ

CE

10

12

> 10

_ 1 1 II 1

1 1 1 1

1 1 y 1 1 l_

SLOPE= 1.00

• /

MAIL: GRAPHITE

1 1 1 1 1 1 1 1 1 1 1 1 1 1 o' 10° _ 10' 10' R X u'*''^x {a**^) X F(cra)

Figure 8. Experimentally observed values of volume

removed per particle (for a wide range of velocity

and particle size) as a function of the predicted

damaged volume. Material: graphite.

proportional to the predicted volume damaged

per particle, we expect the data points to fall on

a straight line of slope unity. For magnesia and

graphite this relation is indeed observed while

for glass the points fall on a straight line of

slope 1.3. We speculate that the discrepancy in

the case of glass is primarily due to the fact

that the strength of glass in dynamic tests

(erosion) is higher than in static tests (ring124
cracking experiments). Possibly the use of the

Weibull parameters obtained from dynamic

tests would improve the prediction for glass.

3.2. Spherical Particles Striking at

Oblique Angles

The extension of the preceding formulation to

the case of oblique impact presents considerable

difficulty. When a sphere is pressed against a

surface with horizontal as well as vertical force

the ring cracks are distorted and at relatively

low values of horizontal force they become

"horseshoe" shaped (open in the direction of

sliding). The stress distribution prior to fracture is known from the work of Hamilton and

Goodman [3] . Since fracture tends to propagate

along a principle stress trajectory [9] (of the

stress state prior to fracture) it is possible to

compute the shape of these "horseshoe" cracks.

We have done this and assumed that cracking

terminates at a point where the maximum

tensile stress prior to fracture equals a value

<T„. On this basis it is possible to predict the

distribution of the "outermost" cracks as was

done for the ring cracks. Although the ideas are straightforward, the equations are cumbersome

and require extensive numerical computation.

The results have been given by K. P. L. Oh [10]

and we merely report them here. In the present

case we take the volume damaged by a single

particle as the area bounded by the outermost

horseshoe crack multiplied by the depth of

1.2 I

1

1

1 1 1 r

ANGLE oi , degrees

Figure 9. Predicted and observed values of

the relative volume removal as a function

of angle of impingement for glass eroded

by No. 110 steel shot at 200 ft/s (60.96

m/s). Predictions are made for three

values of /x, the coefficient of friction. The

inserted sketch shows the contact circle

and the horseshoe cracks.

cracking computed as before from the vertical

component of velocity. The only additional

parameter in the analysis is the coefficient of

friction between the sphere and the surface for

this determines the horizontal force. In figure 9

the predictions are compared with experiment

and it is seen that good agreement is obtained at

angles near 90° with a coefficient of friction of

/i = 0.03. This value for ^ is not unreasonable.

Typical values for rolling friction coefficient

between a ball and plate, including some shear

along the contact area, are 0.001 to 0.01. For

irregular geometry the quoted values are 0.05 to

0.2 [11]. At lower angles the decreased depth of

cracking due to the decreased vertical component of velocity becomes the dominant factor

and the choice of friction coefficient is not so important.

3.3. Angular Abrasive Particles Striking

Perpendicular to the Surface

In practice, erosion is often produced by

angular abrasive particles rather than spheres.

Since a detailed analysis of this probelm, similar to that developed for spheres, does not appear possible at the present time, we present

an approximate solution following the approach

suggested by Sheldon and Finnie [8] which was

later developed by Finnie and H. L. Oh [12].

Surfaces struck by eroding particles typically

have the appearance shown in figure 10. There

is relatively little damage under the particle

which is perhaps due to large hydrostatic compressive stresses at this point. On the surface

surrounding the indentation, a fairly extensive

region of material is cracked. Based on this

observation, we assume that the volume re- moved by an eroding particle, V^, is proportional

2R

U

Figure 10. Schematic view of cracking due to particle

impact.

125
to the damaged volume enclosed in the spherical

cap of radius i?* and depth S. That is

Vp oc R*

In case the particle is not a sphere but irregular

in shape, we may invoke two radii: radius R

which corresponds to a sphere of equal weight

as the particle and radius R' which is effective

in producing fracture.

Now, we have to estimate iS and R* and as

before, we shall take S as proportional to the

depth of indentation X obtained from the Hertz

solution. We consider /?* to be the radius of a much larger indenter than the actual particle

which on the average produces a ring crack at

radius a*. This assumption is perhaps not too

unreasonable because the cracked material is

capable of transmitting load. Thus we write

S oz X o: R* (era*) '

But from eq(l) we see that approximately

0"a oc R'

Eliminating o-a*, we arrive at a relation between

R* and X:

R' X

m+2

m-2

We relate the depth of indentation to the particle velocity, the particle mass and the radius

of curvature R' of the contacting region following the general approach given by Timoshenko

[4] to obtain

X oc [R^/R'y^'^ U'''

where TJ is the particle velocity. Finally, volume

removed per particle is:

oc X^i?* oc X

We distinguish two cases

spherical particle : R' = Ri

Vp oc R

irregular particle: R'

V„ oc R^/'

1-2 / \

-2

JJ-i/S I m-2

j

(3m-2 \

I

3 m-2 \

These expressions for the dependence of volume

removal on particle size and velocity have been

confirmed. In experiments on six brittle solids

over a wide range of velocities and particle sizes

[13] it was indeed found that the volume re- moval per particle could be written as Vp oc R" V for spherical steel shot

Vp oc R"' V for angular silicon carbide particles

A comparison of the predicted and observed

values of a, a' and b is given in figure 11. Except

for glass eroded by spherical steel shot the

correlation between theory and experiment is

good, particularly in view of the number of

assumptions involved.

3.4. The Brittle-Ductile Transition

An interesting consequence of the size-effect

on strength of brittle solids is that if they are

loaded on smaller and smaller regions, the

stress required to produce fracture may eventually exceed that required for yield. Thus, we

can understand the familiar observation that

microhardness or scratch hardness tests on

brittle soldis may produce flow rather than

fracture. In erosion the response of brittle and

ductile solids to changes in the angle of impingement are markedly different, as shown

schematically in figure 12. Thus a transition to

ductile behavior should be easily detectable.

To estimate the particle size at which ductile

behavior should be noticed in erosion testing of

nominally brittle solids, we need only the parameters of the Weibull distribution and indenoa.X

O-J

UJ>

Q

Z

<xa

oQ.X

Q

en

oc<

4 8 12 16 20

WEIBULL FLAW PARAMETER

Figure 11. Comparison of theoretical and experimental values of particle velocity and radius ex- ponents as a function of the Weibull flaw parameter m.

126
ANGLE OF IMPINGEMENT, , degrees

Figure 12. Schematic representation of the erosion

behavior of ductile and brittle solids. The maximum

erosion is taken as unity in both cases.

E E e ^ tn <£) 'o'o'o X X

J. OJ—

LlJUJ LlJ >> >

cn<n CO y, << < ^

CECE q: mm m << < oo o

(mO o

COo

tr

UJ

1 1 1 1 1 1 1 1

120 MESH^^

500 MESH^^ / ~ 1000 MESH-7 yS

/ // 1 y/

1 //

' //

///^ ^

0

1 1 1 0 10 20 30 40 50 60 70

ANGLE o< , degrees

80 90

Figure 13. Weight removal per newton of abrasive as a function of particle approach

angle a for glass eroded by silicon carbide

particles at 500 ft/sec (152.4 m/s).

tation hardness measurements. For glass, this

estimate can be made even more directly from

the data shown in figure 3. The mean value of (7„ (the stress at the rim of the contact area) re- quired for fracture is related to the average

stress p under the indenter through the Hertz

equations as: V = a.2/(l-2,,) ~ 4aa

So, for an indenter of radius 10"^ in (2.54 x

lO-'m) we. would estimate — 250,000 psi

(1,724 MN/mO andp ^ 10« psi (6,895 MN/m-)

from figure 3. Values of the flow pressure

P(yieid) for glass have been given by Marsh [14]

and are typically about 800,000 psi (5,516

MN/m^) for ordinary hardness tests and about

2,000,000 psi (13,790 MN/m=) for hardness

tests at low temperature where time-dependent

effects are absent. The points in figure 3 are of

course for static tests and would be somewhat

higher if dynamic loading had been used. So

direct comparison is impossible but it appears

that yield should occur before fracture for

particles in the range 10"^ in (2.54 x 10-^m) to

10"^ in (2.54 X lO^^m) in radius. Erosion tests

on glass show just such behavior. In figure 13 it

is seen that erosion tests on glass with particles

5 X 10-3 (127 X 10-^m)and 8 x 10^ in

(2.03 X in diameter lead to typically

brittle behavior while 3.5 x 10"* in (8.89 x lO-^m)

diameter particles show erosion behavior which

is typical of ductile metals. Other aspects of the

brittle-ductile transition in erosion have been

discussed by Sheldon and Finnie [15].

4. Analysis of Ultrasonic Cutting

In this shaping process a slurry of hard

abrasive particles and water separates a vibrating tool from the surface being shaped. Typically the tool vibrates normal to the surface at a frequency of about 25,000 Hz with amplitude of

the order of 10-^ in (2.54 x 10-^w). Thus the

particles are driven into the surface to produce

localized fracture. A small average force is

applied between the tool and the workpiece to

keep them close together during the operation.

Much has been written on ultrasonic cutting

including three textbooks and several comprehensive reviews with the most complete theories

having been given by Shaw [16] and by Russian

workers [17]. However, the statistical nature of

the strength of the workpiece does not appear to

have been considered. As Cook [18] points out

"It seems clear that until we understand smallscale fracture of brittle materials, we will not be

able to properly analyze this process."

We will now show that the material properties may be incorporated into an analysis of ultrasonic cutting and will base our treatment on

the cracked configuration shown in figure 1. Since the boron carbide particles, which are used in our experiments, and in most practical

ultrasonic cutting operations, are more nearly

spherical than angular we consider only the case

of spherical particles in analyzing this process.

The formulation of the expression for volume

removal is more difiicult than erosion for the

particle size distribution must be considered. In

erosion all particles may participate in removing material but in ultrasonic cutting only the

largest particles are active. As shown in figure

14, when the tool descends to its lowest position

at a distance c from the workpiece, only those

grits whose diameter ^ > c indent and remove

material, and then the depth of indentation

127
STATIC FORCE, F

ABRASIVE

y(t) = c+ Y(l-sin^)

Figure 14. Configuration involved in ultrasonic cutting.

varies from grit to grit. Because of this aspect

of the process, it is no longer appropriate to

consider volume removal on a single particle

basis as we have done in eq (3). Rather, we

should consider the average volume removal

max

d$ oc /(e) (5)

where ^>(^) is the grit size distribution, ^max

denotes the maximum grit diameter, (Imax, c)

denotes the range of the active particle diameter, and /(c) denotes the values of the integral, a function of c. As indicated by eq (2), (a-*) is a function of

""a and R, given values of the Weibull parameters. In turn, % may be expressed in terms

of X and R thru eq (4). The quantity /(c)

in eq (5) may therefore be evaluated once the

range of the active particle diameter is determined and if an independent relation between

X and R can be established. This may be obtained through a consideration of the mechanics

of the process as outlined below.

We assume that the tool oscillates with amplitude Y and period T at a distance

c + y(l-sin 2 TT -|r from the work piece with

an average force F exerted on it as shown in

figure 14. By geometery.

X + X, = ^ - c - Y\^l-sin

Also, from the Hertz equations,

P{t) OC X^'-R"'- OC Zt'V=/2'V2

It follows that

2H X OC OC ^ - c - r^l-sin

P(0 OC /2'V2|^^_c_y^l_sin^)]'''

When the tool reaches its lowest position c (i.e..

t = T/4) the depth to which a particle indents

the workpiece is

X OC (I-C)

and the impulse delivered is

(6)

SP{t) dt a: i^'^/^y'l^l-c-Y^l-sin^)

Since the momentum developed over one cycle

F T must equal the average total impulse delivered by the A^^ particles underneath the tool,

we have

dt

N

^max d^ii)

d^ (7)

A form for <i>(^) must now be chosen in order

to evaluate the integral. We chose a Beta function because under normal operations, only

particles with diameters close to ^^ax are active

and in the neighborhood of |,nax, the Beta function can describe the actual grit size distribution quite accurately. Thus,

* (I) = 1

\ ^max /

3+1

For example, it is seen in figure 15 that by

taking /? = 2.15, ^n,ax = 39.8 microns (39.8 x

lO '-m) a good description of the 500 mesh boron

carbide that we used is obtained in the upper

15 percent range. Employing a Beta function

for $ (^) in eq (7) and carrying out the integration, we have

F OC (13 + 1)

r(^ + l)r(3) N(W)'

r(/?+4) \ yJ

128
Figure 15. Particle size distribution of the boron carbide used in the

experiment.

where we make use of the fact that (Imax)^

is a constant for constant tool cross-sectional

area. The proportionality constant Ku involves

the cross sectional area of the tool, the slurry

concentration and the elastic properties {E,fx.)

of the tool, grit and workpiece. The value of c may therefore be evaluated for known values of

process variables and together with eq. (6)

would allow us to estimate 1(c) in eq. (5). We

have used the following data in our calculations

and preliminary experiments:

Y ^ 0.001 in (2.54 x 10-\,)

F = 12 oz, 16 oz and 24 oz (3.33 N, AAh N and

6.67 N)

Slurry concentration, water/abrasive by weight = 1.5

Tool cross sectional area = 0.08 sq in (5.16 X

Tool material: mild steel, E = SO x lO*' psi

(2.07 X 10^ MN/m^), ;a = 0.30

Grit material: Boron carbide 500 mesh, E =

42 X 10« psi (2.89 x 10' MN/m^), fi = 0.25,

size distribution as shown in figure 15

Work piece material: MgO and glass, with

properties as reported by Sheldon and Finnie

[8].

The results are shown in figures 16 and 17,

where we have plotted observed volume removal

against 1(c) for three values of F for glass

and magnesia. The data points fall on a straight

line of slope unity indicating that predictions are indeed satisfactory. We intend to conduct

more experiments in the future to confirm the

predictions over a wider range of the variables.

5. Quantitative Prediction of

Volume Removal

In our treatment of shaping processes we

have taken the volume removed by a single

particle as proportional to the cracked volume

produced by a single particle contacting a Q 20 1 III 1

1 1 1 UJ>O ro 2 E

10

1 UJ O

s — 8 SLOPE = 1.00—^

ZDo y

6

> d

UJ 4

ERV PER MATL-- GLASS

COm

o 2 1 III \ 1 1 1

10"•II 10

10 PREDICTED VALUE , I (c)

Figure 16. Observed volume removal as a function of predicted value,

1(c), for ultrasonic cutting of pyrex glass.

129
Q

LjJ 20

O ro ^ t

LU 10

1 Ld O 8 S ~

D

O

UJ ^ 4

> CE

CC UJ

UJ Q_ CO

OB 2

"1 1 r

10

10

SLOPE = 1.00

J \ L

MAIL: MAGNESIA

-J \ L r9

10

PREDICTED VALUE , I (c

-8

Figure 17. Observed volume removal as a function, of predicted value,

1(c), for ultrasonic cutting of magnesia.

smooth surface. Thus at least one test, using the

shaping process of interest, must be made on

the chosen material before quantitative predictions can be made. We took this approach

because the intersection of the fracture patterns from many particles will be a complex

process and more work needs to be done on this

topic before a complete quantitative treatment

of shaping operations is possible. However, it

may be of value to outline some preliminary

considerations in this regard.

Generally, one would expect that the extent

of cracking in any brittle solid would depend

only on the stress state in the part prior to

fracture on the fracture toughness of the material. This latter quantity could be expressed

by an effective surface energy for fracture

under plane strain conditions

in lbs G -(mN/m^) or by the related value of

the stress intensity factor Kjc Ibs/in^/^ {N/m^'^)

where

Gir = E

K 2 E

Normally, the elastic modulus does not appear in elastic solutions for the stresses but in

the present case, because of the change of contact area with load, we would expect the stress

state and hence the crack depth in the solid to

depend on the modulus as well as on the load

(or velocity) and the radius of the indenter.

In our analysis of shaping by spherical particles we took the depth of cracking as proportional to the depth of indentation. That is:

p2/3

where £7 is a weighted average of the moduli

of indenter and surface given by the Hertz

equations. A limitation of this expression is

that the surface energy for fracture is not included and hence the constant of proportionality must depend on Kjc (or G/c)

By contrast, Barenblatt [19] in discussing

the literature on fracture under a cylindrical

punch, where the contact area does not change

with load, showed that the diameter of the

base of the cone crack is proportional to:

2/3 p2/3

{GwEy^^

Very recently, we have made further studies

of the depth, S, of the initial ring crack in glass

loaded by spherical indenters. Rather than finding S proportional to the indentation X, i.e., S

pi/6

P'^'^/R^''^, recent work suggests 5 —zr-j^

Ri-/ ^ Hence from dimensional consideration, we

speculate that:

S

pS/6 ps/e

1/2

This new result will not change any of our

previous predictions significantly but does suggest that the coefficient of proportionality between the volume (a*^) X F (o-a) and the volume removed per particle should involve 1/

(Gic^'^E^/^). Representative values of Gw

are not easy to obtain for brittle solids. However, from the literature we obtained the following approximate values for the materials we

studied

130
in lb „ lb _i/3i7T_i/2\ Observed Damage Per Particle

^IC o— t!j - \Kjric ' Ej ' )

Glass 15x10-^ 10xlO« 3 9*

Graphite 5xl0-i 1x10" 3 7*

MgO 80x10-^ 30x10'' 1 1

*Relative to MgO as unity

Comparison of the last two columns shows that

this method of predicting volume removal is far

from precise but may be worth pursuing in

future work.

6. Conclusions

We have shown that the Weibull statistical

treatment of brittle strength may be used to

predict the loads required to produce fracture

under small indenters. An extension of the

Weibull treatment allows prediction to be made

about the location of fracture. In determining

the region that is damaged by a single spherical

particle we have seen that an important quantity is the outermost ring crack and that this

can also be predicted. Essentially, we use statistical considerations to determine the extent

of cracking on the surface and empirical ex- perimental observations to determine the depth

to which cracks extend below the surface. Improvements in the analysis could perhaps be

made by a better understanding of the interaction of cracks and the extent to which they

propagate. Recent work in this direction is discussed.

We have attempted to analyze these shaping

operations without any consideration of optimizing them, although this has been touched

on briefly in previous work [12]. It could be

argued that all shaping operations we have

discussed are relatively inefficient because they

involve loading the material predominantly in

compression and brittle solids are much stronger in compression than they are in tension.

Bailey and Dean [20] in a study of ice drilling

compared the "specific energy" required in various drilling methods. On this basis a man with

a pick was 100 times better than present drilling methods. Presumably he looks for a weak

spot, inserts the pick and loads the material in

tension. While we could shape brittle solids by

reflecting compressive elastic pulses from free

surfaces as tension waves this process would

be extremely difficult to control if precise shaping was desired. Thus, until the flaws inherent

in brittle solids can either be eliminated or

detected we will have to use shaping operations

which are inherently inefficient as material re- moval processes. Because of the "size-effect" in

strength, if smaller particles are used to obtain a finer surface finish, the processes become even

more inefficient and eventually with small

enough particles some brittle solids will behave in a ductile manner.

7. References

[1] Finnic, I., Review of the metal cutting theories of

the past hundred years, Mech. Eng. 78 (1956)

715-721.

[2] Walker, D. R., and Shaw, M. C, A physical ex- planation of the empirical laws of communition, Mining Eng. 6 (1954) 313-320.

[3] Hamilton, G. M., and Goodman, L. E., The stress

field created by a circular sliding contact, J. Appl. Mech. 33 (1966) 371-376.

[4] Timoshenko, S., and Goodier, J. N., Theory of

elasticity (McGraw Hill Book Co., New York,

N. Y. 1951).

[5] Weibull, W., A statistical theory of the stength

of materials, Handl., Stockholm, 1939). No. 151.

[6] Oh, H. L., and Finnie, I., The ring cracking of

glass by spherical indenters, J. Mech, Phys.

Solids, 15, 401-411 (1967).

[7] Oh, H. L., and Finnie, I., On the location of

fracture in brittle solids I—due to static loading. Int. J. of Fracture Mechanics. Sept. 1970)

287-300.

[8] Sheldon, G. L., and Finnie, I., The mechanism of

material removal in the erosive cutting of

brittle materials. Trans. ASME 88B, 393-400

(1966).

[9] Erdogan, F., and Sih, G. C., On the crack extension in plates under plane loading and trans- verse shear. Trans. ASME, J. Basic Eng.,

519—525 (1963).

[10] Oh, K. P. L., On the Statistical Nature of Brittle

Fracture, PhD. thesis in Mechanical Engineering, University of California, Berkeley,

(1970).

[11] McClintock, F. A., and Argon, A. S., Mechanical

behavior of materials, 665 (Addison Wesley

Publishing Co., 1966).

[12] Finnie, I., and Oh, H. L., An analysis of rock

drilling by erosion, Proc. 1st. Congress Int Soc. Rock Mechanics 2, 99-104 (1966).

[13] Sheldon, G. L., Erosion of Brittle Materials, PhD.

thesis in Mechanical Eng., Univ. of California,

Berkeley, (1965).

[14] Marsh, D. M., Plastic flow in glass, Proc. Roy.

Soc. (London) 279A, 420-435 (1964).

[15] Sheldon, G. L., and Finnie, I., On the ductile behavior of nominally brittle materials in erosive

cutting. Trans. ASME 88B, 387-392 (1966).

[16] Shaw, M. C, Ultrasonic Grinding, Microtecnic, X,

No. 6, 257-265 (1956).

[17] Kazantsev, V. F., and Rosenberg, L. D. The

mechanism of ultrasonic machining, Ultrasonic,

166 (1965).

[18] Cook, N. H., Manufacturing Analysis, 134 (Addison Wesley Publishing Co., Reading Mass. _ 1966).

[19] Barenblatt, G. I., Mathematical theory of equilibrium cracks, in advances in applied mechanics

H. L. Dryden and T. von Karman, editors

(Academic Press 7 1962 ) 55-129.

[20] Bailey, J. J. and Dean, R. C, Jr., Rock mechanics

and the evolution of improved rock cutting

methods, in Failure and Breakage of Rock Proc. 8th Symposium on Rock Mechanics AIMME,

New York, (1967), 396-409.

131
Discussion

Kirchner: Why do you use the volume

theory rather than the surface theory?

Finnie: For the case of glass we use surface.

V would represent an element of surface area. For the case of polycrystalline ceramics we

integrate over the volume stressed in tension.

Kirchner: How do you justify that?

Finnie: There is a fair amount of evidence

from people who have studied bending tests

that the flaws in polycrystalline ceramics are

distributed throughout the volume.

Rice: Would a small amount of plastic deformation prior to fracture explain, at least

partially, the location of ring cracks outside of

of the ball contact area?

Finnie: For the ball sizes for which we observed the location of fracture, the stress levels are such that we would not expect plastic deformation in the steel ball or in the glass.

Pr'^oIelTni\??s?^^

STANDARDS SPECIAL PUB 348, The Science of Ceramic Machining ani, Surface Finishinc,

of Standards hel/rr^^^^^ r^' u'^u"'^''^^TTt'' ^T.T' ^^""^ °^ ^aval Research, and the National Bureau 01 standards, tield at JNBS iNov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

132
Sonic Machining of Ceramics

William B. Campbell'

The Ohio State University

Columbus, Ohio 43210

The effect of 10 kHz sonic vibrations on the shaping and finishing of fired ceramic

shapes was investigated to identify optimum machining conditions commensurate with

product and surface quality. Sonic motors, developed at The Ohio State University and

with outputs up to 15 horsepower, were used to activate the machine tool. Significant

improvements in tool wear were observed in impact-tuned systems that produced

removal rates greater than 45 cu in per min. Template and profile tooling were

successfully used to obtain open cutting surfaces. Over 3000 test cuts provided the

data necessary to identify optimal parameters of operation.

Key words: Abrasive; ceramic machining; sonic machining; vibrational machining;

wheel machining.

1 . Introduction

Among the major methods for obtaining sur- face finish and complex configurations are mechanical finishing and shaping processes. For

ease of application and minimumization of process costs, most ceramic machining processes

are confined to green ware; however, manyfired products require final finishing or shaping. Vitrified grinding wheels, some floor tile

and refractory shapes are among present products requiring final finishing.

A recent program in the Department of Ceramic Engineering at the Ohio State University was concerned with the evaluation and development of machining processes for shaping,

and finishing vitrified grinding wheels. Sonic

cutting methods were evaluated to establish

conditions necessary for optimum machining

rates commensurate with product and surface

quality. The major parameters defining the

cutting conditions were:

1. Wheel speed (surface speed)

2. Feed rate

3. Degree of open surface

4. Uniformity of dimensions

5. Depth of cut

6. Tool shape

7. Tool orientation

8. Power source requirements

Over 3000 test cuts were required to identify

conditions which produced open surfaces and

uniform dimensions. For comparative purposes,

all data were evaluated against a production

removal rate of 20 in per min feed, 0.040 in

depth and 170 RPM resulting in an average re- moval rate of 8.8 cu in per min. Sonic cutting

produced acceptable product properties and ex- ceeded the current rate of removal in produc-

• Associate Professor, Department of Ceramic Engineering.

' Figures in brackets indicate the literature references at the end of this paper.

tion. Roll crushing cuts and profiles were simulated with a template tool; the depth and

degree of open surface surpassed present specifications.

2. Sonic Motors

Energy transfer by mechanical vibrations of

a single frequency has been termed sonic power

[1]^ Although these mechanical vibrations are

identical in character to acoustic vibrations, no

fundamental change occurs when their frequency is raised from the audible to the ultrasonic range. At higher frequencies, processing

applications are limited by material capabilities

because transducers typically operate with amplitudes inversely proportional to the operating

frequency. Inertial forces exceed material property limits because acceleration increases with

the square of frequency. In many practical applications where amplitudes of several mils

(0.001 in) are necessary, conventional transducer power sources are unsuitable [2]

A new transducer design [3-5] , developed at

the Ohio State University Sonic Power Laboratory, incorporates lead-zirconium titanate ce- ramic crystals mounted on a metal "horn" and

concentrates the vibrations at the tip. Operation at 10 kHz and 2600 volts is powered by a commercial rotating inductor generator and a

step-up transformer. A set of crystals, costing

less than fifty dollars [6], are connected electrically in parallel to produce a tip displacement of 0.0035 in with a measured delivered

power of 10 kw (15 hp) in the P-11 motor and

a delivered power of 500 watts (approximately

one-half hp) in the P-7 motor. The measured

efficiency of each motor exceeds 90 percent,

the former approaching 97 percent.

Two prototype sonic motors, figures 1 and 2,

illustrate the range of size and ratings cur- rently available. The one-half hp motor has

about one cubic inch of active piezoelectric ma439-921 O - 72 - 10 133
terial which activates the half wave horn to

dynamic conditions shown in figure 3. With its

steel resonator and force-concentrating unit,

the P-7 motor weighs about three pounds. The

larger motor, P-11, has a piezoelectric volume

of five and one-third cubic inches, operates with

a full wave dynamics shown in figure 4, and

has a total weight of 22 lbs.

Figure 1. One-half horsepower sonic motor of one-half

wavelength stepped-hom design.

Figure 2. Fifteen horsepower sonic motor with

catanoidal full wave horn design.

j Wavelength Steel Rod -| Wavelengtt) of Steel Full Wavelength Transducer and Horn Assembly ( Length = X )

Wavelength Stepped - Horn Tronsducer

Cotenoidol Horn

Figure 3. Dynamic operating conditions in half wave

horn transducer.

Figure 4. Dynamic operating conditions in full wave

horn transducer.

134
3. Machining Criteria

For most abrasive products the term grade

is used to define "the tenacity with which the

bonding material clings to the abrasive grains

in times of stress, and keeps them from breaking out [7]". This characteristic has been re- lated to a microstructural model by Kingery

and co-workers [8] . The similarity of our specimens to the constraints of their model permits

a direct analysis of machining forces.

Assuming the case of a particle having one

residual bond, figure 5, fracture will occur when the fracture strength of the bond is

reached.

a = Mc

max I where M = force movement

C = distance to outer fiber

/ = moment of inertia

and

a = 2RF

(1)

max 7r/iV4

The force for fracture is

F = iaf) S(R/h)

(2)

(3)

Since the load on one grain is inversely proportional to the number of grains in the work

areas, the critical force at the machining interface is

F' = Gas/k

(4)

(5)

where A; is a geometrical constant dependent on

wheel size, tip shape and work area, and

G = nh^/SR'V, (6)

is the analytical grade factor. Thus, the inherent bond fracture strength is more controlling

than the amount of bond present since is

constant.

Extending this relationship to examine elastic behavior, the critical force becomes

F" = iE,ee,)G/K (7)

Figure 5. Applied force F braking a single particle

from a wheel. [5]

If the critical value of ei, is exceeded, fracture occurs in the bond and the rate of fracture, or grain removal, will depend upon the rate of tip

movement or work area. Brittle bond fracture

will occur only when the elastic limit is exceeded; however, the application of vibrations to

the tool increases the applied stress. Thus, the

effective work area may be increased. Moore

and Kibbey [9] reported high rates of ceramic

tool wear for conditions of high amplitude and

frequency. Their tests were qualitative but indicated the direction of vibration and high frequencies to be most detrimental on ceramic

wear rates.

The only factor in eq (7) that is dependent

on system kinetics is the geometrical constant,

k. This constant will have different values for

differing kinetic conditions because it is related

to work area, which is related to feed and speed.

As the work increases, the force per grain de- creases until it is not sufficient to remove each

grain ; at this point tool wear increases and material removal decreases. With impressed vibrations the response is similar except that the

higher stress level (vibrational energy proportional to the cube of frequency) extends the

wear and removal units. Thus, an optimum feed

and speed for sonic machining is to be anticipated.

4. Conventional Machining Rates

The cutting rate reported as representative

of current production operations was calculated

to be 8.8 cubic in per min. This value is the

weighted average between the peripheral rate

of 15.2 inVmin and the spindle rate of 2.55

inVmin. The weighted average is equal to the

value obtained by weight loss measurements

and subsequent volume calculations based on

density. For actual cutting tests, the removal

rate was determined by weight loss and time

measurements, as well as by measurements of

dimensional changes. A comparison of these independent volume values indicated a two per

cent maximum deviation.

5. Tool Shapes and Orientations

Three tool shapes were evaluated for cutting

performance at different orientations. The tool

shape and optimum orientation are shown in

figures 6, 7 and 8. Tools were shaped from

standard carbide tips, No. TNG-334, Grade B106, Ultramet Corporation, Urbana, Ohio. The

135
Side Rake Angle

^Point Cutting

Side Relief Angle

TOP VIEW

Figure 6. Optimum orientation for tip 1.

(J^ of Wtieel

12°

i ^ , 1 End Relief -40'

12° Bacl( Roke TOP VEIW

Feed

-End Cutting Edge Angle

Top View

-Side Cuffing Edge Angle

- 30° Back Rake

Feed

SIDE VEIW

^L2° End Cutting Edge

Bock

Rake .

30 Side Relief

Feed 0° ' Side Rake

END VEIW

Figure 7. Optimum orientation for tip 2.

tips were brazed to the mounting cap or bar

prior to installation on the motor. The carbide

tips were satisfactory for the purposes of this

study; however, it is anticipated that further

improvements in both material and shape re- main for future investigations.

The following nomenclature and tool orientations were utilized in this study.

Rake angle: The wedge angle of the tool

at the point of workpiece

contact.

Relief angle: The angle between the tool

surface and the unmachined

surface which prevents conside View

Side Relief

End Relief Angle

Side Rake

Angle —^,

END VIEW

Feed

Figure 8. Optimum orientation for

tip 3.

tact (usually less than 10

degrees).

Cutting edge angle: The angle describing

the distance that, the nose

trails the leading edge.

Nose radius: Radius at the cutting tip.

The principal tool orientations are illustrated

in figure 9.

6. Optimum Speed

The effect of wheel speed was determined

over the range 150 to 400 RPM for each combination of tool shape, power source, machining

orientation and wheel grade. In all cases, the

optimum removal rate was achieved at speeds

shown below for six inch wheels. Above these

speeds, tool wear was excessive and

136
<{. of Wheel

! Counter-clockwise Rototion

TOP VEIW

30° Bock Rake

SIDE VEIW

2° End Cutting Edge

Feed END VEIW

2° Side Roke Due To Tool Weor 0° Original

Side View of Carbide

Figure 9. Principal tool orientations.

Motor RPM SFPM

P-7

P-11

200

300

50-310

80-475

removal rates slowly decreased; below these

speeds the rate of removal was not maximum.

Actual cutting speeds were controlled within

5 rpm.

7. Sonic Machining Rates

Sonic motors were mounted on a #2 Cincinnati Cutter and Tool Grinder driven by a variable-speed motor. A power drive for tool feeding was installed to obtain constant head

velocities. The arrangement for the P-7 motor

is shown in figure 10 and detailed in figure 11.

The crystal and electrode end of the motor

were contained in the box shown at the left of

each photograph. The motor was attached to

the box at the vibration-free node ring. The

carbide cutting tip is shown at the motor tip.

These positions were used for side cutting

where the tip moved vertically down the wheel.

Figure 12 shows the mounting for the P-11

motor and the impact tool mount. Supported at

the node ring, the P-11 utilized an extension to

contact the impact tool mount. As shown, the

Figure 10. P-?' sonic motor and tool mounted in

position.

Figure 11. Detail view of P-7 tool and wheel for

peripherial cutting.

template tool and orientation for groove cutting are apparent.

The basic differences between the two cutting arrangements are the horsepower ratings

of each motor and the use of an impact reed

mount. The P-11 is rated at 15 hp, whereas the

P-7 is approximately 1 hp. The impact tool

mount permits the cutting tool to obtain the

optimum translation frequency for machining

the workpiece; whereas, the rigid mount maintains the frequency of the motor at the toolworkpiece interface. The impact mount serves

to self-tune the tool to the frequency required

by the changing conditions of the workpiece in

side cutting. By comparison, the impact tool

mount resulted in a substantial reduction

(100% ) of tool wear. Neither the higher power

137
Figure 12. P-ll sonic motor with half wave extention

and impact tool mount.

motor nor the type of tool mount affected the

rate of removal. These observations are to be

expected in the machining of brittle materials.

Peripheral cutting rates were less than side

cutting rates and decreased with increased

depths. It is speculated that the geometry of

the tool contact at the periphery is responsible.

The important factor is the elapsed time to cut

the periphery and sides. The time is similar for

each operation because of the wheel dimensions.

The data are shown in figure 13.

The rates for side cutting with both motors

are shown in figure 14. It should be noted that

the volume is directly proportional to the depth

of cut; thus, indicating that sufficient energy

was available at all tested depths. The feed rate

maximized at 36 in per min and remained con- stant at higher rates. For 0.040 in depth, the

42

t; 36

30

24

18 -

6" Wheel

P-ll 300 RPM

P- 7 200 RPM

Tip 3

Data Data

0.080"

10 14 18 22 26

Measured Feed ( Incties Per Minutes)

10 14 18 22 26 30 34

Measured Feed [ Incties Per Minute ) Figure 14. Side cutting rates (A80-V).

maximum removal rate was above 25 cu in per

min, or about three times present rates. At the

maximum cutting depth, 0.080 in, 48 cubic

in per min were removed at a feed 36 in per

min.

8. Template Cutting

The simulation of roll crushing profiles was

accomplished by the template tool shown in figure 15. The tool was handfed at a slow rate and

typical results are shown in figure 16. The sur- face was very open and the detail exceptional.

Thanks are extended to Mr. Hudson Jayne

and Mr. Robert Meads for their technical as- sistance in all phases of the experimental program. The cooperation of the personnel of the

Sonic Power Laboratory, Ohio State University,

was most helpful.

Figure 13. Peripherial cutting rates (A8b-V).

Figure 15. Template cutting tools, new and 3 hours

wear.

138
Figure 16. Simulated roll crush cut using template tool.

9. References

[1] McMaster, R. C, Sonic power—university research

with an industrial payoff, news in Eng., College

of Engineering, The Ohio State University,

Columbus, Ohio, July, 1967.

[2] Houck, F., Vibration as an asset in design engineering, 66-MD-18, Am. Soc. Mech. Eng., (1966).

[31 McMaster, R. C, U.S. Patent 3,368,085.

[4] Libby, C. C, U.S. Patent 3,368,086.

[5] Minchenko, H. M., U.S. Patent 3,396,285.

[6] Belong, R. B., Sonic riveting, news in Eng., May

1968 (College of Engineering, The Ohio State

University, Columbus, Ohio).

[7] Lewis, K. B., Grinding wheel. Grinding Wheel Institute Cleveland, Ohio (1959).

[8] Kingery, W. D., Sidhwa, A. P., and Waugh, A.,

Structure and properties of vitrified bonded

abrasives. Bull. Am. Cer. Soc, 42, 5, (1963).

[9] Moore, H. D., and Kibbey, D. R., Development of

Accelerated Testing Procedures for Ceramic

Cutting Tools, E. E. S. Report 169 (The Ohio

State University, Columbus, Ohio) 1962.

Discussion

Bates: During the erosion machining of

these ceramics, was it a dry or a wet medium?

Campbell: All of this work was dry.

Bates: I've been doing similar work on

cavitation erosion of ceramic materials using a

similar device as yours. This is a magnetostrictive device, however, and I think that the use

of a wet media would introduce cavitation ero- sion and increase the overall erosion rate.

national bureau OF standards SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau'

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

139

Environment-Sensitive Machining Behavior of Nonmetals

A. R. C. Westwood and R. M. Latanision

Research Institute for Advanced Studies

Martin Marietta Corporation, Baltimore, Maryland 21227

Liquid environments can influence the efficiency of machining nonmetallic solids

in a variety of ways, e.g., by serving as lubricants, coolants, or particle dispersants.

More importantly, however, certain environments can markedly increase or decrease

the hardness of the nearsurface regions of such solids, and thereby exert a profound

influence both on the rate of material removal and on tool life. Because of its technological potential, the characteristics and possible mechanisms of this latter phe- nomenon—the Rebinder effect—receive primary consideration in this paper.

For crystalline ceramics, Rebinder effects in machining arise because of the in- fluence of the environment on nearsurface dislocation behavior. Effects resulting from

adsorption-induced changes in the surface free energy of the solid are of minor importance. Rebinder effects can also occur in noncrystalline solids, however, and recent

observations on such effects in various glasses are described. The possibility that these

effects are caused by a stress-plus-chemisorption-induced redistribution of sodium ions

in the near-surface region is discussed.

The importance of considering the total cutting system, environment-solid-tool,

in any account of environment-sensitive machining is stressed, for environments which

facilitate material removal when one type of tool is used can be detrimental to the

effectiveness of another tool with a different cutting action.

Key words: Adsorption; ceramics; drilling; environmental effects; fracture; glasses;

machining; mechanical behavior.

1. Introduction

Liquid environments can significantly influence the rate of material removal from ceramic

or glassy solids by an abrasive tool. Sometimes

the cause of their effects is relatively straightforward, for example, when they serve as coolants, lubricants, or dispersants. However, as Rebinder has pointed out [1] \ certain environments can act literally as "cutting liquids", be- cause they interact with the solid in such a manner as to facilitate its mechanical destruction.

Such effects can be quite spectacular, figure 1

[2]. Yet, despite the potential technological

value of this phenomenon—the Rebinder effect —a survey of the literature reveals that, apart

from their application to rock drilling in the

U.S.S.R. [1] and to a limited extent in the

U.S.A. [3, 4], the fundamental parameters underlying the environment-facilitated machining

of ceramic and vitreous materials have yet to

be studied in a sufficiently systematic manner

to allow predictability or optimization of cutting rates for any practical engineering material.

This situation makes the preparation of any

review of the "scientific aspects" of the influence of environments on machining on the one

hand relatively simple . . . because there is lit- tle to review, and on the other rather diflScult

. . . because the few investigations undertaken

' Figures in brackets indicate the literature references at the end of this paper.

to date have barely scratched the surface of

this field of study, and it is likely that many

important factors have yet to be considered

and evaluated. Thus, only a limited amount of

information exists upon which can be constructed any mechanistic understanding of the phenomena involved.

Recognizing this situation, the authors have

recently initiated a study of the influence of

simple organic liquids on the flow, fracture and

machining (primarily drilling) behavior of

such model solids as MgO and CaFo monocrystals and a conventional soda-lime glass. This

paper will be primarily concerned with a discussion of the results and implications of this

work. It should be appreciated that at this

stage in the development of the field, much of

what is said in regard to mechanisms should

be considered current opinion—not established

fact. Progress in understanding has been made,

and some general concepts can be stated with

reasonable certainty. For example, it has become apparent that adsorption-induced reductions in surface free energy are of limited con- sequence, while adsorption-induced changes in

the flow behavior of the near-surface regions

of the solid are of considerable importance.

However, it has also become evident that the

details of the mechanisms involved are sufficiently complex that it will be some time before

the specific behavior of any particular cutting

system (environment-solid-tool) can be predicted.

141
Alcohol

Alcohol

Alcohol

TIME , SEC

Figure 1. Illustrating the significant differences

in rate of bit penetration that can he produced

by changing drilling environment. The material is Corning "Pyrex". Diamond-studded

bit. Load 450g. [2].

2. Environment-Sensitive Machining of

Crystalline Ceramics

2.1. Mechanisms

The rate of penetration of a bit at constant

thrust into a ceramic body depends on (i) the

type of bit, and (ii) the hardness of the sohd.

Some bits are designed to cut most efficiently

through solids which exhibit a significant

amount of plasticity, and their operation primarily involves ploughing and shearing of the

material. Such bits include twist and spade

drills. Other bits are designed for use with more

brittle solids, and their effectiveness appears

to depend on how easy it is to initiate and

propagate a brittle crack through the material.

Such tools include diamond-studded bits, twocone rock microbits [3], and those tools which

remove material via repeated impacting.

Now, since the hardness of ceramic solids

can be significantly affected by the adsorption

of surface active species [5-7] , the observation

of adsorption-sensitive drilling or machining

behavior in these materials is to be expected.

Presuming, then, that such effects on machining behavior probably are a consequence of the

effects of environments on hardness, the first

step towards understanding the mechanisms of

environment-sensitive machining behavior must

be a clarification of the mechanisms underlying

the adsorption-sensitive hardness of nonmetallic solids.

The first example of the latter effect was

noted by Rebinder in 1928 [7], and in the intervening period several possible explanations

have been proposed for this phenomenon. (For

reviews see [8, 9] ) . The most widely known

hypothesis is that due to Rebinder himself, who

suggested that the decreases in microhardness

' However, preliminary data from studies of the slow growth of "stable" cracks in glass suggest that this process may be influenced by adsorption-dependent flow processes [2].

usually observed are caused by adsorption-induced reductions in the surface free energy of

the solid. This suggestion is based in part on

the premise that most of the energy involved

in creating a hardness impression in a brittle

material is utilized in creating new surface

area. There are several reasons why this hypothesis is questionable. First, in reality, only

a small fraction of the energy utilized in the

removal of material is required to generate new

surface area, typically less than 1 percent [10].

Most of the energy is absorbed in plastic flow

and heat generation processes. Second, adsorption-induced increases in hardness also have

been observed, both in crystalline solids such as CaFa [11], and in soda-lime glass [12]. Such

effects are thermodynamically unlikely on the

basis of the reduction in surface free energy

hypothesis [13]. Finally, it has been shown that

environments which produce significant influ- ences on drilling behavior have no significant

effect on the energy to propagate cleavage

cracks in MgO [14], or "super-critical" cracks ^

in a soda-lime glass [12].

In view of the limitations of the Rebinder

hypothesis, therefore, attention is now being

focused on the fact that in crystalline solids

which exhibit some degree of plasticity, the

deformation processes which lead to fracture

and removal of material usually involve the

generation and motion of dislocations in the

near-surface layers of the solid. Thus a feasible

basis for the occurrence of Rebinder effects is

that they are related to adsorption-induced

changes in the ease with which dislocations are

either able to (i) move and interact so as to

initiate cracks, or (ii) be generated at and move

around in the vicinity of crack tips, and so influence propagatability. Also, in recent years a

better understanding of the factors involved in

chemisorption on ionic solids, and in determining the mobility of dislocations within these

materials has become available, and these advances in knowledge have now been incorporated by Westwood, Goldheim and Lye (WGL)

[11, 15, 16] into an alternate mechanism for

Rebinder effects in crystalline ceramics. This

model assumes (i) that the mobility of dislocations in ionic crystals is controlled predominantly by interactions with such point defects as vacancies and impurity atoms, etc., [17-19],

and (ii) that chemisorption on such materials

involves—at least in part—electron transfer

between nearsurface point and line defects and

the adsorbate [20, 21].

The suggestion is, essentially, that adsorption-induced changes in the electronic state of

nearsurface point defects and dislocations

whether caused by direct interactions with the

adsorbate or as a result of band bending effects —introduce variations in their mutual inter142
MgO (a)

100

(t -2) , SECONDS

Figure 2. Variation of aL with log (indentation time,

t, mimis 2 s) for freshly cleaved Norton MgO crystals exposed to toluene or various aqueous environments. Indenter load = lOg. Room temperature.

Insert shows typical distribution of dislocation etch

pits around indenter impression. [15]

actions. Such variations are then manifested as environmentally-induced changes in dislocation

mobility, and hence in microhardness, i.e., as Rebinder effects.

The validity of this hypothesis has been ex- amined by studying the effects of adsorbed ions,

complex-ions, and organic molecules on the

room temperature mobility of nearsurface dislocation halfloops introduced by a diamond indenter into freshly cleaved surfaces of MgO

[15, 16] and CaFa [11]. The principal experimental variables were time of indentation and

composition of the environment, and the parameter measured was aL, where (aL = Lt —

Lj) ; Lt being the extent of edge dislocation

glide occurring during an indentation time t, and that during a two second indentation,

this being the minimum practical time for re- producible results. In effect, the larger the value

of aL after a given time of loading, the softer

the crystal. Values of Lt and Lo were determined by optical measurements of the distribution of edge-dislocation etch pits surrounding

the indentations (e.g., see the insert to fig. 2).

Figure 2 illustrates the variation of aL with

indentation time for MgO exposed to several

environments. When the surface was exposed

to water-free toluene, no significant dislocation

motion occurred following that produced within

the initial two seconds of loading. When the

crystal was exposed to moist air, water or an

aqueous salt solution, dislocation creep oc- curred, the edge dislocations moving some 3^

in 4000s. However, exposure to environments

' This is the "anomalous indentation creep" observed at room

temperature when MgO, AUO, or TiC, etc., are indented in a

laboratory atmosphere, and shown by Westbrook and Jorgensen [22] to be associated with the presence of adsorbed water.

Aq CoCIa PRESAT. AgCI

Aq. AgNOj PRESAT AgCI

Ag,Cr

14

^2 CLEAVED NORTON /

MgO

20 1 10"' 10-' [Ag]+ [01]- NORMALITY OF AQUEOUS ENVIRONMENT

o HjO - DMF SOLNS.

H2O - DMSO SOLNS.

ott—ai

10 1 10-2 PURE rr,..^^ DMSO [DMSO

MOLALITY OF AQUEOUS ENVIRONMENT

10-2 " 1 10 t PURE [DMF] DMF

Figure 3. (a) Variation of AL (i.000 sec) with

normality of aqueous solutions of either CaCL

or AgNOs presaturated with AgCI for Norton

MgO. The predominant complex species present in these solutions are indicated, (b) Variation of aL (1000 sec) with molality of aqueous

solutions of DMSO or DMF. Compare with

(a). Note similarity in form of curves. [16]

containing highly-charged complexes resulted

in significant increases in mobility. For a 7.2N

solution of CaCla presaturated with AgCI, the

latter being present predominantly in the form

of AgCli'- complex ions, edge dislocations propagated some 13/x in 4000 s. For a 17N AgNOs

solution containing Ag4Br'+ complex ions, dislocation loops propagated some in the same

period of time. The actual variation of the ex- tent of near-surface dislocation creep in some

fixed time t, ^L{t), with charge on the predominant complex species in the environment is

quite complicated, as is evident from figure

3(a).

Experiments [16] in organic liquid environments revealed that dimethyl formamide

(DMF) or dimethyl sulphoxide (DMSO) molecules absorbed on {100} MgO surfaces exert

the same influence on the mobility of near sur- face edge dislocations as do negatively or positively charged silver chlorocomplexes, re- spectively. This correlation may be seen by

comparing figures 3(a) and (b). It appears

from these data that DMF molecules act as electron-donor adsorbates with a respect to (impure) MgO, as AgClr^ complexes would be ex- pected to do, while DSMO molecules act as elec143
MgO

: / '. : / L

" A

1 1 1 1 DMSO 20 40 60 80 dmF

VOL % DMF

Figure 4. Variation of aL {1000

sec) with composition of DMSODMF environment at room temperature for two types of MgO.

Data points are omitted to improve clarity. Curve I is for "as

received" Norton MgO crystals.

Curve II is for hydrogen-treated

Norton MgO. Note shift in position of minimum A with state of ionization of impurities in

MgO [16].

tron-acceptor adsorbates, as would Ag^CV^ complexes.

The characteristic double minimum in the re- lationship between aL and environment can

also be produced in water-free DMSO-DMF en- vironments, figure 4, Curve I, and—as would

be expected on the basis of the WGL model

the exact locations of the minima at A and B

depend on the concentration and state of ionization of the impurities present in the material

[16]. In figure 4, for example, Curve I is for

"as received" Norton MgO containing, among

other impurities, '—150 ppm of iron predominantly as Fe^+. Curve II, on the other hand, is

from otherwise similar crystals which have

been heat-treated in hydrogen at 1650° for 24

h. prior to indenting to reduce trivalent impurity cations to the divalent state.

Another alternative mechanism for Rebinder

effects has been suggested by Westbrook and

Hanneman [23] , and this model is based on the

premise that the coefficient of friction between

an indenter and a crystal may be significantly

altered by the presence of adsorbed species.

Such an effect could lead to variations in the

complex distribution of stresses around the indenter, and so to a dependence of dislocation

mobility—and hence microhardness—on environment. However, studies [11] of the relaxation behavior of near-surface dislocations in

CaFo have revealed that active environments

influence dislocation mobility even in the ab- sence of stresses imposed by a loaded indenter.

Moreover, for any particular environment, a

correlation exists between the extent of dislocation motion induced by a loaded indenter and

the amount of relaxation which occurs after the

indenter is removed. It appears, therefore, that

indenter-lubrication effects are not fundamental

to the occurrence of Rebinder phenomena [11].

Rebinder effects can also occur in crystalline

semiconductor materials and in glasses. For the

former types of solid, it is well known that

chemisorption can lead to significant variations

in surface charge and charge carrier density

[20], and also that variations in bulk carrier

concentration can lead to variations in hardness (dislocation mobility) [24]. Thus, Rebinder effects in such materials may be a con- sequence either of electron transfer directly

between the adsorbate and localized energy levels associated with dislocation cores, or of the

chemisorption-induced bending of energy bands

in the vicinity of the surface which leads to

changes in the population of the core states of

dislocations [8] . However, it is not yet known

whether dislocation-like defects play any significant role in determining the flow behavior

of glasses, and current opinion is that this is

unlikely. Thus, a mechanism not dependent

upon the mobility of dislocations must be

sought to explain the existence of Rebinder effects in this class of solids, and some possibilities will be considered in 3.1.

2.2, Observations on the EnvironmentSensitive Drilling Behavior of Ceramics

The experimental observations reported in

2.1 demonstrate clearly that active environments can markedly influence the mobility of

nearsurface dislocations in ceramic monocrystals. The next logical step, therefore, is to determine the relationship between such effects

and the occurrence of environment-sensitive

machining behavior. With this in mind. Westwood and Goldheim recently undertook a study

of the adsorption-sensitive drilling behavior of

MgO and CaF. monocrystals [25]

a. Monocrystal Drilling Studies

These experiments utilized a precision drill

(30 rps) fitted with either a freshly sharpened

carbide spade bit (2mm diam.) or with a diamond-studded, hemispherical-headed bit (3mm

diam.). These two types of bit were selected

because their differences in cutting action might

144
be expected to lead to differences in response to

environmental changes. Carbide spade bits of

the type used in this work function more effectively when the solid exhibits some sensible

degree of plasticity. The diamond-studded bits

used, on the other hand, perform better in

harder and more brittle solids. Conceptually,

therefore, one might expect that an environment which softens the test piece would facilitate its drilling with a spade bit, but be detrimental to drilling efficiency when a diamond

bit is used.

The experimental procedure was as follows:

A freshly cleaved crystal was immersed in the

test environment and the bit lowered gently onto

its surface. A load was then applied to the drill

shaft by means of a lever system and the drill

started. Holes were drilled normal to the freshly cleaved surface of the specimen, and the

depth of penetration, D, under some fixed load

Figure 5. Influence of DMFDMSO environments on

(a) extent of edge dislocation motion produced during indentation period between 2 and 1000 s around

indentation in freshly

cleaved (100) MgO sur- face {16'] and (b) depth

of penetration of a carbide

spade bit in 600 s, D(600),

in [iOO] direction into

similar crystal. Room temp.

[25]. Note similarity in

variation of aL (1000) and

D(600) with environmental composition.

monitored as a function of time by an LVDTrecorder system.

Some results from drilling experiments with

carbide spade bits are presented in figures 5

and 6. Note first that certain DMSO-DMF solutions which soften (increase aL) and facilitate

drilling in MgO with respect to its behavior in

water, e.g., 50 v/o DMF in DMSO, figure 5, • harden and make drilling more difficult in

CaFo, figure 6. It follows that the variations in

drilling efficiency shown in figures 5(b) and

6(b) are not related to differences in coolant

properties of the solution environments. On the

other hand, a comparison of the data in figures

5(a) and 5(b), and in 6(a) and 6(b), reveals

that aL(IOOO) and the depth of penetration in

600 sec, D(600), vary in an essentially identical manner with composition of the environment surrounding the crystal. It is apparent,

therefore, that, when a carbide spade bit is

used, the efficiency of the drilling process in

both MgO and CaFa is related directly to the

mobility of near-surface dislocations in these

materials. A similar correlation between drilling efficiency and dislocation mobility has also

been noted when dilute aqueous solutions of

AICI3 (- 0.1 w/o) are used with MgO [25].

Such solutions have for some time been recogDMSO 20 40 60 80 VOL % DMF

Figure 6. Data similar to that in fig. 5, but for CaF:.

This material cleaves on

fill) planes and was

drilled in a direction perpendicular to this plane

[25] . Dislocation mobility

data from ref. [11].

145
Figure 7. (a) Influence of various environments

on D(600) in MgO, and of n-alkane environments on AL (1000). Nc is the number of

carbon atoms in the molecule [2, 25] —(b).

Effect of binary alcohol solutions on D(600).

Note that solutions of butyl alcohol in methyl

alcohol can reproduce the maximum in D(600)

occurring in propyl alcohol (Nc = 3). [2].

nized by Rebinder [1] as hardness reducers for

drilling quartz and granite.

The influence of n-alkane environments on

the drilling and near-surface dislocation behavior of MgO and CaF, monocrystals has also

been studied [25], and it has been found that

both aL(IOOO) and D(600) increase regularly

with chain length for the series from n-heptane

(number of carbon atoms, Nc, = 7) to n-hexadecane (Nc = 16), figure 7. In n-alcohol en- vironments, hovi^ever, a maximum in D(600)

occurs at Nc = 3, figure 7 [2] . Such an effect

is reminiscent of an effect noted previously in

toluene-dodecane solutions (fig. 8, Ref. [25]

)

and, on the basis of the earlier work, it implies

that alcohol environments act as electron donor

adsorbates with respect to MgO surfaces. Also

of especial interest is the observation that by

mixing together alcohols oi Nc = x and Nc =

If (in this case a; < < 5), it is possible to re- produce the characteristic variation of D(600)

for alcohols having values of Nc between x and

y, figure 7(b). The generality of this phenomenon will become evident from other results to

be presented below, and a possible interpretation is suggested in 3.3.

In experiments conducted using diamondstudded bits, the efficiency of drilling was as- sessed not in terms of penetration after a given

time, D{t), but in terms of the average rate of

drilling during the period 150-250 sec, this being designated D(200). It was possible to use

this more meaningful parameter because the

rate of drilling with diamond bits was relatively constant after the first 100 sec. As mentioned earlier, diamond bits drill relatively

hard and brittle materials more effectively than

they do softer and more ductile solids (an attempt to drill brass or aluminum with a diamond bit will provide proof of this). Thus, one

might expect the observed tendency for drilling

efficiency to increase with Nn for alkane and alcohol environments having Nc > 5 when spade

bits are used (figure 7) to be reversed when

diamond bits are used. Furthermore, one might

expect that, for the diamond-studded bits, the

alcohols would be superior cutting environments to the alkanes when Nc is < 10, and inferior when Nc > 10. These predictions are borne out by the results shown in figure 8. Note

that the maximum in D(600) for alcohol en- vironments having l<Nr<4, figure 7, is re- flected as a minimum in D(200) in the data of

figure 8. Note also the reversal in relative cutting efficiencies of toluene and water for the

spade bits and diamond-studded bits, cf. figures

7 and 8.

b. Other Observations on Ceramic Monocrystals

To examine the possibility that hardnessreducing environments influence crack initiation in notch-sensitive solids, comparative

scratch tests were performed on freshly cleaved

{100} MgO surfaces exposed to air (i.e., adsorbed water) and surface active n-hexadecane

[23]. A Kentron microhardness machine was

used, with a 20g load on the indenter, and a

traverse rate of 0.03 cm /sec. The results are shown in figure 9. Cracks were not formed in

surfaces exposed to air, figure 9(a), except

where the track crossed a cleavage step. Under

hexadecane, however, cracking was prolific, fig- ure 9(b), fracture occurring on both fllO)

planes intersecting the surface and {100)

planes below the surface. Thus, under conditions in which material is removed via ploughing, it is clear that "softening" environments

facilitate crack initiation in MgO.

c. Drilling of Polycrystalline MgO

Figure 10 presents some results from drilling

experiments on polycrystalline MgO (Kodak

CO 6 E

a. o

o

MgO • -n- Alcohols

* -n- Alkanes

Water or Toluene

1 I L Water 2

Toluene

Figure 8. Influence of environmental composition on rate of drilling after 200 sec, 0(200),

with a diamond-studded bit in MgO. Note that

the characteristic variation in D(200) for the

n-alcohol and n-alkane environments is the converse of that obtained when spade bits are

used; cf. these data with those in fig'ure 7. [2].

146
a

Figure 9. Indenter tracks in freshly cleaved

MgO surfaces produced (a) in air (i.e., with

adsorbed water -present), and (h) under sur- face-active n-hexadecane. Loan on indenter was 20g, and it was moved from left to right.

Room temperature. Note prolific cracking in- duced by softening hexadecane eii ^'ronment.

[25].

12 - Polycrystalline MgO

£ In n-Alcohols o D(200)- Diamond Bit £ 10- D(600)- Spade Bit o

o

S 8-

FiGURE 10. Influence of n-alcohol environments and type of bits on efficiency of

drilling polyo-ystalline MgO. Note opposite

influences of environment for the two bits. [2].

IRTRAN) in n-alcohol environments [2]. These

data demonstrate again the necessity of con- sidering the total cutting system, environment —solid—tool, in any discussion of environmentsensitive machining behavior. Clearly, environments which facilitate cutting with a spade bit

tend to reduce the effectiveness of a diamondstudded bit.

2.3. Comments on the Environment-Sensitive

Drilling of Ceramics

A liquid environment might influence drilling efficiency in several ways, for example, (i)

as a consequence of its lubricant or coolant

action, reducing bit wear, (ii) by decreasing

the surface free energy of the solid, or (iii) by

changing the flow and/or fracture characteristics of the solid. Certainly, cooling effects introduced by liquid environments are beneficial

to cutting. It has been noted, for example, that

drilling MgO with a spade bit in air is very

inefficient. In these experiments the bit became

blunt within 15 s after very limited penetration. In surface-inactive toluene, on the other

hand, the drill penetrated three times further

before blunting occurred [25]. However,

neither the coolant nor the lubricating action

of environments are considered of primary importance in controlling environment-sensitive

machining behavior because (i) certain environments which facilitate drilling in MgO, say

50 v/o DMF in DMSO, make drilling more difficult in CaFo (of figs. 5 and 6), (ii) additions of

an alkane to toluene improve drill penetration

in MgO but reduce it in CaF^ (compare figs.

6(a) and 8(b) in Ref. [25]); and (iii) Selim,

et al. [26] have demonstrated that the addition

of certain organic compounds to water which

facilitate the cutting of quartzite actually in- crease the coefficient of friction in drilling. As

mentioned in 2.1, there are also valid reasons

to reject the view that variations in drilling

behavior are related to adsorption-induced re- ductions in the surface free energy of the solid.

It does appear, however, that the results of

several investigations [1-3, 25-6] are consistent with the view that Rebinder-type effects in

drilling are associated with adsorption-induced

variations in the flow and flow-dependent fracture characteristics of nonmetallic solids. In

particular, we consider that the correlation between drilling eflficiency and near-surface dislocation mobility demonstrated by the data for

carbide spade bits in figures 5 and 6 does represent a real and important relationship between these two properties. A possible explanation for this correlation is that for a tool

which removes material primarily by ploughing and shearing, any environment which facilitates dislocation motion in the nearsurface

region of a notch-brittle solid will effectively

facilitate crack initiation, and hence material

removal. Conversely, for a tool which operates

primarily as a consequence of repeated high

energy impacts and brittle crack propagation,

a softening environment will be detrimental to

cutting, causing crack blunting and reduced

propagatability.

Robinson's observations on the environmentsensitive drilling behavior of Indiana limestone

[3] may also be interpreted in this manner. He

found that 2 percent solutions of sodium

adipate increase slightly ( + 5 % ) the compressive yield stress of this limestone with

respect to its value in water, whereas solutions

147
of sodium azelate reduce this stress (—19%).

In comparative drilling tests using a microdrag

bit—which causes failure to occur predominantly by plastic shearing processes—he found

that sodium azelate solutions increased the rate

of drilling by —- 70 percent with respect to that

in a pure water environment, while sodium

adipate solutions reduced it by about 28 percent. However, the use of a two-cone rock

microbit, which is considered to cause rock

removal primarily by brittle failure processes,

resulted in smaller but opposite influences on

drilling behavior.

Unfortunately, very little work of a mechanistic nature has been conducted on the environment-sensitive machining behavior of polycrystalline ceramics, despite its potential

practical value. Thus, at this time, it is only

possible to speculate on some of the possible

modifying influences of internal surfaces.

It is known, of course, that grain boundaries

can serve as relatively "soft" sources, so that

many of the dislocations involved in plastic

flow in polycrystalline solids are likely to be

produced at sites remote from the free surface.

In this case, environments are less likely to

exert a significant influence on deformation

behavior. Also, by acting as stable obstacles to

dislocation motion, grain boundaries are known

to play an important role in crack initiation

processes in ceramic solids. Whether the process of crack nucleation at grain boundaries is or is not affected by environmental conditions

will depend on a number of factors, including,

for example, the influence of non-equilibrium

concentrations of point defects and impurities

at or near the boundaries [30] on local adsorption behavior. Such segregation could affect the

local relaxation or concentration of stresses by

dislocations.

Other possible effects could arise from the

known anisotropy of adsorption on different

crystallographic surfaces—in this case, different grains [31]. It has been shown that this

phenomenon is responsible for the apparent

anisotropy in microhardness when nonmetallic

crystals are tested in a moist atmosphere; the

effect disappears when the tests are conducted

on dry, clean crystal surfaces [9, 31]. Nevertheless, such an anisotropy of adsorption could

conceivably cause one particular grain to be

hardened, while its neighbor is softened.

In short, the as yet unknown influences of

such factors could significantly hinder the rapid

transfer of understanding gained from monocrystal studies to an interpretation of the behavior of polycrystalline ceramics. It is evident

that mechanistically oriented studies of the en- vironment-sensitive machining behavior of

polycrystalline ceramics of various compositions, purities and grain sizes are urgently re- quired.

3. Environment-Sensitive Machining

of Glasses

3.1. Mechanisms

Explanations for the environment-sensitive

fracture behavior of oxide-type glasses usually

fall into one of two categories (i) those dependent upon adsorption-induced reductions in sur- face free energy [7, 32], and (ii) those which

assume the presence of water to be critical, and

of which the stress-enhanced corrosion mechanism of Charles and Hillig [33] is perhaps most

widely accepted. For explanations in category

(ii), the role of testing media other than water

is usually considered to be simply that of

screening the highly reactive water molecules

from the glass surface [34-6]

However, there is ample evidence that, under

conditions where fracture is induced by abrasion or grinding, certain organic environments

are considerably more active than water [37-8]

For such examples of environment-sensitive

fracture behavior it would certainly appear

that explanations dependent upon the dominant

presence and corrosive influence of water could

not be valid. However, it does not follow that

such effects must therefore be caused by adsorption-induced reductions in surface free

energy, because it is well known that flow can

occur in certain glasses at room temperature

[39-42] . Thus, an alternative possibility is that

certain examples of the environment-sensitive

fracture behavior of glass may be a conse- quence of adsorption-induced variations in the

flow and flow-dependent fracture properties of

the near-surface regions of this solid [12].

As mentioned earlier, current opinion does

not favor the view that dislocations, or dislocation-like defects, are involved in flow processes

in glass. Indeed, very little appears to be known

about the actual mechanisms of flow, or of

flow-sensitive fracture, in this class of solids.

Of course, this means that any mechanism proposed for the environment-sensitive flow

and/or fracture behavior of glasses must of

necessity be speculative in nature. Recently,

however, the authors [12] have considered the

relative merits of a model based on the fact

that stress [43] and potential [21, 44] gradients can cause redistribution of the sodium ions

in the near-surface layers of a glass. Recognizing further that the flow of sodium ions towards an electrode causes an increase in volume

of the glass there [45], then a feasible con- sequence of any stress-plus-chemisorptioninduced flow of sodium towards the surface of a glass specimen should be some relaxation of

any tensile stresses created there [46] , and by

virtue of the increased concentration of sodium

ions, a decrease in the hardness of this region.

Then, depending on the mode of action of the

148
cutting tool, such a change could either facilitate or hinder machining. Presumably any en- vironment which interacts to produce a positive

surface charge on the glass should tend to induce sodium ions to diffuse away from the surface. Such movement should then result in the

opposite influences on drilling efficiency, dependent upon the tool.

The possible relevance of this mechanism,

and of the surface energy reduction and stressenhanced corrosion models, to the environmentsensitive machining behavior of several glasses

is under current investigation [2] . Some of the

results obtained so far v^ill nov^ be described.

E

5-

4-

3-

• -n-Alcohols A - n-Alkanes

- Water or Toluene ^~~a

/ Water 2

Toluene 6 8

Nc

10 12 14 16

Figure 12. Variation in fracture energy, 0, of a

soda-lime glass at 25 °C with environmental composition. 0 determined by double cantilever technique. [12].

3.2 Observations on the EnvironmentSensitive Drilling Behavior of

Oxide-Type Glasses

This v^'ork has been concerned with the influence of water, toluene, the n-alkanes, nalcohols, and methyl esters on the drilling behavior of various glasses, and, to date, most of

the experiments have been conducted on a con- ventional soda-lime glass containing 13 percent NaoO and 9 percent CaO. With one exception, all of the data to be discussed were

obtained using the hemispherical-headed, diaEo

o

z

O

i LI

3-

1 -

i i i

J- I L

f Water 2

Toluene

• - n- Alcohols

- n- Alkanes

- Water or Toluene

Figure 13. Water content of environments used to produce data shown in figures 11, 12, H.-17 [121.

The Karl Fischer technique was used in these

measurements.

Toluene

Figure 11. Variation in D(200) with environmental composition for a soda-lime glass at room temperature. Diamond-studded bits. Note

that solutions of octyl alcohol (Nc = S) in

pentyl alcohol (^( =5) can reproduce the peak

in drilling efficiency which occurs in heptyl

alcohol (iVr = 7). [12].

mond-studded bits and experimental procedures described in 2.2.a.

The variation of D(200) with test environment for the soda lime glass is shown in figure

11. For this material, water and toluene are

relatively poor drilling environments, all of

the alkanes being superior. Note particularly

that n-alcohols having xVc>5 are extremely

effective drilling environments, and that

D(200) for glass in heptyl alcohol is some

twenty times greater than that for water.

Again it was observed that, by mixing two

n-alcohols of different chain length, it is possible to reproduce the drilling characteristics

of the n-alcohols of intermediate chain length

in the homologous series, figure 11(b).

To examine the relevance of the reduction in

surface energy hypothesis to these results, ex- periments have been conducted to determine

the energy, required to propagate pre-existing cracks in this glass as a function of the

same test environments. The double-cantilever

technique was used [47-9], and the data from

these tests, as well as the water content of the

environments (determined by the Karl Fischer

technique) are presented in figures 12 and 13.

The values of </> obtained for this soda-lime

glass in water, 2.2±0.3J/m^ and in the n- alkane series from Nc = 6 to Nc = 16, 4.2 ±0.4-

J/m~, figure 12, are essentially identical to

those found by other workers for similar en439-921 0 - 72 - 11 149
vironments [50-1]. Note particularly, however, that cj) does not vary significantly v^ith Nc

for the alkane environments. For the n-alcohols,

values of 4> intermediate betw^een that for

w^ater and those for the alkanes w^ere obtained,

and consideration of the data in figures 12 and

13 suggests that these values reflect the higher

concentration of v^^ater present in the alcohols.

The value of <^ determined for this glass in

relatively waterfree toluene was 4.0J/m^.

Now, if the data in figures 11-13 are compared, it becomes clear that no sensible relationship exists between drilling efficiency and

either the fracture energy of the glass in these

testing environments, or the water content of

the environment. If, on the other hand, the

drilling behavior of this glass is indeed related

to adsorption-induced variations in its nearsurface flow characteristics, then one should be

able to demonstrate that its hardness. H, also

is environment-sensitive. This possibility was

examined using a single fulcrum (tungsten

carbide) pendulum sclerometer [10]. The relative magnitude of H is evaluated with this instrument from measurements of the damping

behavior of the loaded, oscillating pendulum.

When a hard and brittle solid is tested, relatively little energy is absorbed in near-surface

flow or fracture processes, and so the rate of

pendulum damping is small. For a softer and

more ductile solid, however, relatively more

energy is absorbed by such processes, and so the

rate of damping is greater. In this work, the

relationship between H and pendulum damping

behavior derived by Rebinder (see Ref. [10]

was used to estimate the relative values of H,

and the results are presented in figure 14.

Significant variations in the magnitude of H

occur with Nc. particularly for the n-alcohol

environments, and it is evident that this sodalime glass is markedly harder in heptyl alcohol

(Nc = 7) than in any of the other alcohols or alkanes tested. Note that the variation in H

with environment is not related to the water

16-

co 14

; 12

10

• -n-Alcohols A -n- Alkanes

Water or Toluene

(a)

Ov/oOctyl Alcohol '°°

in Pentyl Alcohol

* »

I L

Water 2

Toluene 6 8 Nc

10 12 14 16

Figure 14. Influence of various environments on the

pendulum hardness, H, of a soda-lime glass at 25 °C. Note that variation in H with Nc is similar

to that of D(200) in figure 11. Compare also

figures 11(b) and 14 (b). [12].

E

E 1.5 O

o

0.5

Sodo-Lime Gloss

In n-Alkanes

10Nc

12 14 16

Figure 15. Influence of n-alkane

environments on drilling behavior of a soda-lime glass us- ing carbide spade bits at 25 °C.

Note opposite variation with No

to that obtained with diamondstudded bits, fig. 11. [2].

content of the environment, cf. figures 13 and

14, but is identical in form to the variation in

D(200), figure 11. Thus, environments which

harden glass (relative to some standard value,

e.g., its hardness in surface-inactive toluene)

facilitate its drilling by diamond-studded bits.

Based on experience, therefore, such environments should be deterimental to the drilling of

glass by carbide spade bits, a prediction verified by the data shown in figure 15.

Data from tests conducted in solutions of

octyl alcohol in pentyl alcohol are also included

in figure 14. Note particularly the existence of a maximum in H for such solutions, the value

of H,„ax being similar to the value of H for an

heptyl alcohol environment.

Only a limited amount of work has been

conducted on other glasses *, and some of the

data obtained are shown in figures 16 and

17 [2]. It can be seen from figure 16 that the

drilling efficiency of each of these materials is

greatest in heptyl alcohol, though the peak is

least pronounced for fused silica. For n-alkane

environments, on the other hand, the data indicate that the principal peak in drilling efficiency occurs in undecane, A^c = ll, figure 17. A similar peak occurs when soda-lime glass is drilled

in the methyl ester homologous series of environments. In this case, octyl acetate is the most

effective cutting medium [2]

3.3 Comments on the Environment-Sensitive

Drilling of Glasses

Consideration of the data presented in fig- ures 11 and 15-17 leads the authors to conclude

that the efficiency of drilling various oxide-type

glasses in alcohol or alkane environments is in

no way related to the water content of these

environments. Nor is it likely that the energy

* CER-VIT (Registered trademark, Owens-Illinois, Inc.) is a

polycrystalline glass ceramic material, produced by controlled

crystallization and used in the manufacturer of telescope lenses [52].

150
Figure 16. Effects of n-alcohol environments on the efficiency of drilling Pyrex, fused

silica, and the glass ceramic material Cervit at 25 °C. Diamond-studded bits. [2].

Note characteristic peak at Nc = l also

found with soda-lime glass, cf, fig. 11 [12].

• Cer-Vit

L X J. I—I 1 I 1 I I I / Water 2 4 6 8 10 12 14 16

Figure 17. Effects of n-alkane environments on

efficiency of drilling Pyrex, fused silica, and

Cer-vit with diamond-studded bits at 25 °C.

Load 300g. Note characteristic peak at Nc = 11

also found with soda-lime glass, cf. fig. 11. [2].

to propagate cracks in a brittle manner (i.e.,

the surface energy) is an important factor in

the drilling process. It appears, therefore, that

neither of the two mechanisms customarily relied upon to explain the environment-sensitive

fracture behavior of glass is relevant to the re- sults described in 3.2, and so other possibilities,

including the flov^-dependent mechanism, must

be considered.

It is of interest, then, to determine whether

any of the more obvious environmental variables might be the factor determining the magnitude of D(200). Possibilities include the dielectric constant (^) [53,541, dipole moment

(p), interfacial energy (ysi), heat of adsorption

(Qa), or wetting (Qw) [39], viscosity (ry), or

coefficient of friction (fiF)- Considerations of

elementary theory or previous work reveal,

however, that for the alcohols or alkanes none

of these parameters varies with Nc in such a manner that one might reasonably expect a

maximum in D(200) at any particular value of

Nc such as is evident in the data presented. For

example, as Nc increases, ysi, e and fxp decrease

monotonically, Qa and rj increase monotonically,

and V does not vary significantly. The further

possibility that variations in D(200) are re- lated to differences in the effectiveness of these

environments as coolants can probably be eliminated in view of (i) the correlation between

the data from pendulum sclerometer and drilling experiments, cf. figures 11 and 14, and fig- ures 11(b) and 14(b) —thermal dissipation effects are unlikely to be important in the former

case, and (ii) the influence of mixed alcohols

on D(200), figure 11(b)—it is difficult to see how mixing two "inferior" cutting environments produces a "superior" one if coolant ac- tion is the factor determining drilling efficiency.

It remains, therefore, to consider further the

hypothesis that the observed variations in

D(200) are a consequence of environmentallyinduced variations in nearsurface deformation

behavior. This possibility appears reasonable

because: (i) previous work has established that

adsorbed species can influence the anelastic behavior of glass when tested in bending [55],

(ii) the observed variations in D(200) are reminiscent of the environment-sensitive variations in drilling behavior of MgO which, as shown earlier, have a demonstrable relationship with adsorption-induced changes in near

surface flow and fracture behavior, and (iii)

the hardness of a solid is a nearsurface or bulk property rather than a true surface property—^yet figure 14 shows that this property, as manifested by pendulum sclerometer measurements, can be significantly influenced by the

test environment.

The problem then resolves into determining

the mechanism by means of which the hardness

(i.e., flow or stress-relaxing behavior) of the

near-surface regions of a glass can be influenced by an environment such that drilling behavior is significantly affected, but </> is not.

Now it may be assumed that <t> for soda-lime

glasses is relatively insensitive to sodium ion

concentration, since the values of cf) for fused

silica and a soda-lime glass containing 14 percent NajO are essentially identical at 77 °K,

namely 4.5±0.2//m- [49]. On the other

hand, soda-lime glass is considerably softer

than fused silica (typically by about 40%).

151
Thus if, for some reason, the concentration or

distribution of sodium ions in the glass surface

becomes different from that in the bulk, the

hardness of the surface regions also will be

different. It is also known that the capacity of

the alcohols and alkanes to act as electron

donors varies with molecular chain length [56]

,

thus the surface charge resulting from their

chemisorption will vary—in some as yet unpredictable manner—with Nr. Then, if the surface

charge influences the local sodium ion concentration, as postulated by Volkenstein [21] and

others, H also will vary withN^o, and so will

drilling efficiency.

The observation that binary solutions of alcohols or alkanes produce drilling efficiencies

equivalent to those of pure, intermediate alcohols or alkanes appears to be consistent with

this hypothesis in that such environments

might be expected to produce, via adsorption,

surface charges intermediate between those of

the two participating species, the specific

charge depending on the relative surface con- centrations of these species.

4. Conclusion

In our view, the potentially most important

application of the Rebinder effect is facilitation

of the machining of hard ceramics, glasses, and

minerals, and the intent of this paper has been

to demonstrate that, for such nonmetallic

solids, an important correlation exists between

the influence of an active environment on near

surface flow behavior and on ease of drilling.

Although the specific mechanisms involved re- main poorly understood, it appears that many

experimental observations can be interpreted

conceptually in terms of the influence of adsorbed species on the distribution or state of

ionization of point defects and impurity ions,

and hence on near-surface dislocation or flow

behavior. These are, however, complex phenomena, and much sophisticated research will be

required before it will be possible to predictably

formulate a cutting environment capable of,

for example, increasing the conventional material removal rates of a specific solid by, say,

ten fold. The need for further studies on polycrystalline ceramics is especially urgent, since

very little information is availalale on their behavior. It has also become clear that one must

always think in terms of the total cutting

system, environment-solid-tool, in order to appreciate the nature of environment-sensitive

machining behavior. Given sufficient effort,

however, it should be possible for Rebinder

effects to finally assume their rightful place as a conventional approach to machining nonmetals

in time for the golden anniversary of their discovery.

The sponsorship of the U. S. Office of Naval

Research, (currently under Contract N00014-

70-C-0330, NR-032-524) during both the preparation of this review and the performance of

much of the work described, is much appreciated. It is also a pleasure to acknowledge the

capable and enthusiastic assistance of R. D.

Huntington in the recent RIAS studies of en- vironmental effects in machining.

5. References

[1] Rebinder, P. A., Schreiner, L. A., and Zhigach,

K. F., Hardness Reducers in Rock Drilling

(C.S.I.R.O., Melbourne, Australia, 1948).

[2] Westwood, A. R. C, Huntington, R. D., and

Latanision, R. M., RIAS, Baltimore, Md., unpublished work (1970).

[3] Robinson, L. H., Effect of hardness reducers on

failure characteristics of rock, Soc. Petrol. Eng.

J., 294-300 (Sent. 1967).

[4] Boozer, G. D., Hiller, K. H., and Serdengecti, S., Effect of pore fluids on the deformation behavior of rock subject to triaxial compression

(Fifth Symp. on Rock Mechanics, U. of Minnesota, May 1962).

[5] Rebinder, P. A., Effect of surface energy changes

on cohesion, hardness, and other properties of

crystals, Proc. Sixth Physics Conf., p. 29 (State

Press, Moscow, USSR, 1928).

[6] Engelhardt, W. von, Versuche zur Verinderung

der Bohrarbeit in Harten Ouarzgesteinen, Oel

und Kohle, 39, 707-709 (1943).

[7] Engelmann, W. H., Terichow, 0., and Selim,

A. A., Zeta potential and pendulum sclerometer studies of granite in a solution environment, RI 7048, USBM, 16pp (1967).

[8] Westwood, A. R. C, Effects of adsorption on hardness and the mobility of near-surface dislocations in nonmetals, Adv. in Materials Res.,

vol. 2, Microplasticity, pp. 365-382 (Interscience Publishers, Inc., New York, 1968).

[9] Westbrook, J. H., Surface effects on the mechanical properties of non-metals, Surfaces and Interfaces, vol. II, Physical and Mechanical Properties, pp. 95-138 (Syracuse University

Press, Syracuse, N. Y., 1968).

[10] Kuznetsov, V. D., Surface Energy of Solids, pp. 45 and 74-110, (Trans. H.M.S.O., London,

1957).

[11] Westwood, A. R. C, and Goldheim, D. L., Occurrence and mechanism of Rebinder effects in

CaF2, J. Appl. Phys. 39. No. 7, 3401-3405

(1968).

[12] Westwood, A. R. C, Parr, G. H., and Latanision,

R. M., Adsorption-sensitive mechanical and

machining behavior of a soda-lime glass (Proc.

Third Intl. Conf. Physics of Noncrystalline

Solids, Sheffield, U. K., Sept. 1970) to be

published.

[13] Rebinder, P. A., Institute for Physical Chemistry,

Academy of Sciences, Moscow; private communication, (January 1969).

[14] Shockey, D. A., and Groves, G. W., Origin of

water-induced toughening in MgO crystals, J. Am. Ceram. Soc. 52, No. 2, 82-85 (1969).

[15] Westwood, A. R. C, Goldheim, D. L., and Lye,

R. G., Rebinder effects in MgO, Phil. Mag. 16,

No. 141, 505-519 (1967).

[16] Westwood, A. R. C, Goldheim, D. L., and Lye,

R. G., Further observations on Rebinder effects in MgO, ibid, 17, No. 149, 951-959 (1968).

152
[17] Johnston, W. G., Effect of impurities on the

flow stress of LiF crystals, J. Appl. Phys. 33,

No. 6, 2050-2058 (1962).

[18] Pratt, P. L., Chang, R., and Newey, C. H., Eff"ect of divalent metal impurity distribution, quenching rate and annealing temperature on flow

stress of ionic crystals, Appl. Phys. Letters 3,

No. 5, 83-85 (1963).

[19] Pratt, P. L., Harrison, R. L., and Newey, C. H.,

Dislocation mobility in ionic crystals. Disc.

Faraday Soc, No. 38, 211-217 (1964).

[20] Volkenstein, F. F., Electronic Theory of Catalysis on Semiconductors (Pergamon Press,

Oxford, U. K., 1963).

[21] Volkenstein, F. F., Electronic processes at the

surface of a semiconductor during adsorption,

Sov. Phys. Uspekhi 9, No. 5, 743-751 (1967).

[22] Westbrook, J. H., and Jorgensen, P. J., Indentation creep of solids. Trans. AIME 223, No. 2, 425-428 (1965).

[23] Hanneman, R. E., and Westbrook, J. H., Effects

of adsorption on the indentation deformation

of nonmetallic solids, Phil. Mag. 18, No. 151,

73-88 (1968).

[24] Seltzer, M. S., Influence of charged defects on

mechanical properties of PbS, J. Appl. Phys.

37, No. 13, 4780-4784 (1966).

[25] Westwood, A. R. C, and Goldheim, D. L., Environmental control of drilling in MgO and

CaFi monocrystals, J. Am. Ceram. Soc. 53, No.

3, 142-147 (1970).

[26] Selim, A. A., Schulz, C. W., and Strebig, K. C,

Effect of additives on impregnated diamond

performance, Soc. Petrol, Eng. J., Preprint No.

2387 (January 1969).

[27] Clarke, F. J. P., and Sambell, R. A. J., Microcracks and their relation to flow and fracture

in single crystals of MgO, Phil. Mag. 5, No. 49,

697-707 (1960).

[28] Clarke, F. J. P., Sambell, R. J., and Tattersall,

H. G., Mechanisms of microcrack growth in

MgO, ibid, 7, No. 75, 393-413 (1962).

[29] Stokes, R. J., Johnston, T. L., and Li, C. H.,

Effect of slip distribution on the fracture behavior of MgO single crystals, ibid, 6, No.

61, 9-24 (1961).

[30] Aust, K. T., Mechanical effects at grain boundaries, Surfaces and Interfaces, vol. II, Physical and Mechanical Properties, pp. 235-269

(Syracuse University Press, Syracuse, N. Y.,

1968).

[31] Westbrook, J. H., and Jorgensen, P. J., Effects of

adsorbed water on indentation hardness anisotrophy in crystals, Antisotropy in Single

Crystal Refractory Compounds, vol. 2, pp. 353-

360 (Plenum Press, Inc., New York, 1968).

[32] Orowan, E., The fatigue of glass under stress,

Nature 154, 341 (1944).

[33] Charles, R. J., and Hillig, W. B., The kinetics

of glass failure by stress corrosion, pp. 511-

527, Symp. on Mechanical Strength of Glass

and Ways of Improving It (Union Scientifique du Verre, Charleroi, Belgium, (1962).

[34] Moorthy, V. K., and Tooley, F. V., Effect of cer- tain organic liquids on strength of glass, J. Am. Ceram. Soc. 39, No. 6, 215-217 (1956).

[35] Culf, C. J., Fracture of glass under various

liquids and gases, J. Soc. Glass Tech. 41, 157-

167T (1957).

[36] Mould, R. E., Strength and static fatigue of

abraded glass under controlled ambient conditions; IV-Effect of surrounding medium, J. Am. Ceram. Soc. 44, No. 10, 481-491 (1961).

[37] Engelhardt, W. von, Schleiffesstigkeit und

Grenzflachenenergie fester Stoffe, Naturwiss

33, No. 7, 195-203 (1946).

[38] Ramsauer, R., Uber den Einflub von Fliisigkeiten

auf die mechanische oberflachenvergroberung von Glas, Glastechnische Berichte 24, No. 10,

239-247 (1951).

[39] Weyl, W. A., Surface structure and surface

properties of crystals and glasses, J. Am.

Ceram. Soc. 32. No. 11, 367-374 (1949).

[40] Marsh, D. M., Plastic flow in glass, Proc. Roy.

Soc. A279, 420-435 (1964).

[41] Dick, E., Neue Versuche Zur Mikroplastizat von

Glas, Glastechnische Berichte 43, No. 1, 16-21

(1970).

[42] Lange, F. F., Non elastic deformation during

fracture of glass. Tech. Rept. No. 3 on Contract

NOOO14-68-C-0323, NR-032-507. Westinghouse Elec. Co., R. & D. Center, Pittsburgh,

Pa. (1970).

[43] Weber, R., and Goldstein, M., Stress-induced migration and partial molar volume of sodium

ions in glass, J. Chem. Phys. 41, No. 9, 2898-

2901 (1964).

[44] Weyl, W. A., The rheology of inorganic glasses,

Rheology, vol. 3, pp. 299-340 (Academic Press,

New York, 1960).

[45] G. Quinke, Ann. d. Physik 10, 160, 374, 512

(1880). See Weyl, Ref. [44], p. 339.

[46] Westwood, A. R. C, Latanision, R. M., and Lye,

R. G., Adsorption-sensitive anelastic effects in

glass, Phys. Stat. Solidi, (a) 3, K17 (1970).

[47] Gilman, J.J., Direct measurements of the surface

energies of crystals, J. Appl. Phys. 31, No. 12,

2208-2218 (1969).

[48] Weiderhorn, S. M., Fracture surface energy of

soda-lime glass. Mat. Sci. Res., vol. 3, pp. 503-

528 (Plenum Press, New York, 1966).

[49] Weiderhorn, S. M., Fracture surface energy of

glass, J. Am. Ceram. Soc. 52, No. 2, 99-105

(1969).

[50] Linger, K. R., and Holloway, D. G., The fracture

energy of glass, Phil. Mag. 18, No. 156, 1269-

1280 (1968).

[51] Weiderhorn, S. M., and Johnson, H., Failure of

glass under biaxial loading, N.B.S. Rept. No.

10145 on Contract NAonr-7-70, NR-032-517.

[52] Stewart, D. R., and Davis, W. L., Evaluation of a large glass-ceramic mirror blank, Appl. Optics 9, No. 4, 938-941 (1970).

[53] Berdennikov, W. P., Measurement of surface ten- sion of solids, Zhur Fiz. Khim. 5, No. 2-3,

358-371 (1934).

[54] Cox, S. M., Glass strength and ion mobility, Phys.

and Chem. of Glasses 10, No. 6, 226-239 (1969).

[55] Aslanova, M. S., and Rebinder, P. A., Effects of

adsorption on elasticity and creep of glass

fibers, Dokl. Akad. Nauk 156, No. 2, 299-302

(1954).

[56] Ingold, C. K., Structure and Mechanism in Organic Chemistry (Cornell Univ. Press, Ithaca,

N. Y., 1953).

Discussion

DiNESS: What was the concentration of sodium ion in the fused silica?

Westwood: I don't know. It must be quite

low, however.

DiNESS: Do you think that your rationalization can account for this effect in such a low

concentration of sodium?

Westwood: That is a good question, Art,

153
and I can't give you a straightforward answer.

All commercial e-lasses have some sodium present, however, and our calculations indicate that

increases in sodium ion concentration of only

10 pnm in a surface layer —- 0.1 ixm thick are sufficient to relax quite large tensile

stresses* and hence, presumably, to influence

the crack initiation step.

DiNESS : One further point, have you looked

at any non-oxidic glasses?

Westwood: No.

WiEDERHORN : Regarding the sodium concentration in fused silica, I believe the range to be

of the order of 10 parts per million and less.

You could also get rather large hydroxyl concentrations. I don't know how this would affect

your results.

Westwood: Yes, that's a good point. In

fact, while one might expect a negative surface

charge to attract sodium ions to the surface, a

positive surface charge probably would attract

hydroxyl ions instead. The presence of either

species in the surface layers is likely to lead to

some reduction in hardness.

WiEDERHORN : Let me ask you another question regarding hardness. Did you try the con- ventional pyramidal type of indenter?

Westwood: Yes, but it proved very difficult

to find the indentations made under the test en- vironment with indenter loads of only 1 —10 g

after the environment was removed. Accordingly, our hardness determinations were made

with the pendulum device.

Rice: Just briefly, you have addressed yourself mainly to glasses and ionically bonded materials. Do you expect somewhat similar effects

in the hard materials such as carbides, borides,

and nitrides which are more often characterized by covalent bonding? In these materials

lattice friction is probably much more important.

Westwood: Yes, there is considerable evidence for Rebinder effects in solids which have

(Westwood, A. R. C, Latanision, R. M., and Lye, R. G., Phys.

Stat. SoHdi (a) 3. KIT (1970). **(For a review of such effects see Westwood, A. R. C, in "Microplasticity", Interscience Publishers, New York, pp. 365-82

(1968).

some degree of covalent bonding**. One has to

be careful when talking about carbides though

because they have a significant amount of metallic bonding. However. Jack Westbrook has

shown such effects in TiC. One further point:

Rebinder effects are sensitive to both light and

imposed potentials—two other variables that

may prove exploitable in technology.

Heuer: I'd like to ask, since there is a lot

of interest in the machining of alumina,

whether you intend to do any work on alumina.

Westwood : Yes, we plan to. KiRCHNER ; When your report came out about

a year ago, we became interested in the effect

of these media on the strength of alumina. We,

however, found essentially no effect on the

strength. Now, do you have any comments on

the relationship of this and the role of microplasticity in the failure of alumina?

Westwood: Was this polycrystalline alumina?

KiRCHNER: Yes.

Westwood: And you measured tensile

strength ? KiRCHNER: Well, we measured flexure

strength.

Westwood: Of course, when one is dealing

with polycrystalline soli(Js, there are a lot of

new factors involved in the problem. Imagine

that we had many, many grains throughout

the solid, but only the surface layer was affected by the adsorbing species. Imagine also that

the actual mechanical properties were dominated by the presence of voids, cracks, and

other sort of defects within the solid. In this

case, one would not expect to see any effect

of environment on the tensile strength. On the

other hand, if one were to study, say, the bending fatigue behavior of such a solid, in which

case the deformation process is initiated and

concentrated near the surface, then this property may well be dominated by the environment, and one might expect to see effects. Machining is, of course, a near-surface dominated

process—that's why it is possible to see such

spectacular effects.

national BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

154
B. Nonmechanical Methods

Shaping or Figuring Ceramic Surfaces by Ion-Beam Bombardment

Paul W. Levy

Brookhaven National Laboratory, Upton, New York 11973

The shaping or figTiring of glass and ceramic surfaces by ion-beam erosion is a

relatively new technique. All operations are carried out in a vacuum. Usually the

optic is bombarded by a collimated beam of rare gas ions whose energy is between

1-100 keV. Selected areas can be eroded by electrostatically deflecting the beam or, less conveniently, by moving the optic. The erosion area can be controlled by focusing

the beam with electrostatic or magnetic lenses. The basis of the erosion process is the physical phenomenon known as sputtering. This occurs when the incident ions scatter elastically with atoms near the surface of the target and some of the displaced atoms, after participating in one or more collisions, acquire sufficient velocity

to escape from the solid. Other dislodged atoms become interstitials and, together with

vacancies created in the recoil process, create a radiation-damaged layer just below

the eroding surface, whose thickness is roughly the range of the bombarding ions.

For ceramic materials typical erosion rates are on the order of angstroms/^tamp/

cm'/min and increase with increasing angle of incidence. Ion-beam currents up to 100 /xamp/cm- are practical. Under these conditions the material struck by the beam

may reach temperatures between 300 and 700 °C. The maximum lattice disruption occurs at a depth of one incident ion range from the original surface, just as the

eroding surface reaches this point. Additional bombardment does not increase the

damage beyond this maximum. Furthermore, the damage is minimized by using

bombardment conditions which maximize the temperature in the target. It is not

known if there are any radiation damage effects which reduce the usefulness of

ion-beam figured optics. At the moment, it would appear that ion-beam figuring

is a practical technique for "touching up" conventional optics and, more importantly,

for fabricating non-spherical or non-axially symmetrical optics. Also, the ion-beam

erosion process is particularly amenable to computer control.

Key words: Figuring; ion-beams; optical surfaces; polishing; radiation damage;

sputtering.

1 . Introduction

For approximately two hundred years ce- ramic surfaces, particularly glasses, have been

shaped and polished by methods employing

abrasive substances. During the last decade or

tw^o a number of nev^ shaping or figuring techniques have been introduced. One of the most

interesting, and promising from a technological point of view, is "ion-beam sputtering" or "ion-beam erosion." Very briefly, this process

is based on the physical phenomenon knov^^n as "sputtering," v^^hich occurs when metal, ce- ramic, or even plastic surfaces are bombarded

by energetic ions and, incidentally, even electrons. When a fast moving ion enters the

surface of a solid the interaction is sufficiently

violent to eject one or more atoms from their

normal positions in the crystal or glass lattice.

Some of these atoms will be scattered towards

the surface and will escape from the solid. Other

atoms will penetrate into the solid to produce

"radiation damage effects." Still other atoms

will make one or more collisions and eventually come to rest in a normal lattice position.

'Supported in part by Advanced Research Projects Agency

(ARPA) Contract with Kollsman Instrument Corp. and in part

by the U.S. Atomic Energy Commission.

^Figures in brackets indicate the literature references at the end of the paper.

As the bombardment continues the surface is

gradually eroded and a layer of disrupted or radiation-damaged material develops just below the slowly eroding surface.

The literature on ion-beam erosion or sputtering is, with very few exceptions, confined to

papers which deal with the physics of this process. However, it is closely related to radiation

damage processes, and a considerable amount

of useful information can be obtained from

publications on this subject. A good example

in this category is the book by Carter and Colligan [1].^ In fact, this book contains background information on almost every subject

discussed in this article. A useful compendium

of papers on sputtering physics is contained in

the "Proceedings of an International Conference on Atomic Collisions and Penetration Studies with Energetic (keV) Ion-Beams" [2].

The most comprehensive source of information

relating directly to ion-beam figuring, which

appears to be available at this time, is a Kollsman Instrument Company Report, "Ion-Beam

Optical Figuring" [3] . More than a few publications could be listed which describe the application of electron- and ion-beam bombardment to cleaning, deoxidizing, degassing, etc.,

of metal and non-metal surfaces for a variety

155
of purposes. There are indications in a large

fraction of these papers that sputtering, i.e.,

bombardment-induced erosion, is the process

actually producing the observed effects. Since

they have only peripheral applicability, they

have not been included. Apparently there are

only two papers [4, 5] which are concerned directly with the application of ion-beam bombardment to the figuring of optics.

The next section will contain a description of

the basic elements in an ion-beam figuring system. The following section will describe the

basic physical processes occuring during sputtering. This will include some practical information on the rate of material removal as a

function of ion energy, angle of incidence, etc. During this discussion, it will become apparent

that the material immediately below the eroded

surface is in a radiation-damaged condition.

Consequently, a separate section will be devoted

to radiation damage considerations. This will

lead to some conclusions, or recommendations,

for figuring pieces in a way that minimizes ra- diation damage effects. Finally, the last section

will be a brief discussion on the practical applications of the ion-beam figuring technique.

2. The Basic Ion-Beam Figuring

Process

The basic elements that are contained in a

typical ion-beam figuring arrangement are shown schematically in figure 1. First, all operations are performed in a vacuum chamber.

The piece to be worked or figured is securely

fastened to a holding device which is usually

mounted on the center line of the ion beam.

The bombarding ions are first ionized and then

accelerated by an electrostatic accelerator. The

TO VACUUM PUMP

Figure 1. The basic elements in an ionbeam figuring system.

position of the beam on the working piece is

usually controlled by electrostatic and/or magnetic lenses. The surface configuration changes

occuring during bombardment can be determined with an interferometer by removing the

piece from the chamber at desired intervals or monitoring during irradiation. Some additional

information on each of these basic elements of

the system is described below.

Typically, ion-beams are produced by low

energy electrostatic accelerators. For figuring

purposes the energy is usually in the range of

1 to 100 keV. The accelerator contains sections

for producing the ions by ionizing rare gas

atoms, accelerating them to the desired velocity or energy, and focusing them into a relatively

narrow beam which can be focused on the sur- face of the target. Almost always, the bombarding beam consists of rare gas ions since they

possess several advantages over other atoms or

molecules. Rare gases are chemically inert and

do not react with the working piece or exposed

parts of the vacuum chamber. Also, they consist of monoatomic atoms which can be singly

ionized without producing a troublesome number of doubly ionized atoms. Furthermore, they

are convenient to handle and, with the exception of helium, are relatively easy to pump.

Two methods are commonly employed to fix

the position where material is removed from

the surface of the working piece. The first

method employs a beam maintained in a fixed

position with the piece moved in a controlled

fashion by the holding device. Alternatively,

the piece is held in a fixed position and the

beam directed to the point on the surface where

material is to be removed. Actually, for both

methods it is desirable that the apparatus include a device for directing, or steering, the ion

beam. Either electrostatic or magnetic deflecting arrangements can be used. Electrostatic deflection is the simpler but magnetic deflection

may provide more precise control.

If the ion-beam is maintained in a fixed position and erosion control obtained by moving the

working piece, the holding device must include

several important features. First, it must be

able to control precisely the position of the

piece or, alternatively, it must be capable of

rotating the piece about a precisely controlled

axis. Furthermore, it is essential that vibrations

be reduced to an absolute minimum and it is

highly desirable to be able to control these

movements from outside of the sputtering

chamber.

Clearly, it is advantageous to follow the progress of the figuring during bombardment.

There are a number of different interferometer

arrangements that can be used for this purpose.

The fringe patterns appearing in an interferometer can be observed visually, or with a

"television camera" and displayed on a televi156
sion monitor, and they can be recorded photographically.

Since it is possible to follow the progress of

the figuring during bombardment, it is desirable to be able to control where the erosion is

taking place. As mentioned above, this can be

accomplished by maintaining the beam in a

fixed position and mechanically moving the

piece about. However, it would appear that

there are distinct advantages in keeping the

piece fixed and controlling the erosion process

by electrostatically or magnetically deflecting

the beam. This arrangement reduces the complexity but does not eliminate the mechanical

target-moving apparatus. It is apparent that

video fringe registration and electrical control

of the erosion process are features which can

be combined in a closed-loop computer controlled figuring system. In fact, a study completed in 1968 indicates that a completely automatic computer control system is feasible [3].

Also, systems combining manual and computer

control would appear to be particularly desirable. It would appear, at least to the writer,

that one of the major diflRculties which will be

encountered in an automatic system is the distortion of the surface occurring at the erosion

point due to heat deposited by the ion-beam.

This distortion disappears as soon as the beam

is removed, and, in principle, it should be possible to include a time-dependent correction for

the distortion in the computer control program.

Inasmuch as the ion source continuously

emits atoms, the system must be continually

pumped. For efficient erosion the vacuum re- quirements are not great. The pressure must be

maintained at a level where collisions between

the residual gas atoms and the ions in the beam

are negligible.

One additional feature, not included in figure

1, should be described. This is the device for

preventing charge build-up on the target piece.

During bombardment the piece becomes positively charged, first because the rare gas ions are positively charged, and secondly because of

electron emission from the surface during bombardment. This positive charge build-up can be

prevented by including an electron gun in the

chamber which continuously directs negatively

charged electrons on the target.

3. The Basic Atomic Sputtering

Process

3.1. Interactions Involving Energetic

Ions and Solids

Almost all of the physical processes which

occur when energetic ions interact with a solid are illustrated in figure 2. The energy of the

incident ion will be quite high as it enters the

target. It will interact with the atoms in the

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(D INCOMING ION CONTINUES INTO SAMPLE

0 ATOM RECOIL EJECTS

Si ATOM FROM LATTICE FORMING Si VACANCY

Figure 2. The collision processes occurring during

a typical sputtering event.

solid in two distinct ways. First, both the incident particle and the stationary atoms may be

considered as hard "billiard balls" that collide or scatter elastically. Each collision will reduce

the energy or velocity of the incident particle

and it will gradually slow down. The slowing

down rate depends on the energy and mass of

the incident particle and the mass of the atoms

in the solid. This particular process is characterized by the term, "nuclear stopping power."

If the stopping power is large, the incident ion

quickly loses energy. Conversely, if the nuclear

stopping power is low, the ion will lose energy

slowly and penetrate relatively deeply into the

solid. A second important energy loss mechanism results from the fact that the incident ion

is charged and usually remains charged until

its velocity is quite small. Specifically, there is

an electrostatic interaction between the charged

ion and the electrons and shielded nuclear

charge of the atoms in the solid which causes

the incident ion to lose energy. This second

energy loss mechanism is characterized by the

term, "electronic stopping power." Actually, the

electronic stopping power is roughly proportional to the velocity of the ion or, i.e., to the

square root of its energy.

If, in an elastic collision, suflficient energy is

transferred to a lattice atom, it will be ejected

from its normal position. The exact energy that

must be imparted to any atom to cause this to occur depends on several factors. These include:

the strength of the binding forces acting on the

lattice atoms, the masses of the atoms involved,

the geometrical arrangement of the surrounding atoms, etc. Very roughly, an average of 25

eV must be imparted to a typical glass atom

to cause it to be ejected from the lattice.

All of the collision processes violently excite

the atoms along the trajectoiy of the incident

ion and the trajectory of any displaced atoms.

157
This local excitation can be viewed as a rapid

and large increase in temperature and is often

referred to as a "thermal spike." The energy

in a typical thermal spike is dissipated in nanoseconds or even shorter times. Some of the

properties associated with thermal spikes will

be discussed in more detail below.

The slowing down processes described above

limit the distance that an energetic incident

ion can penetrate into a solid. All particles having a specific energy E will penetrate very

nearly the same distance into the solid. However, the trajectory followed by one ion will

differ from that of all other ions. Intuitively,

one can readily understand that the sequence

of scattering events will differ for each particle. The average distance penetrated by a beam

of ions, incident normal to the surface and having a specific well-defined energy E, is defined

as the range. The statistical variation in penetration deDth, about the range, is called the

"straggle." More precisely, the straggle is defined as the mean square variation in the penetration distance and is usually on the order of

1 to 5 percent of the range. Both of these concepts are suggested, but not precisely described,

in figure 3.

3.2. Conventional Range and Energy

Relations

A very large fraction of the experimental

and theoretical papers that deal with the basic

sputtering process make use of the so-called dimensionless range and energy relations:

= 7?|4^ N M,

= 4-

Mo

In these expressions:

R = Range, p = dimensionless "range"

E = Energj^ e = dimensionless "energy'

N = number of atoms per unit volume

DAMAGE vs DISTANCE

PROFILE Q (x) USED

FOR CALCULATIONS

X, DISTANCE FROM ORIGINAL SURFACE

Figure 3. A typical damage versus penetration distance profile determined experimentally and the

damage versus depth profile used in the calculations in the text.

Mo

Zo

a =

0,0 =

mass of the bombarding ion

atomic number of the bombarding ion

mass of the atoms in the target

atomic number of the atoms in the

target

0.885 ao where

(Zt2/3 -I- Zo2/3)l/2

0.529 X 10-* cm, i.e., the Bohr radius.

The first of these is the range (p for range)

relation and the second the energy (e for en- ergy) relation. These two relations contain a surprising amount of information and in particular indicate how the range depends on the

energy, mass, and atomic number of the incident particle and the mass and atomic number

of the atoms in the target solid. Although these

two relations are useful, one does not have to

understand them in detail to obtain a good

grasp of the basic sputtering process described

in the next section.

3.3. The Basic Sputtering Process

The basic physical processes producing sputtering are illustrated in figure 2. In this diagram the first interaction occurs when the incident particle makes a glancing collision with

an oxygen atom quite near the surface. In turn,

this oxygen recoil makes an additional collision

with a nearby silicon atom and passes out of

the surface, i.e., it escapes from the solid. Although there are relatively few individual

events in this particular sequence, it demonstrates one of the most important features of

the sputtering process, namely the multiple

collision sequence or cascade. Single-event sputtering collisions can occur, especially if the

incident ion enters the surface at a glancing

angle. However, theoretical treatments on sputtering indicate that multiple collision processes

account for a very large fraction of the es- caping atoms. It should be apparent that the

efficiency of the sputtering process will be least

for normally incident ions and should increase

as the angle of incidence is increased. This is

demonstrated by numerous measurements. In

order that a recoiling atom escape from this

solid, it must possess two properties. First, it

must have a velocity component away from the

surface. Second, it must have sufficient energy

to overcome both the nuclear and electronic

energy losses that it will suffer in traversing

the solid and, in some materials, it must have

sufficient additional energy to overcome surface

binding forces. The surface binding forces are between 1 and 100 eV and can usually be neglected. For most practical sputtering situations

the incident ion energy is 20 keV or more and

thus the surface binding effects are completely

negligible. The disrupted lattice which remains

after a typical sputtering cascade, often re158
ferred to as radiation damage, cannot be ignored. Some comments will be made about it immediately; however, detailed radiation damage considerations are contained in the next

main section.

3.4. Sputtering-Produced Lattice Distortion

The sputtering process, outlined above and

illustrated in figure 2, obviously produces considerable lattice distortion. Actually, a large

number of distinctly different processes, all involving the movement of atoms, contribute

to this distortion. Only the more important of

these will be described. First, as is indicated in

figure 2, the incident ion can penetrate into the

lattice and come to rest in an interstitial position. If it happens to locate in a sufficiently large

empty volume, typically found in glass lattices,

it will not distort the surrounding lattice appreciably. If it comes to rest in a tightly packed

region it can produce appreciable distortion.

These comments apply equally well to any lattice atoms which have been ejected from their

normal positions during the sputtering cascade

and do not escape through the surface. In fact,

since the number of lattice atoms ejected by

an incident ion is usually much greater than

one (1), there will be many more lattice-atom

interstitials than incident-ion interstitials. The

other principal contribution to the lattice distortion arises from the vacancies, i.e., the empty

lattice sites which remain after atoms have

been ejected from their normal positions. Incidentally, vacancies created during one sputtering event have a high probability for capturing mobile atoms from later events. It is

easy to understand that interstitials have a tendency to expand the radiation damage part

of the lattice. Surprisingly, under certain circumstances the introduction of vacancies also

has a tendency to expand the lattice. Perhaps

the easiest way to explain this is to point out

that the removal of an atom also removes the

binding forces holding together the atoms on

either side.

3.5. Sputtering Yield and Sputtering

Efficiency

It is essential for anyone interested in the

sputtering process or ion-beam erosion to differentiate clearly between two commonly used

definitions. These are the sputtering yield and

the sputtering efficiency which are defined by

the following expressions:

Sputtering Efficiency =

Sputtering Yield =

number of sputtered

atoms

number of incoming

ions

energy of the

sputtered atoms

energy imparted

by incoming ions

Note that the sputtering yield refers to the

number of sputtered atoms produced by each

entering ion, whereas the sputtering efficiency

is defined in terms of energy. More specifically,

to determine the sputtering efficiency one must

determine, or add up, the energy of all of the

atoms which escape from the target and compare it to the total energy of the incident ions.

The sputtering yield is the more common ex- pression and will be used throughout the re- mainder of this paper. However, both of these

quantities are functions of numerous parameters, e.g., incident ion energy, mass, atomic

number, angle of incidence, and, in particular,

the mass and the atomic number of the atoms

in the target piece. For practical purposes it is convenient to use a quantity called the ero- sion rate, which is usually expressed in angstrom (A), of material removed per square cm

per unit beam current per unit time. Typical

sputtering yield values for fused silica, as a

function of incident ion energy, are shown in

figures 4 and 5. This illustrates three important

features: first, the sputtering yield increases

as the mass of the incident particle increases

second, the yield vs. energy curve has a distinct maximum ; and third, the maximum yield

increases to higher energies with increasing

incident particle mass. The data given in this

figure is for the incident ions striking the sur- face perpendicularly, i.e., for zero angle of

incidence. Above it was implied that the sputtering yield should increase as the angle of

incidence is increased. This is clearly demonstrated by figure 6 which shows how the erosion

20 40 60 80 100 ION ENERGY (keV)

Figure 4. Range of argon and krypton

ions in fused silica.

159
z

40 80 120 ENERGY OF BOMBARDING ION (keV)

Figure 5. Typical sputtering yield versus

incident ion energy curves for fused

silica.

rate varies with angle of incidence for three

common lens materials. The yield versus angle

of incidence curves also depend on the mass

and energy of the incident ion.

3.6. Target-Area Temperature During

Bombardment

During bombardment numerous sputtering

cascades are taking place, and the volume immediately below the area struck by the beam

can attain surprisingly high temperatures. This

is readily understood by pointing out that the

sputtering process converts the incident ion

energy into the motion of atoms in the target

and to electronic excitation. In reality, both

of these energy-transfer processes can be re- garded as increasing the local energy temperature. As mentioned above, the rapid increase

and decrease in temperature associated with a

single cascade event is often called a thermal

spike. When a given area is under bombardment and each incident ion produces an individual spike, the net result is a large temperature increase directly below the surface. The

actual temperature will depend on numerous

parameters, e.g., beam size and energy, thermal

conductivity of the target piece, and the temperature of the holding arrangement which

might act as a heat sink. Consider just one

<

a.

ino

FUSED SI LICA-.,^/ y/ 1 BK-7GLASS\x/ /'

/

"^^CERVIj/

1 1 1 20 40 60

ANGLE OF INCIDENCE 80

Figure 6. Typical sputtering yield versus ion-beam angle of incidence c^irves.

example, the maximum temperature attained

by a typical glass or fused silica optic under

continuous bombardment. A 100-microamp

beam of 100 keV ions will impart 10 watts to

the surface. If the beam strikes an area of one

square centimeter the following approximate

temperatures occur at the center of the beam

and at various points near the surface. On the

surface the temperature is 700 °C, 1 mm below

the surface it is 400 °C, and 6 mm below the

surface it is 150 °C. Incidentally, these values

apply to a piece approximately 10 cm in diameter and 1 cm thick. As soon as the beam is

"shut off" the temperatures drop very rapidly.

In fact, when the surface temperature is approximately 300 °C during bombardment it will drop to between 50 and 100 °C within one- tenth of a second after the bombardment is

terminated. These considerations indicate that

it is possible to erode an optic without raising

its temperature excessively at any point by the

simple expedient of moving the eroding beam

back and forth on the surface. However, as will

be demonstrated in the next section, radiation

damage effects can be minimized by figuring

the piece under conditions which maximize the

temperature near the surface.

3.7. Channeling and "Amorphous

Sputtering"

One of the more interesting radiation effects

discovered during the last decade is the phenomenon known as channeling. It is now known

that the concept of the penetration depth or range introduced above can be applied precisely

only to glassy or amorphous substances. More

specifically, it needs to be modified when applied to most crystalline materials. When crystalline substances are bombarded by particles

which are accurately aligned along certain

crystallographic directions, it is found that a

large fraction of the incident particles penetrate into the crystal considerably beyond the

range defined above. In fact, the particles may

penetrate to a depth of 10, 20, or even 30 times

160
greater than the "range." The reason for this

is now well understood. The atoms in many

crystals are aligned in a way that produces

well defined channels. Once an energetic particle enters such a channel with its velocity

vector closely aligned with the channel axis, it can proceed in an almost uninhibited fashion.

As long as the particle is traversing a part of

the crystal which is free of defects there is

nothing in its path to restrict its progress. Furthermore, even if the velocity vector does not

correspond precisely to the axis of the channel,

the particle is assisted along the channel by

various "focusing effects." Inasmuch as such

channels do not exist in glassy or amorphous

materials, channeling effects are not observed

in these substances.

It is reasonable to expect that channeling

effects will play a role in the ion-beam figuring

of crystalline materials. It might be expected

that they would be particularly important in

radiation damage considerations. However, in

this paper all comments on channeling are confined to this subsection. In crystalline materials

it is convenient to define a range which does

not include channeling effects. Unfortunately,

the range defined in this way is usually referred

to as the amorphous range. Also, the parameters

associated with the sputtering process, e.g., the

sputtering yield or the sputtering efficiency,

are defined in a way which obviates (actually

it ignores) all channeling effects. These parameters are said to apply to the "amorphous"

materials. Thus, in the literature one sees ref- erences to the sputtering parameters for amorphous AloO:,, a substance which, as far as the

writer can determine, does not exist in a glassy

or amorphous form.

4. Radiation Damage in Ion-Beam

Figured Materials

4.1. The Necessity for Radiation Damage

Considerations

The description of the sputtering process,

contained in the previous section, indicates that

the material just below the surface of the target

is in a "radiation-damaged" condition. This

region contains vacancies and interstitials

which will, almost certainly, distort the lattice. It may contain numerous color centers or, i.e.,

may have become colored. The glass technologists would say that it has been browned by

the bombardment. Intuitively one might expect

that the damaged or distorted lattice would be

unstable in one or more respects. A sufficiently

large change in dimension occurring after an

optic had been figured by ion-beam erosion

would render the piece useless. Consider an- other example: for certain applications the index of refraction of the material at its surface

must be stable, e.g., for dielectric coating or

thin-film applications. Thus, even if the material were dimensionally stable, a change in

dielectric constant would be intolerable. Radiation-induced changes in the density, index

of refraction, and the dimensions of quartz

and fused silica samples have been extensively

studied [6]. These and other radiation effects

usually observed in ceramic materials have

been discussed previously [7]. However, and

this point must be made explicitly clear, with

only one or two exceptions, all of the available

radiation effects information on glasses and

other ceramic materials has been obtained

from materials bombarded in a reactor or with

gamma rays or x rays. Perhaps the most

pertinent work is a study of ion-bombardment

induced changes in the index of refraction of

fused silica [8] . However, it is diflftcult to apply

this work to practical figuring applications.

There are studies on the release of trapped

ions implanted during bombardment, the diffusion of implanted ions, etc., that provide useful information but only peripheral to the considerations described above.

4.2. Nature of the Radiation-Damaged

Material

The process leading to the formation of va- cancies and interstitials along the trajectory of

the bombarding ion and recoil atonis has been

amply described. However, the distribution of

vacancies and interstitials along the trajectory

of the incoming ion varies considerably. This is

shown in figure 3 where Qix) is the number

of displaced atoms along the trajectory a distance X from the surface from the target. It

is difficult to approximate the true Q(x) function analytically, and consequently the approximations shown in this figure will be used

throughout the remainder of this paper. Specifically, it will be assumed that Q{x) is a constant for all values of x less than the range

R and is 0 for values of x greater than r. Note

that Q(x) describes the number of displaced

atoms produced by each individual incident ion.

Inasmuch as more than one bombarding ion

may contribute to the damage at any specific

point X, where x<R, the damage may be greater

than Q{x). The quantity D(x,t) will be used

to describe the damage at point x, at time t, after the bombardment has commenced. If the

bombardment has proceeded for a considerable time and the concentration of vacancies

and interstitials has reached an appreciable

level, the vacancies and interstitials produced

late in the bombardment have increased probability for combining with their opposite or "antipodal" defects and the rate of damage

formation will be decreased. In other words,

if the bombarded material contains an appreci161
able concentration of vacancies and interstitials, any vacancies or interstitials produced in

subsequent events can re-enter or rejoin the

normal lattice. Thus the damage produced by

each incident ion vi^ill diminish as the bombardment progresses. This phenomenon is sometimes called "radiation annealing." However,

in the considerations that follow it will be

assumed that the induced damage is linearly

proportional to the number of incident ions.

The annealing of radiation damage by radiation-produced processes should not be confused

with the removal of damage by the thermal

processes which will be discussed in a later

section.

The sputtering process includes an appreciable number of ionization events. That is,

electrons are excited to an extent that they

escape from bound atoms and peregrinate, or wander about, the target material. Also, the

"missing charges" or (electronic) holes are

mobile. When electrons or holes are trapped

by certain entities such as vacancies, impurities, or more complex defects, color centers are formed. As mentioned above, this is referred

to as coloring or browning. For many practical

purposes, e.g., in reflecting optics, the coloring

is unimportant. Scientifically this coloring can

be quite useful. When defects such as oxygen

atom vacancies are converted into color centers

by charge trapping, the color-center concentration can be used as a measure of the damage

concentration. This technique has been described in a previous paper [9] and, as will be

described below, provides information directly

applicable to radiation-damage considerations

in ion-beam figured fused silica.

4.3. The Measurement of Radiation Damage

in Figured Optics

Radiation damage can be characterized by

color-center measurements under some circumstances. However, a really good indicator for

measuring radiation damage in ion-beam eroded

materials does not exist. Furthermore, a method

might be devised which would apply to one

damage characteristic but not another. For

example, an interferometer measurement might

indicate that a figured piece does not contain a detrimental amount of damage. However, an

entirely unsatisfactory amount of damage may

be present when assessed by another method.

For example, a piece which is acceptable in

most respects may be unacceptable when evaluated in terms of its resistance to attack by

water vapor. This discussion has been included

to emphasize the point that it is not possible to

conclude that radiation damage in figured optics can be ignored on the basis of one or two

types of tests.

4.4. Kinetics of Radiation Damage

Formation

The discussion on the kinetics of radiation

damage formation during ion-beam figuring,

contained in this section, will be based on the

assumption that neither radiation annealing nor

thermal annealing processes are important. This

is equivalent to restricting the derivations to

the case in which the ion-beam currents are low enough to prevent the target from being

heated more than a negligible amount. Also,

it is assumed that the total or integrated flux

is low enough to prevent extensive overlap of

individual damage regions. In practice this

latter condition is hard to realize. However, by

adopting it the calculations can be very greatly

simplified without eliminating any important

features of the description of the damage formation process.

4.4.1. Definitions

The following parameters will be used in the

derivation of the equations describing the rate

of radiation damage formation in a target

being eroded by ion-beam bombardment.

t = time, t = 0 corresponds to the start

of the erosion bombardment

X = distance or depth measured from

the coordinate of the surface pri- or to bombardment, i.e., the sur- face is at x = 0 when t = 0.

^sit) = position of the surface at time t, x,it) =0 at t = 0. R = range or penetration depth of the

incident ion.

Q(x) = number of defects (per unit length)

introduced along the trajectory

of each incident particle. It will

be assumed that Q(x) is constant, or more specifically:

Q{x) =Q, if Xs<x<Xs + R, and

Q{x) =0, if x>Xs + R.

D{x,t) = concentration of defects at time t and point x. If the eroding sur- face has passed the point x, i.e.,

if x<Xs, D(x,t) has no meaning.

/ = flux (number of ions per unit time,

per unit area) of bombarding particles striking the target.

S = surface erosion rate. Specifically S

is the depth of material removed,

per unit area, by one incident ion.

Also, S is assumed to be constant,

i.e., independent of the amount

of prior bombardment.

4.4.2. The Basic Erosion Rate Equation

The rate at which the target surface erodes

during sputtering can be described by a sur- prisingly simple expression. First, it is assumed

that the incident ion flux is uniformly distrib162
uted over the target area and that the beam

current is constant. In a time period Ai, the total number of ions striking a unit surface area

is /Ai, and this number of ions will, using the

definitions given above, cause the surface to

erode a distance Aa: = SlAt. If the total bombardment time is t, the surface will move a

distance Xsit) given by

SIdt = Sit

It will become apparent that the simple equation

Xsit) = Sit

is very useful.

4.4.3. Damage Formation at a Depth Greater than the

Incident Ion Range

The description of the rate of radiation damage formation at a point below the surface

of the target will be divided into two parts.

First, in this section the damage formation will

be described for a point Xi which lies farther

from the original surface than the maximum

penetration depth or range of the incident ions.

Second, the damage formation will be described

for a point which lies between the original target surface and the maximum penetration

depth. Note particularly that the position or coordinate of a particular point in the target is

always measured with respect to the position

of the target surface before bombardment. In

other words, the surface of the target is receding during erosion but distances are always

reckoned from the surface coordinate at irradiation time ^ = 0. The geometry of the situation to be considered first is illustrated in figure 7. At the time when the irradiation is

started, the incident ions will penetrate a disX, DISTANCE FROM ORIGINAL

SURFACE

Figure 7. Spatial relations for determining the rate of damage formation at a point which lies at a

distance Xi, measured from the original surface, greater than the

range of bombarding ions.

TIME

CAD CI IDCA^C" TO COME WITHIN DISTANCE R OF POINTx,

X| >R

TIME REQUIRED

TO ERODE SURFACE DISTANCE R

X 1

POINT X| REMOVED

FROM TARGET

A

1 ^ M- " ^

1

1

1

1

X —R 1

1 / ^-isX^j

"^MAXIMUM

DAMAGE

= 0(x)-|-

t BOMBARDMENT TIME

Figure 8. Damage formation at the point Xi,

described in the previous figure, versus bombardment time.

tance R, and there will not be any radiation

damage formation at x-^ until the surface has

eroded to the extent that the ions will penetrate

to this point. Thus, the surface must erode a

distance x^ — R before damage is formed at

X,. The total amount of damage at the point x

as a function of time is illustrated in figure 8. A time t= (x^-R) /IS will elapse before the

surface has moved sufficiently to permit the

penetrating ions to reach a depth, measured

from the moving surface, equal to the range R.

Once the ions penetrate to the depth x^, the

damage will increase with time. Inasmuch as it was assumed that the damage produced by each

incident ion is the same at every point along

its trajectory (see fig. 3), the damage at a

point X, will be proportional to the total number of ions passing this point, and, if the cur- rent is constant, the damage will increase

linearly with time. The time required for the

surface to erode from its original position to a depth X is simply x/IS. Clearly, the maximum

damage at the depth Xi occurs just as this point

is reached by the eroding surface. Also, it is

meaningless to talk about the damage concentration at any point after it has been eroded

away. Thus, the damage at the point x,, as a function of time D{x,t) is given by the following expression:

Q{x)dt=IQ{x) V x,-Rl

Xt — R

IS

when^<i<-§

D{x,t)=0, when i<^i^or *>-g .

The maximum damage which occurs during

sputtering occurs when t = x,/ (IS), or

g

163
TIME REQUIRED TO ERODE SURFACE DISTANCE R

/ '

TIME

REQUIRED

TO REACH

POINT

1

K 2

t > POINT

1 IS'

r ^

1 REMOVED FROM TARGET

y

" \/

1 /

1 y 1 X

/p-—MAXI MUM

1 DAMAGE =Q(x)

R

"s~

1 R

t , BOMBARDMENT TIME

Figure 9. Damage formation versus bombardment time at a point X2 which

is between the original surface and a

depth equal to the range of the bombarding ions.

Note that D^ax does not contain the time t or

the ion current / ; thus this simple computation

leads to the following important generalization.

Namely, the total amount of radiation damage

introduced into a target during sputtering has

a maximum value, given by Q(x)R/S, and this

quantity is independent of the ion current or bombardment rate.

4.4.4. Damage at a Depth Less than the Range

Consider next the rate of damage formation

at a point X2 which lies between the original

surface and the maximum penetration depth

of the bombarding ions. This situation is illustrated in figure 9. Inasmuch as the point Xo

lies within the range of the bombarding particles, the damage will begin to accumulate as soon as the bombardment has commenced and

will accumulate for a time t = X2/IS, when the

surface reaches the point X2. Thus in this case the damage D(x2,t) is given by the expression

t

D(x2,t) = J

IQ(x)dt = IQ(x)t,0<t<j^ '

0

Again, it is meaningless to discuss the damage

at the point X2 after this point has been re- moved by the erosion process. In this case the

damage is always less than the maximum obtainable.

4.4.5. Spatial Distribution of the Radiation Damage

Consider next the distribution of radiation

damage below the surface at various times

during the bombardment. The damage distributions are illustrated in figures 10 and 11. The

X, DISTANCE FROM ORIGINAL SURFACE

Figure 10. Damage versus depth distribution for a time t < R/(IS) after

the bombardment has commenced,

X, DISTANCE FROM ORIGINAL SURFACE

Figure 11. Damage versus depth distributions for a sequence of times U, t2, . . . t„ after the bombardment has commenced.

line from the origin to the intersection of the

maximum defect line and the line parallel to

the defect axis at x = R is the locus of the eroding surface during bombardment. The heavy

lines, indicating the damage level, were computed from the D{x,t) expressions by inserting

the appropriate values of t^, tz, etc. These considerations lead to the following important generalizations. First, as stated above, when the

material removed by ion-beam erosion reaches

a depth equal to the range of the bombarding

ions, the radiation damage is a maximum. Secondly, the distribution of damage with respect

to the surface of the figured material will re- main unaltered by additional figuring. Incidentally, one can readily show that these conclusions are valid for any reasonably smooth

Q{x) curve, and more specifically they are valid

for the observed Q{x) curve shown in figure 3. It is important to reiterate that the radiation

damage formation kinetics described in this

section are subject to at least one important

limitation. Namely, it was assumed that all

annealing effects were negligible. Annealing

effects are described in the next subsection and,

it will be seen, they can be appreciable.

164
4.5. Radiation Damage Annealing

4.5.1. Measurement of lon-Bombardment-Induced

Radiation Damage and Annealing

As explained above, very few, if any, measurements of ion-beam bombardment-induced

radiation damage have been made which are

directly applicable to figuring considerations.

Furthermore, very few studies have been made

which provide information about the annealing

of ion-beam-induced damage that can be applied to figured optics. Specifically, the dimensional stability of an ion-beam-eroded surface

has not been thoroughly investigated. Consequently, the discussion on annealing which follows is based entirely on annealing studies

made on reactor-irradiated glasses and especially fused silica. In fact, this treatment can be

regarded as an extrapolation of information on

the annealing of reactor-radiation-induced

damage to the ion-beam bombardment case.

4.5.2. Annealing of Reactor-Induced Damage in

Fused Silica

Studies on the annealing of the reactorinduced radiation damage in fused silica will be

described very briefly since most of this material has been described elsewhere [7,9] . During reactor bombardment both oxygen and silicon atoms are ejected from their normal lattice

positions and both vacancies and interstitials are formed. In this respect, ion-beam-induced

and reactor-induced radiation damage are very

similar. Part of the oxygen vacancies which are produced during the original bombardment

will, at the same time, be converted into color

centers by trapping electrons. The remaining

oxygen vacancies can be converted into color

centers by exposing the sample to a source of

purely ionizing radiation at a later time. Thus,

the concentration of oxygen vacancies can be

determined by making optical absorption measurements on these color centers.

If, after a sample has been irradiated, it is

heated at a temperature above the irradiation

temperature some or all of the color centers

will disappear. The reduction in color-center

concentration is attributable to two factors.

First, some of the oxygen vacancies are "annealed out" by capturing peregrinating oxygen

atoms. Clearly for each vacancy destroyed in

this manner one color center will be destroyed.

However, consider the following complication.

At certain temperatures the heat treatment will

cause thermal untrapping of electrons from

vacancies. In this case the vacancies remain in

the material but they do not contribute to the

color-center absorption. However, the remaining vacancies can be reconverted to color centers by exposing the sample to ionizing radiation after each heat treatment. Usually this

conversion would be carried out at room temperature.

>-<

ccH

CD<

<

<

1^' ANNEAL AT TEMPERATURE T,

Q^Z"'' ANNEAL AT TEMPERATURE T2>T|

o

° o

I I I

o o

J L ^

o O O o

\ L

5 10 ANNEALING TIME ( m i n )

Figure 12. Typical damage versus annealing time data for two different temperatures.

A typical and somewhat idealized radiation

damage annealing measurement is illustrated

in figure 12. This might apply to the following

sequence. The sample is irradiated in a reactor

at 70 °C. Subsequently at room temperature it

is exposed to ionizing radiation and the oxygen

vacancy concentration determined from optical

absorption measurements. Next, the sample is

annealed at temperature T^, greater than 70 °C,

returned to room temperature, exposed to iono

cc<

ccH

CD

or<u

o

<<

I .0

0.5 —

0.2

0.1

SLOW COMPONENT)

\ y (FAST COMPONENT)

\

I I I ""^i \ \ \ \ L

5 10 ANNEALING TIME (min)

Figure 13. The damage versus annealing time

data, shown in the previous figure, plotted on a semilog paper. This illustrates that the

annealing occurring at a fixed temperature

can be approximated by two decreasing

exponential functions representing a fast

and a slow annealing component.

439-921 O - 72 - 12 165
izing radiation, and the optical absorption measured again. Thus the radiation damage or, more specifically, the oxygen vacancy concentration has been determined after irradiation

and after the initial anneal. These damage

measurements correspond to the uppermost two

points in figure 12. The remaining points on

this curve are obtained by repeated anneals at

temperature T^. Next, the annealing temperature is raised to a value T2 greater than and

repeated anneals carried out at the higher temperature. In this vi^ay the lovi^er curve in figure

12 is obtained. These annealing curves can be

approximated by decaying exponential expressions. Thus v^^hen the data points on figure 12 are replotted on semilog paper, it is seen that

each of the curves can be divided into two components. This is illustrated in figure 13. Obviously it is appropriate to refer to these as the

"fast" and "slow" components. In other words,

the radiation damage annealing occurring at

temperature can be approximated by the following expression:

D{t) = (Initial Damage) — (Fast Component) — (Slow Component)

or D it) = Do —DofS-^f^—Dose'^'^

where

D{t) = damage remaining after total an- nealing time t. Do = initial damage concentration.

Dof = that part of the damage which anneals rapidly at a given annealing

temperature.

A/ = fast annealing component decay

constant.

Dos = that part of the damage which anneals slowly at a given annealing

temperature.

As = slow annealing component decay

constant.

This separation into components can be demonstrated analytically if it is assumed that the

activation energy for annealing is not constant

but can be expressed as a temperaturedependent function. This demonstration is too

long to present here. In addition, it should be

mentioned that radiation damage annealing

studies employing both chemical and physical

measurements provide empirical evidence for

the existence of fast and slow components.

An additional useful annealing characteristic

can be described with the aid of figure 13. Both

the magnitude and the slope of the fast-decaying

component of the To curve increases as the

temperature difference T2 — Tj becomes larger.

The intensity of the slow To component also decreases; however, its slope changes very

slightly.

The results of a careful study of the annealing of reactor-induced radiation damage in

fused silica are shown in figure 14 [7, 9] . When

this material is irradiated in a reactor or with

sufficiently high energy gamma rays, the absorption band, i.e., the color center, at 5.82 eV

increases as the dose increases. Additional studies employing optical and ESR techniques provide evidence indicating that this band should

be attributed to an electron trapped on an oxygen atom vacancy [10] . It is a good example of

the type of defect described above which can be

observed when it has trapped a charge but

cannot be detected when it is uncharged. Furthermore, since it is associated with a fundamental defect, namely a vacancy, it should be a

good "indicator" of the extent of the damage

in fused silica. The data in figure 14 were obtained by repeating the procedure described

above until all of the radiation-induced change

in absorption had been "annealed out."

Several useful comments can be made about

the results in figure 14. First, note that all the

radiation-induced defects are removed by heating slightly above 700 °C. This provides illustrative material for a point made previously.

Namely, this particular set of data may not be

a true indicator for all of the radiation damage

which could be present in this sample. For ex- ample, it is known that fused silica contains

radiation-induced absorption bands in the vac- uum ultraviolet. These were not studied during

these measurements and could possibly be attributed to still other radiation-induced defects.

Thus, if the sample had to be heated above

700 °C to remove all of the vacuum ultraviolet

bands, one would have to conclude that they

were better radiation damage indicators. Second, note that the damage decreases slowly

from the irradiation temperature to approxi-

» DEFECT CONCENTRATION, FROM SATURATION COLORING

• COLOR REMAINING AFTER FIRST ANNEALING SEQUENCE

CO

. 1-o < UJ .0-

S 0.8|- o

o R 0.6 S 0.4 <

a:

cr.

0.2

5.82 eV BAND

IRRADIATION - TEMPERATURE

t 70 °C _l 1 I 0 200 400 600

TEMPERATURE. °C

FiGUFE 14. The annealing of reactor- radiation-induced defects in fused

silica. The defects removed by each

increase in annealing temperature

corresvonds to the fast component

described in the previous figure.

The 5.82 eV band has been attributed to a color center formed when

an electron is captured by an oxygen vacancy.

166
mately 700 °C. Thus, to remove all of the damage it is necessary to heat the piece above this

temperature. However, it is entirely possible

that a reliable optic can be achieved by annealing at a considerably lower temperature. Third,

note that all of the coloring is removed by

heating at 550 °C but approximately 1/3 of the

radiation-induced defects remain at this temperature. Thus, in this case the removal of the

radiation-induced browning does not indicate

that the sample is free of defects. It does show

that the coloring can be eliminated at a temperature considerably below the defect annealing temperature.

4.5.3. Figuring Conditions for Minimizing

Radiation Damage

Some practical conclusions about the figuring

conditions that will produce optics containing

a minimum amount of radiation damage re- sult from the considerations given above.

First, it is apparent that the damage will be

minimized if, during bombardment, the temperature of the working piece is allowed to become as high as possible. To attain a high temperature one would maximize the beam current

and adjust the impact area size to attain the

maximum desirable temperature for as long as

possible. Also, the beam should be moved

across the working pieces as slowly as possible.

From the damage removal viewpoint, it is

clearly better to make a few passes at high

current density than many passes at low cur- rent density.

Second, consider the situation requiring that

a large amount of material must be removed

from a few areas and very little or none at all

from other areas. After the required material

has been removed from a given area the radiation damage immediately below the surface

will be a maximum. In contrast, other areas which have been eroded very little or not at all

will contain little or no damage. Thus, in the

worst possible case, this non-uniform distribution of damage may be conducive to subsequent

surface distortion. This situation might be alleviated by covering the entire surface with a uniform damage distribution. This could be

achieved by first figuring the surface to the

desired configuration and then removing a

layer one range thick from the entire surface.

Alternatively, it might be possible to remove

most of the required material at high-beam

energy and then, as the actual surface approaches the desired configuration, reduce the

beam energy so that the damaged layer becomes

very thin or approaches zero. Third, it is entirely possible that radiation

damage introduced during the figuring process

could be removed by subsequently heat-treating

the entire piece. This approach might be useful

for removing the radiation-induced coloring.

Often this can be removed by heating to a relatively low temperature.

4.6. Future Radiation Damage Studies

As indicated in the introduction to this section, there is very little, if any, reliable information on radiation damage effects in figured

optics. In particular there is very little information on the annealing of damage. Obviously,

if the figuring process produces a negligible

amount of damage it will be unnecessary to

study the annealing characteristics in detail.

However, before ion-beam figuring can become

an accepted technique it must be demonstrated

that the figured optic possesses long-term dimensional stability. Furthermore, it must be

demonstrated that ion-beam erosion does not

produce any unanticipated effects. To give just

one example in this latter category, consider

the possibility that the eroded surface is con- siderably more susceptible to chemical attack

than the normal surface.

These few comments on future work have

been included to make it explicitly clear that

additional research on ion-beam erosion is re- quired before one can regard this as a practical

technique for figuring optics.

5. Applications of Ion-Beam Figuring

It is appropriate to conclude this paper with

a very brief discussion on applications of ionbeam figuring. To begin, inasmuch as it is possible to follow the progress of the figuring during ion-beam erosion, this immediately

suggests several unique applications. First, this

would appear to be a particularly promising

alternative to "hand finishing" or for "touching

up" optics which had been shaped close to their

final configurations by conventional polishing

techniques. Second, it can be used to produce

aspherical or nonspherical surfaces from spherical surfaces if the amount of material to be

removed is not too great. Figure 15, which was

published previously [11], is a paraboloid

formed by ion-beam erosion from a spherical

surface which had previously been prepared

by conventional polishing. Third are optics

containing non-axially symmetric configurations. For example, figure 16 shows grooves cut

into an originally flat surface. Actually, this

picture was obtained as part of a series of ero- sion rate measurements [3] ; however, it clearly

illustrates non-symmetric figuring. A large

number of interesting possibilities become apparent when partial or complete on-line computer control is combined with the ion-beam

figuring process. In principle, it would appear

that this arrangement makes it possible to obtain almost any non-spherical and non-axially

symmetric surface configuration. In fact, it

167
I 5 min 100 kV

Figure 15. The fringe pattern, observed with an

interferometer, from a paraboloid produced by

ion-beam figuring. From Ref. [11].

would appear to be particularly well suited for

the production of Schmidt plates and similarly

shaped pieces. Finally, one "far out" application must be mentioned. Ion-beam figuring

would appear to be a particularly good technique for correcting astronomical reflectors on

satellites or on the surface of the moon.

6. References

[1] Carter, G., and Colligon, J. S. Ion Bombardment

of Solids (Heinemann, London, 1968).

[2] Proceedings of an international conference on

atomic collisions and penetration studies with

energetic (keV) ion beams, Can. J. Phys. 46,

449 (1968).

[3] Ion beam optical figuring. KoUsman Instrument

Corp. Tech. Rpt., Contract No. F33615-68-C1146 (1968).

[4] Meinel, A. B., Bashkin, S., and Loomis, D. A.

Controlled figuring of optical surfaces by en- ergetic ionic beams, App. Optics 4, 1674 (1965).

30 m i n 50 kV

BK-7 GLASS

65° ANGLE OF INCIDENCE

50/iA BEAM

Figure 16. The

fringe pattern, observed with an interferometer, from

BK-7 glass after

it had been subjected to ion-beam

erosion using different bombardment conditions.

[5] Schroeder, J. B., Bashkin, S., and Nestor, J. F.

Ionic polishing of optical surfaces, App. Optics

5, 1031 (1966).

[6] Primak, W., and Bohmann, M. Radiation Damage

In Ceramics, Progress in Ceramic Science,

Vol. 2, p. 104, Ed. J. E. Burke (Pergamon,

London, 1962).

[7] Levy, P.W., Physical properties of irradiated

ceramic materials, nuclear science and technology for ceramists, Nat. Bur. Stand. (U.S.),

Misc. Publ. 285, 21 pages (1967).

[8] Hines, R. L., and Arndt, R. Radiation effects of

bombardment of quartz and vitreous silica by

7.5-keV to 59-keV positive ions, Phys. Rev.

119, 623 (1960).

[9] Levy, P.W., Use of color centers for detection

and measurement of defects, the chemical and

physical effects of high-energy radiation on

inorganic substances, p. 3 (Amer. Soc. for

Testing and Materials, Philadelphia, 1964).

[10] Lell, E., Kreidl, N. J., and Hensler, J. R. Radiation effects in quartz, silica, and glasses. Progress in Ceramic Science, Vol. 4, p. 1, Ed.

J. E. Burke (Pergamon, London, 1966).

[11] Narodny, L., and Tarasevich, M., Parabolid figured by ion bombardment, App. Optics 6, 2010

(1967).

Diiscussion

Wiederhorn: As a practical consideration,

what rate of thinning do you get on material

such as alumina, glass, etc.?

Levy: 100 angstroms per minute is quite

feasable.

GiELiSSE: You stressed the point that the ex- pression for the maximum concentration of defects contains neither the time or ion beam

current. Does not the range which does occur

in the expression, depend upon the level of ion

beam current?

Levy: No, the range depends only on the energy, not on the current density.

K. Smith: This application, doesn't it pro168
duce strain on the surface due to the differential

of the temperature?

Levy: Yes, it does. In fact, the principal

practical problem that seems to occur is the distortion of the surface during bombardment due

to thermal effects.

K. Smith: Does that have a v^eakening effect

on whatever you are trying to make?

Levy: I doubt it. I can visualize certain materials which seem to be somewhat more sus- ceptible to crazing and things like that. But any

material can be bombarded at a rate high

enough to melt the surface.

Heuer: We have used a similar technique to

make electron microscope foils. By taking

stereo pairs, the radiation damage can be seen

as very small defects and they are indeed confined to the top and bottom of the foil and not

anywhere in the interior bulk.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 1970, Gaithersburg, Md. (Issued May 1972).

169

The Effect of Sputtering on Surface Topography and Strength of Ceramics

R. W. Rice

Chemistry Division, Naval Research Laboratory

Washington, D.C. 20390

The topography developed as a result of R.F. sputtering on a variety of ceramic

surfaces is described showing that surface cracks, scratches, pits, etc., are fairly

rapidly rounded out. Generally, this occurs without differential grain boundary

sputtering. However, some bodies develop a rough finish due to nonuniform microstructural removal. Limited trials with ion beam sputtering show it also rounds

cracks, etc., but does not develop a rough finish on bodies that will with R.F.

sputtering.

Because of the rounding of stress concentrating features and generally similar

effects on single and polycrystalline bodies, sputtering was investigated as a substitute for flame polishing, but with much wider applicability. Flame polished

strengths were not obtained. However, some materials (e.g., MgAl204) do show greater

improvements in strengths than others (e.g., AI2O4) as a result of sputtering. These

differences will be discussed and results compared with other methods of surface

finish.Both the detailed nature of the sputtered surface as well as enhancement of

grain boundaries and other microstructural features under certain conditions indi- cate that sputtering may also be useful as an etching technique.

Key words: Carbides; glass, ion beam sputtering; nitrides; oxides; RF sputtering;

strength ; surface finish ; twins.

1 . Introduction

Obtaining nearly perfect surfaces by flame

polishing or chemical polishing is very valuable

in the study of mechanical properties of ce- ramics. However, the applicability of these techniques is limited. High melting points, or

reactivity, prevent flame polishing of many materials. On the other hand, the large, often

columnar, grains that result from melting the

surface of polycrystalline bodies generally preclude the use of flame polishing polycrystalline

bodies (e.g., AI2O3) which can be flame polished

in single crystal form. Further, flame polishing

generally works well only for round rods. Techniques of chemical or electropolishing (of con- ductive ceramics) for producing good surfaces are known for only a few materials. Preferential attacks on certain crystal surfaces and at

grain boundaries also limit the application of

these.

Etching to reveal grain bounda/ries subboundaries, dislocations, and twins in ceramics

is also important in studying their mechanical

behavior. While chemical etches have been developed so that a substantial number of ce- ramics can be suitably etched, improvements are

still needed.

Sputtering ^ which is the process of removing

material from a surface by bombardment of the

surface with ions, appeared to be a way to obtain near perfect surfaces on many ceramics.

The achievement of such surfaces was sug-

• Several reviews of the sputtering process are available. Some

of these are tabulated in references [1-5]. (Figures in brackets

indicate literature references at the end of this paper.)

gested by both the fact that (a) removal is

normally atom-by-atom, and (b) sputtering

has been very successful in thinning a wide

variety of ceramic materials for electron microscopy [6, 7] which require specimens of such

thinness that little variation in thickness can

be tolerated. It was also felt that etching,

which tends to be contrary to obtaining near

perfect surfaces, might be obtained by varying

sputtering conditions (e.g., bombarding ion en- ergy), especially on certain materials (e.g.,

those with two phases). A search of the literature indicated other examples of obtaining

good surfaces [1, 7-9] and some examples of

etching [1, 7, 8, 10-12] by sputtering. However, these past studies have generally not

studied the effect of prior surface finish on the

resulting sputtered topography, and none have

considered what effect sputtering has on mechanical properties. The latter is important not

only because sputtering may affect surface sensitive strengths, but also because sputtering is

being considered as a means of machining and

finishing ceramics [14-16].

This paper presents results of a study of

sputtering of ceramics. The initial emphasis

was on determining whether or not strengths

achieved by flame or chemical polishing could

be achieved by sputtering. When it was found

that such high strengths were generally not obtained, the emphasis was on determining what

strength changes did result from sputtering.

This has also resulted in considerable information on the type of topography developed. While

it has not been possible to study etching in detail, some examples of it are also reported.

171
Table 1. Materials studied *

Zr02, unstabilized crystals

ZrOo, CaO stabilized (9 W/O)

crystals

Zr02, Y2O3 stabilized (14.5 W/O)

crystals

ZrOz, Partially stablized

(2.9% CaO)

ZrO:, MgO stablized

Skull

melting

Sintered

Sintered

Zircoa C

Zircoa 1027

W. & C. Spicer Ltd.

St. Marys Winchcomb

Glouchestershire

England

, Zircoa Corporation

Solon, Ohio

' Other materials studied have been listed in table 1 of ref. [17].

2. Experimental Technique

2.1. Specimen Preparation

A variety of samples were used as listed in

table 1 and reference [17]. Grinding, gas polishing, and flame polishing of specimens were

as discussed elsewhere [17-19]. Mechanical polishing was done by standard methods.

2.2. Sputtering

Samples were sputtered by both radio frequency (RF) and Ion Beam (IB) sputtering 2. RF sputtering uses ions generated in a glow

discharge which are accelerated to the sample

(target) by the potential difference between

the target electrodes. The simplest type of ar- rangement and the one used here is sketched in

figure 1 ; however, several other variations

exist [2-5]. RF rather than dc potentials, are used to prevent possible charge buildup on nonconducting samples. Since glow discharges depend on the pressure, voltage, and electrode

separation, the variation of ion energies is limited. Also, ion incident on the target is essentially normal. IB sputtering uses an ion gun to

generate and accelerate ions (fig. 1). Since the

latter is completely independent of the target,

the ion energies and angle of incidence are

quite variable. Specimens are usually rotated

to maintain uniformity. Although charge buildup is generally not a problem with IB sputtering despite its dc nature, there are techniques

of neutralizing the beam, if necessary.

RF sputtering ^ was done with a commercial

unit (manufactured by Materials Research Corporation, Orangeburg, N.Y.), using electrodes

approximately 5 cm in diameter (2 in).

' It should be noted that sputtering here refers to the removal

of surface material from a bulk specimen, not the deposition of the removed material. Sputtering has come into wide usage,

especially in electronics, as a technique whereby the atoms re- moved from one surface are deposited on another surface similar

to vacuum deposition. Such sputter deposition is not concerned

with the surface left on the material from which the atoms are removed (the "target")- Confusion sometimes results, since sputter

deposition is often referred to only as sputtering and the re- moval (sputtering) process as sputter etching which should not

be confused with grain boundary, etc., etching discussed in the

introduction.

' RF sputtering was done by Drs. G. Sigel and W. Anderson of the Solid State Division of the U.S. Naval Research Laboratory.

^ IB sputtering was done by Dr. B. Hockey of the Institute for Materials Research, National Bureau of Standards.

RF SPUTTERING ION BEAM SPUTTERING

Figure 1. Schematic representation of RF and ion

beam sputtering.

Chamber pressures were typically a few fiin,

with potential difi'erences of about 2500-3000V

(at 13.56 MHz). Because considerable heating

(e.g., to of the order of 1000 °C) could occur with some specimens, especially at power levels

of 0.4-0.5 KW, most sputtering was at 0.1-0.3

KW. Sputtering times were normally about 50

hours in steps of 2 to 6 hrs.

Because the specimen holder is mounted in an

inverted position, it was necessary to hold the

specimens to the holder. This was done with

wires, as shown in figure 2. IB * sputtering was done with a commercial

unit (manufactured by ALBA, Paris, France)

normally used for thinning specimens for electron microscopy. The unneutralized argon beam

was typically accelerated at 6 kV and struck

the specimen holder at an angle of 15-20 from

the horizontal (the thickness of specimens and

the specimen holder-ion gun geometry did not

allow lower angles of incidence.). The aluminum or stainless steel specimen holders were

rotated at about 15 rpm. In order to avoid possible loss of specimens due to rotation, specimens were held down by wires near each end.

Five specimens were mounted, separated by

1-3 mm, on the approximately 3 cm diameter

holder. Because of the beam diameter, sputtering was greatest in the central half of the

holder. Sputtering times were typically 50-80

172
Figure 2. RF sputtering holder. Note ceramic bars

held in place by wires.

hours. A few preliminary 2-hour trials were

also made with a neutralized beam unit ^

In order to sputter as many specimens as

possible, and to avoid possible differential sputtering on different surfaces of round single

crystals, flat bars were used. Initially these

were about 2.5 x 5 mm in cross section, but

most specimens were about 1.8 x 3.8 mm. The

bar edges were rounded in order to obtain some

sputtering of the edges.

2.3. Specimen Evaluation

Surfaces were examined by optical and replica electron microscopy. It should be noted that

most optical photos were made using Nomarski

interference contrast to emphasize surface topography, and hence should not be taken as representative indication of the roughness of the

surface.

Ambient mechanical strengths were measured in three-point bending using an Instron

test machine with a head travel speed of 0.127

mm/min. Most specimens were tested on a span 1.27 cm and a few were tested on a span

of 0.89 cm.

3. Experimental Results

3.1. RF Sputtering Topography of

Single Crystals

RF sputtering of single crystals of oxide

ceramics generally developed a topography best

described as a series of shallow rounded depressions. The depth of these averaged about 1 fxm,

but the number, size and extent of these depended primarily on the prior surface finish.

Thus, for example, they completely covered

the surface of ground specimens as shown in

figures 3 and 4, with a remnant pattern of

grinding stria generally being retained (fig.

3). Progressively finer mechanical finishing

prior to sputtering reduced the density of these

depressions. Even with a relatively poor polish,

most of the depressions are separated as shown

in figure 5. This figure also shows that each pit

left in the mechanically finished surfaces becomes a rounded depression on sputtering, and

that considerable rounding occurs within two

hours of sputtering. Electron micrographs

(fig. 6) show more clearly the extent of rounding features with extended sputtering. While

some smaller depressions normally occur inside

of larger ones during sputtering, this became

particularly pronounced (fig. 7) when significant specimen heating occurred.

This development of depressions is generally

true for various orientations of ALOg, spinel

(AlaOs-rich), zirconia (CaO or Y2O3 stabilized),

TiOo, and MgO crystals. Electron microscopy

showed that a variety of precipitates were re-

° Tests run by Dr. H. Garvin of Hughes Research Laboratory,

Malibu, Calif.

50/im

Figure 3. RF sputtering of ground surfaces.

Ruby ground in vertical direction. Note remnants of grinding streaks. Note also the

apparent imperfection revealed by sputtering (approximately rectangular object near

top of photo).

173
Figure 4. RF sputtering of ground spinel crystal. (A)

as diamond ground (grinding direction vertical) (B)

similar area of same specimen sputtered approximately 54 hours.

vealed in the stabilized Zr02 crystals by sputtering, corroborating optical indications of

their presence. Unstabilized ZrOa crystals

tended to develop a pattern that was apparently modified by differential surface energies

giving a partially faceted appearance. Some

Ti02 crystals, and to a lesser extent MgO, also

tended to do this some. Figure 8 shows that

sputtering did occur part way down the sides

of the beveled bars. While the Nomarski contrast exaggerates the edge irregularity, many

sections of the edges were somewhat irregular.

Particles of debris, apparently accumulated

from material removed from the surface, were

often found in the grooves formed between

specimens by their beveled edges. Some of this

appeared to be bonded to the specimen.

Flame polished sapphire generally did not

develop depressions until after several hours of

sputtering. Some depressions did develop fairly

early in areas where the flow of a wave of

melt had been interrupted [19] and similar areas of expected stress concentrations. When

Figure 5. RF sputtering of ruby. (A) Specimen as

polished with 1 micron diamond paste. (B) Most

of the same area sputtered 2 hours. Note cor- respondence of pits and their rounding with sputtering (portion of scribe mark and arrows to same largest pit index photos).

more depressions began to occur with continued sputtering, it was most commonly as clusters in such areas where densities of 10 «-10 ^ per sq. cm. were reached in contrast to 10 *-10 ^ away from the areas. Gas polished specimens

[18] developed, even fewer depressions, and

they were generally random, not occurring in

clusters.

Besides rapid rounding out of damage from

scribing ALOa and MgAloOi, some features

were often found associated with the scribe

marks. As shown in figure 9, features which

appear to be twins emanate from the tips of

many lateral cracks along scribe marks. Though

these twins could be seen by careful examination prior to sputtering, limited sputtering

(e.g., 2 hrs) appeared to bring them out more

clearly. Some could also be seen along a few

deeper grinding marks where surrounding

areas were clear enough to distinguish them.

These were typically 20 /xm long and varied

from 2 to 10/xm apart. Some unusual features

also developed around some of the larger lateral

174
25yu.m

Figure 7. Typical lengthening of major depressions

and development of many smaller depressions within

larger ones in RF sputtered samples that got hot

during sputtering. This suggests that depressions

may be due to dislocations, which move or multiply

due to heating.

Figure 6. Electron micrographs of crystal surfaces.

(A) As ground. (B) Polished (3 micron BX). (C)

Sputtered surface subsequent to polishing.

cracks as shown in Fig. lOA. Figure lOB

shows that lateral cracks emanating from

scribe marks on MgAl204 often ended in a depression, indicating that limited sputtering was

"etching" preferentially at or near the tips of

these cracks.

A variety of sporatic features was observed

that appear to be due to impurities and defects

in the crystals. A few more regular objects

such as the rectangular feature near the top of

figure 3 were observed. These are believed to be

voids. A variety of circular to very irregular

and very sharp to very diffuse objects were

observed that are believed due to impurities.

3.2. RF Sputtered Topography of Glass

Sputtering of soda lime glass generally developed shallow crescent-shaped depressions,

Figure 8. RF sputtering of specimen edge.

Top of photograph is looking down the edge

of an AI2O3 rich MgAl204 crystal specimen.

The edge was rounded prior to polishing,

then the specimen was sputtered approximately 54 hours.

175
Figure 9. Possible RF sputtering of twins along scribe

marks on ruby. (A) Specimen polished with 1 micron

diamond paste, then sputtered—surrounding areas

sputtered 5 hours, scribe mark sputtered 2 hours. (B)

A similar ruby with background sputtered 5 hours,

and scribe mark sputtered 2 hours. Note that apparent twins generally emanate from the tips of

lateral concoidal cracks.

as shown in figure 11. These followed the general trend found for depressions in oxide crystals. That is, their density was clearly increased

by abrasion. For example, completely covering

a glass section with scratch marks by hand

sanding produced a crescent density comparable to the density of depressions on ground

crystals. In partially abraded areas, these cres- cents tended to start from scratch marks and

spread out to cover intervening polished areas. All specimens, including previously flame polished ones, tended to develop a craze-type pattern, as shown in figure 12 after several hours

of sputtering. Crescents tended to form near

these "craze lines" and spread into intervening areas not already covered with them, e.g., on

areas of mechanically polished and especially

on flame polished specimens.

Small dots, which were apparently hillocks,

averaging about 0.2/xm in diameter were observed in densities of 10 '-10 ^ per cm " over

many areas by electron microscopy. Much

larger and less dense dark spots, sometimes

A

B

Figure 10. RF sputtering of cracks in ruby and spinel

crystals. (A) Features sometimes developing around

tips of large cracks on the side of a scribe mark on

ruby. (B) Apparent cracks seen along scribe marks

on some AUOs-rich MgAl204 crystal (approx. 3 MgO:

3.5 AI2O3). Note there always appears to be a round

sputtered depression at or near the tip of these

cracks. Background polished and sputtered 5 hours,

scribe marks sputtered 2 hours in both cases.

surrounded by several smaller satellite spots

were often seen optically, usually associated

with the craze pattern as shown in figure 12B.

On specimens sputtered at higher power (0.3-

0.5 KW), cracks (fig. 12) were often associated

with the craze pattern and these spots, which

appear to be pores. Longer sputtering (e.g.,

after several hours) generally developed a terrace-type topography (fig. 12B).

3.3. RF Sputtering Topography of

Polycrystals

In general, polycrystalline specimens developed the same general pattern of depressions,

but with two important differences. The first of

which is related to grain size. As seen in figure

13, sputtering of dense fine grain (approximately 3 fjm) hot pressed AloO., with a ground

surface results in a high density of very fine

176
Figure 11. RF sputtered soda lime glass. The specimen was sputtered 2 hrs.

before this electromicrograph was made.

depressions (e.g., compare with figs. 2-5). On

the other hand, a large grain (approximately

60 fxm Lucalox) specimen developed depressions

comparable with those on single crystals as shown in figure 14. The decreasing size of depressions with decreasing grain size appears

to be due to the fact that pits left from mechanical finishing which are a major source of

depressions usually decrease in size with decreasing grain size. The density of depressions

on the Lucalox was also quite high due to a high density of pits left from semi-polishing.

Subsequent thermal etching (by annealing) of

the AloOs showed that the depressions have no

definite association with grains or grain boundaries as shown in figure 14C.

Annealing of dense hot pressed ALOa bodies

prior to sputtering clearly reveals grain boundaries, which are then completely obscured by

long sputtering as shown in figure 15. A high

density of depressions also develops on such

annealed surfaces, but at a much slower rate

than on a machined one. The second difference between single and

polycrystals appears to be related to the pres- ence of grain boundaries or varying crystal

orientation in certain bodies. Thus, for ex- ample, commercial ZrO, stabilized with CaO or MgO initially had grain boundaries enhanced

as shown in figure 16, but these then became

obscured with further sputtering. Figure 16

also shows that previous pits (and voids) in the

177
B

50/i.m

Figure 12. Craze pattern and cracks on RF sputtered soda lime glass.

specimen result in depressions on sputtering,

and that polishing results in fewer depressions

as with oxide crystals.

In contrast to this, dense hot pressed MgO

specimens, progressively developed a rougher

topography. In the most extreme case, this appeared to be due mostly to a very different

sputtering rate of grain boundary material

(fig. 17A), while in the other case, this appeared to be due partially to differential re- moval from boundaries and to differential re- moval from different crystal surfaces. SiC

showed a somewhat similar trend (fig. 17B),

with a rougher topography forming a network

around patches of normally appearing sputtered areas. This network appears to start at

the grain boundaries, and widen with increased sputtering. Observations of fracture

surfaces normal to the edge of the SiC sputtered

surface showed that, while fairly well rounded,

the grain boundary areas were lower than the

center of the grains.

Trials with a medium grain commercial sintered AI0O3 (Alsimag 614), and dense hot

pressed MgAloOj, SijNj, and B4C, all indicated

that they behaved similar to that of the dense

alumina of corresponding grain sizes discussed

above. In BiC, clusters of pores within grains

(observed on fracture surfaces) were also re- vealed by sputtering.

3.4. Ion-Beam Sputtering Topography

Most specimens sputtered with the unneutralized ion beam developed topography very

178
Figure 13. RF sputtering of dense fine grain {Sixm)

AliOi. Specimen hot pressed from Linde A powder,

then air annealed at 1315 °C, then sputtered approximately 54 hours. (A) Optical photo. (B) Electron micrograph.

similar to that obtained with RF sputtering;

i.e., a series of depressions rapidly developed.

Their density and surface coverage generally

decreases with finer finish, and their size generally decreases with decreasing grain size as with RF sputtering. The main difference in

these depressions was a tendency for an elongated and apparently shallower nature. Length

to width ratios varied from just over one to as high as three to four (figs. 18, 19). As with RF

sputtering, remnants of grinding striations re- mained, and a few grain boundaries were re- vealed (figs. 18A, 19).

The main differences that were observed between RF and IB sputtering were with glass

and polycrystalline MgO and SiC. IB sputtering

did not result in the crescents observed in RF

sputtering. Instead, the glass developed what

appeared to be hillocks very similar to depressions in other materials such as ALOs and

MgAlsOi. The density and surface coverage of

these again decreased with finer finishing. However, as shown in figure 20, rounded trenches

were observed between some of the hillocks like

Figure 14. RF sputtering and annealing of Lucalox.

(A) Surface as semi-polished (with Linde C). (B)

Similar surface sputtered approximately 30 hours.

Note faint lines (especially on right) suggestive of

grain boundaries, which are more pronounced in this

photo than average. (C) Similar specimen annealed

in air one hour at approximately 1300 °C to reveal

grain boundaries. This shows that sputtering depressions are not related to grains or grain boundaries.

features, and the latter generally had several

dots on them (fig. 20A). The polycrystalline

MgO, which had developed fairly to extremely

rough topography in RF sputtering, responded

more like other materials in IB sputtering.

(There was, however, some preferential grain

removal as shown in fig. 21.) Polycrystalline

SiC developed a surface topography quite similar to other materials in contrast to RF sputtering.

Sample trials with a neutralized ion beam

sputtering device indicated the same general

trends in initial 2-hr runs.

179
B

Figure 15. RF sputtering of annealed AUO,. Linde

A AI2O3 powder hot pressed to near theoretical

density, diamond ground, then annealed in air at 1790 °C. (A) As annealed surface showing rounding

of (vertical) grinding marks and thermal etching of

grain boundaries. (B) Similar area of same annealed

surface after approximately 54 hours of sputtering.

3.5. Sputtering Rate and Behavior

Initial sputtering was done at 0.4-0.5 KW;

however, most was done at 0.1-0.3 KW since

sample heating was more frequent and excessive at the higher power. This varied from run

to run and from sample to sample, apparently

due to varying thermal contact with the holder.

Some specimens glowed a dull red, indicating

temperatures of the order of 1000 °C were

reached. This glow was clearly not luminescence since it slowly built up during initial

sputtering and diminished fairly slowly when

sputtering was stopped. Some specimens also

luminesced, e.g., MgO crystals gave a light

green color, and different spinel specimens gave

different degrees of a light green color. During

most of the initial runs at higher powers, ZrOg

specimens turned black in many areas, some

of which were on the side and back as well as the front. The depth and distribution of these areas was very consistent with reduction due

to excessive heating because of poor and inhomogeneous thermal contact. TiOa crystals

also darkened substantially during RF sputtering.

While measurements on sputtering rates

were not a primary goal, some estimate of these

can be made from micrometer measurements

(at the ends of bars) before and after sputtering. In RF sputtering, the overall sputtering

rates averaged slightly over 1 /xm/hr, or about

200A min. This is in good agreement with lit- erature values for ALOs which constituted the

greatest single portion of the samples. There

was a fair amount of variation. For example,

glass measurements indicated a rate about

twice the average. Also, there was a trend for a higher (e.g., 10-20%) sputtering rate over the

outer regions of the sample holder. In general.

B

25/im

Figure 16. RF sputtering of CaO stabilized ZrOi

(Zircoa C). (A) Specimen as polished with 3 micron

B4C. (B) Same area sputtered 2 hours. (C) Same

area sputtered 5 hours. Note diamond scribe marker

across upper left corner in all photos.

180
Figure 17. Examples of rough topography on cer- tain RF sputtered polycrystalline samples (A)

MgO (B) SiC.

20 to over 100 ju.m were removed from samples.

Cutting power from 0.4-0.5 KW to 0.2-0.3 KW

did not appear to make significant changes in

sputtering rates.

Sputtering rates estimated for IB sputtering

from measurements on the ends of samples

were about a fifth that of RF sputtering, i.e.,

about 0.2 ju,m/hr. However, these are definitely

low since the amount of sputtering clearly diminished toward the outer edges of the sample

holder as would be expected. Differences in

sputtered topography between the center and

ends of the specimens indicated that the differ- ences in sputtering rate were of the order of

two, which would be in approximate agreement

with a maximum expected rate of the order of

0.5 fim/hr. Thus actual total removal over much

of the sample was estimated at 10-30 fxm.

3.6. Strength and Fracture of

Sputtered Specimens

The strength tests of RF and IB sputtered

glass and oxide crystals are summarized in

table 2 along with comparative data for other

finishes. These clearly show that, in general,

flame polished strengths are not achieved. The

levels of strengths achieved are typically about

the same as achieved by mechanical polishing.

However, this level of strength was generally

achieved regardless of prior finish. Thus, for

example, samples such as spinel, sapphire, and

glass ground perpendicular to the tensile axis

were clearly increased in strength. There was

also an indication that ground specimens

tended to give higher strengths after sputtering

than did polished specimens. This is indicated,

for example, by the spinel in table 2, and in individual specimens of sapphire and glass.

There are two other possible trends indicated

by these tests. First, the limited IB sputtering

results indicate that it makes more improvement in strength than does RF sputtering. This

is indicated particularly by the averages of

ruby and glass, where maximum values

achieved by IB sputtering of these materials

(with ground surfaces) were about 2 to 3

times the level achieved by mechanical polishing. Second, TiO^, ZrOs, and MgALO, indicate

possible increases in strength over mechanically polished levels with either sputtering technique. While the scatter of the ZrOo crystal

data makes results uncertain, part of this may

be due to the fact that specimens from two different crystals of unknown orientation were

used. The sputtered specimen from one crystal

was twice as strong as the polished specimen.

Much smaller differences were observed in the

other crystals. The spinel data of table 2 suggests that some crystals were weakened. However, direct comparison of polished and sputtered tests (cut parallel to the boule axis) on

the same bar showed the sputtered specimens

average 10-20 percent stronger. This is due to

the fact that these specimens from that particular boule had lower than average strengths.

Annealing of sputtered rubies in air at

1300 °C did not improve strengths. In fact, this

may have decreased them some, though the

scatter makes this uncertain. Strengths of

sapphire that was flame-or gas-polished prior

to sputtering averaged over 100,000 psi

(7 X 10«N/m^).

Tests were made on several polycrystalline

AloO, and MgAloO^, samples, as well as a few

polycrystalline ZrOo, B,C, SiC, and Si,N.

samples. Strengths of a variety of sputtered

(IB and RF) ALO:, bodies were clearly not significantly better than achieved by mechanical

polishing. Limited trials with commercial ZrOo

showed the same results. Limited tests with

439-921 O - 72 - 13 181
Table 2.

Strength of RF and IB sputtered glass and single crystals

(strengths in 1000 psi (and in WN/W))

Specimens prior to sputtering Specimens after sputtering

Orien- As-ground surfaces Aspolished No. of

l/CO I/O No. of

tests Material tation ' Parallel :Perpendicular RF XJt>

Slr^rlo OvfUd xliiic Mine* gidoo 0*10 co 14±1 (10) 10±1 (7) X f _ 0 y J.^ f 14±6 (10) 91 +9. aD MgAUO, II 21±1 (14) 44±12 (30) 7

MgAUOi II 34±2 (23) 37±9 (26) 28±8 (19) 5 38 ±15 (26) 4

MgA1.04 1 34±5 (23) 20±1 (14) 39±11 (27) 46±6 (32) 11 40±4 (28) 3

Sapphire II 48±9 (33) 25±7 (17) 53±15 (37) 58±8 (40) 3 52±3 (36) 4

Ruby 1 62±21 (43) 70 ±12 (48) 15 100 ±34 (69) 2

ZrO. (Y,03-

stabilized)

28±7 (19) 32±1 (22) 4

TiO. 1 13±1 (9) 11±1 (8) 22±7 (15) 33±3 (23) 2

'

II means the tensile axis was parallel with the boule axis; and 1 that the tensile axis was perpendicular

to the boule axis.

BtC, SiC, and Si3N4 indicated no change with

RF sputtering, but possible improvements with

IB sputtering with B,C giving somewhat

clearer results (e.g., a possible 20% increase).

Fine grain spinel bodies indicated improvements in strength of the order of 25-50 percent

due to RF sputtering, while large grain bodies

of dense spinel showed no definite change.

Fracture origins were located in many specimens, especially in glass and single crystals.

Many origins were from the edges with these

being more common in higher strength specimens, especially in glass and spinel crystals.

Edges of some RF specimens had debris attached at or near the origin.

4. Discussion

4.1. Surface Topography

Depressions similar to those observed in this

study have been reported by Dugdale and Ford

[10] on AI2O3 and (fused) SiOj. These were

attributed to porosity in the ALOg, but were

noted as an "interesting but unexplained phenomenon" in SiOa. This work, e.g., on ZrOa,

shows that pores can be sources of these depressions, but they are more frequently caused

by pits left from mechanical finishing which is

the likely source in SiOo. However, the appearance of depressions on certain areas of flame

polished specimens, and generally on annealed

specimens, as well as within existing depressions, suggests other sources. All of these observations are consistent with dislocations also

being sources of pits. This was corroborated by

etching studies of flame polished specimens

which showed a high density of dislocations in

the same type of areas in which many depressions developed [19]. The multiplicity of

smaller depressions within larger ones when

samples were heated during RF sputtering may

then be due to migration or generation of dislocations.

The apparent enhancement of features that are believed to be twins is reasonable since

sputtering rates are dependent on crystal orientation. Thus a twin will have a different orientation and hence different sputtering rate on

many intersecting planes. Besides apparently

etching twins and dislocations, some grain

boundaries (e.g., in ZrO.) were also etched by

RF sputtering for shorter periods (e.g., 2 hr).

Pores and impurity areas also appear to be

etched. The extremely irregular topography of

the polycrystalline MgO also is probably due to

impurities, since the author observes variable

concentrations of Ca and Si at grain boundaries

in this material. The less extreme MgO sputtered topography may be due in part to this

also. However, differential sputtering rates between different crystal surfaces may be more

important.

While the origins of the features on sputtered glass are uncertain, they are suggestive.

The association of crescent-shaped depressions

with abrasion suggests that they may be flaws.

The "craze-type" patterns may also represent

more extended flaws but these may well have

developed due to heating during RF sputtering.

Clearly, the black spots and the fine hillocks

suggest different degrees of inhomogeneity in

glass, with the latter possibly being related to

phase segregation.

These examples of etching indicate sputtering can be a useful etch technique. However, it

should be noted that once a surface feature is

182
B

Figure 18. Typical ion beam sputtered surfaces. All specimens sputtered

about 50 hours. (A) Polycrystalline Zr02, partially stabilized with

CaO previously mechanically polished. Note evidence of a grain

boundary which was not too common. (B) Sapphire, previously

mechanically polished. (C) Ruby previously ground in the vertical

direction.

developed it will continue to be preserved even

after its original cause is gone. Thus etching

by sputtering tends to give an integrated effect,

that is, showing features from various previous

levels on the same surface. This can be useful,

but might also be misleading. It is clear that

etching is best with a fairly high angle (e.g.,

normal) incidence. IB sputtering should also be

feasible in this orientation, and hence should

be as useful, or more so (due to acceleration

control) than RF sputtering.

On the other hand, while RF sputtering does

a very similar job on many materials, IB sputtering produces flatter surfaces on some materials. Heating during RF sputtering was also a problem. This could be reduced, especially in a system in which the specimens rested on a

flat surface rather being inverted, so gravity

would pull them onto the electrode rather than

away from it. However, electron bombardment

during RF sputtering means there will always

be greater heating there than in IB sputtering.

Thus, except for greater sputtering rates, RF

sputtering is less versatile than IB sputtering.

Recently, Schmidt et al. [20] have shown

that IB sputtering can be very useful in more

clearly revealing magnetic domains and providing more accurate magnetic measurements in

183
/Lt.m

Figure 19. Ion beam sputtering of a medium grain (about 10-20 m) commercial AliOi

(Alsimag 6H). Electron micrograph showing elongation of depressions, exposed

pores, and some traces of grain boundaries.

100 /im

B

H

10/i.m

250 /^m

Figure 20. Ion beam sputtered soda lime glass.

Figure 21. Ion beam sputtered polycrystalline MgO.

thin films (by removing mechanical damage and

related strains). One would also expect that the use of reactive gases during sputtering might

be very beneficial for etching purposes. Reisw^ig's [21] work showing excellent etching of

grain structure in graphitic materials, possible

deformation twins in graphite, and of the interface and -grain structure in Au-Ag diffusion

couples by sputtering with hydrogen is very

suggestive of what might be done. Thus sput184
tering to reveal microstructure and related

features appears to be quite promising.

4.2. Strength of Sputtered Specimens

The rapid rounding of surface pits and

scratches is clearly consistent with the fact that

substantially weaker specimens are strengthened by sputtering. However, most sputtered

specimens did not exhibit strength substantially

greater than mechanically polished specimens.

While complete explanations cannot be given

for this or for some of the indicated trends,

some factors can be identified and others suggested.

Since many specimens, especially the higher

strength ones, failed from the edges and sides,

edge effects appear to have been a limiting

factor. This is reasonable in view of the somewhat irregular nature of the edges, especially

on RF sputtered specimens. The fact that sur- faces were not flat also means that some

strength will be lost; however, this effect

should be relatively small.

The general tendency for previously ground

specimens to give higher sputtered strengths

than previously polished specimens is also rea- sonable. Since specimens tend to be more sensitive to a few isolated stress concentrations

than to many, the high density of depressions

on previously ground surfaces and edges would

favor higher strengths. The possibility that

sputtering may remove surface imperfections

from ground surfaces, but leave some work

hardened material on ground specimen which

might strengthen them must also be considered.

The observation of twins in sputtered sapphire and ruby bars and the fact that they generally showed no increase in strength is consistent with other results [19] indicating that

twins can severely limit strength in these materials. The heating in RF sputtering may aid

the growth of twins left from machining. The

absence of such heating in IB sputtering may

have resulted in 1 ruby section free of twins

and hence with higher strength. Such twinning

effects would also explain that lack of improvement of polycrystalline AI2O3 on sputtering.

The increase in the strength of the rutile

crystals may he due, at least in part, to heating.

Strengths approaching those obtained by sputtering have been observed by the authors as a

result of annealing similar specimens. The

darkening indicating reduction (and heating)

may also be a factor. This may also be a factor

in the possible increases in strength in the

ZrOa crystals. However, the fact that little or or no reduction was apparent in many areas

of strengths similar to reduced areas indicates

that reduction effects are at best small in ZrO-i.

Heating of the spinel crystals could cause precipitation of AI2O3, which would be expected to

strengthen them some at the expected temperatures [22]. However, the fact that nearly as high values were obtained in IB sputtering

where heating is not a problem indicates this

cannot be the major effect. It would also not

explain the increase observed in the fine grain

polycrystalline sninel since this material was

essentially stoichiometric. Heating to cause some annealing of these samples may have been

a factor there. These bodies do show substantial increases in strength on annealing, but only

at much higher temperatures.

The strength of RF sputtered glass appears

to have been limited by the cracks that often

developed. Heating was probably a major factor

in this since these cracks were not observed on

IB sputtered specimens. The latter clearly tended to have their strength increased ; however,

the irregular surface was probably a major factor in limiting these increases.

There is a possible trend in most of these in- creases. Recent work by the author (to be published) indicates that the strength of MgO and

AI2O.S are more frequently controlled by microplasticity than by flaws, while flaw mechanisms

appear to be controlling in MgAlo04, ZrOo, B^C,

SiC, and SisN,. These latter materials, along

with glass would then be expected to be more

sensitive to surface finish than MgO and AUO3.

The data does indicate such a trend for Zr02

crystals and single and polycrystalline spinel,

and for glass. The poorer quality of most of the

polycrystalline Zr02, SiC and Si-,Nt samples

are probably factors in their showing little effect.

5. Summary and Conclusions

Both radio frequency (RF) and ion beam

(IB) sputtering of a number of ceramics have

been investigated. RF sputtering normally produces a series of symmetrical depressions in

the surface. Pits left from mechanical finishing

and subsequently rounded (rapidly) by sputtering are a major cause of these depressions.

While these depressions are generally not as- sociated with grain boundaries, they do de- crease in size with decreasing grain size since

pits left from mechanical finishing normally

decrease in size with decreasing grain size.

Ground surfaces develop a high density of pits

and surfaces of finer finishes progressively

fewer depressions. However, the development of

depressions inside of others, on surfaces without pits, and their change due to heating during sputtering all suggest that dislocations may

be another source of these depressions. These

depressions are typically about 1 micron deep,

so sputtered surfaces are normally not too

rough. However, some polycrystalline materials

can develop a very rough topography, apparently due mostly to impurities, but differential

185
sputtering- between different plans may also be

a factor (e.g., in MgO). Soda lime glass develops crescent-shaped depressions that may originate from flaws.

Besides the apparent etching of flaws and

dislocations noted above, several other examples of etching (e.g., of impure areas, pores,

twins, and some grain boundaries) were observed. These were generally clearest after limited (e.g., 1-3 hrs RF sputtering). Besides

normal ion incidence, the lower accelerating

potential used in RF sputtering may also have

been favorable to such etching. Such etching

tends to be cumulative since once a feature develops it tends to be preserved even after its

original cause is gone. Longer sputtering nor- mally rounded such features sufficiently to make

them less distinct, or to obliterate them.

Ion beam sputtering at low angles of incidence generally developed similar topography.

However, depressions were less symmetrical,

and somewhat shallower. There was less evidence of etching, and much more regular sur- faces were obtained on those polycrystalline

bodies that gave rough surfaces in RF sputtering. IB sputtering at high angles of incidence may also be very useful for etching. Use

of reactive gases may also be very beneficial.

Either type of sputtering clearly improved

strengths of weaker samples (e.g., those ground

perpendicular to the tensile axis). Generally,

strengths comparable to those achieved by mechanical polishing were achieved by sputtering

any mechanically finished body. Annealing (e.g.,

1300 °C) after sputtering did not increase

strengths and may have decreased them some.

Limited trials showed no obvious degradation

of the strengths of flame or gas-polished specimens.

Results do suggest that some materials (e.g.,

glass, MgAloOj and possibly ZrOo) can have

their strengths increased by sputtering. Sapphire might also be strengthened, but this may

depend on the absence of twins or easily activated twin sources. However, this will generally

require improvement of edges separating

specimens so debris does not accumulate on the

sides. Ion beam sputtering appears to be favorable to producing better edges. Study of round

rods, or of plates tested biaxially near their

center would be good for further study.

The higher rates (e.g., about 1 ;u.m/hr) of RF

sputtering were favorable to it. However, IB

sputtering otherwise appears more desirable.

At low angles of beam incidence on the specimen, smoother topography is obtained, along

with better edge finish. Heating and attendant

problems (reduction of material and possible

expansion of twins and dislocations) are much

less. Further, since lower accelerating potentials

and at least near normal incidence are obtainable, in IB sputtering, it should be useful for

etching.

The author is extremely grateful to Drs.

Sigel, Anderson, Hockey, and Garvin for sputtering specimens. The aid of C. Herbert, J. Breen, and S. Slawson in preparing specimens

and of J. Robinson for electron microscopy is

gratefully acknowledged as are the initial discussions of the sputtering process with Dr. G.

Wehner.

6. References

[1] Bebrisch, R., Festkorperzerstaubung durch ionenbeschuss, Ereebnisse der Exaton Naturwissenchaften 35, 295 (1964).

[2] Weston, G. F., Cold Cathode Glow Discharge

Tubes, ILIFFE Books, Ltd., London, 1968.

[3] Wehner, G. K., Sputtering, Sciences and Technology, 32 (1968).

[4] Wehner, G. K., and Anderson, G. S., The nature

of physical sputtering, to be published in handbook of thin film technology, McGraw-Hill.

[5] Davidse, P. D., Theory and practice of RF sputtering. Symposium on the Deposition of Thin

Films by Sputtering, University of Rochester,

Rochester, New York, 1966.

[6] Tighe, N. J., Microstructure of fine-grain ceramics, Ultrafine-Grain Ceramics, Edited by U.

Burke, Reed, N., Weise, N., 109-33, (Syracuse

Univ. Press., Syracuse, New York, 1970).

[7] Paulas, M. and Reverchon, F., Study of the Parameters of the Ionic Bombardment of ferritesapplication of the thinning at almost grazing

Incidence of Porous materials. International

Symposium on Ionic Bombardment Theory and

Application, 324-35, (Gordon and Breach, New

York, 1962).

[8] Trillat, J. J., Ionic bombardment, a new method

for the study of surfaces, ibid., 13-51.

[9] Navez, M., Sella, C, Shaperot, D., Electron microscope study of the effects of ionic bombardment of glass, ibid., 339-53.

[10] Dugdale, R.A., and Ford, S.D., The etching of

alumina and fused silica by sputtering. Trans.

Brit. Ceramic Society 65, 165 (1966).

[11] Keig, G. A., and Haines, H. R., The Cathodic

etching of plutonium ceramics, Trans. Brit.

Ceramic Society 62, 363 (1963).

[12] Bierlein, T. K., Newkirk, H. W., Jr., and Mastel,

B., Etching of refractories and cermets by ion

bombardment, J. Am. Ceram. Soc. 41 (6) 196

(1958).

[13] Tarpenian, A., Electrochemical and ion bombardment etching of pyrolytic graphite, ibid., 532.

[14] Beecham, D., Sputter machining of piezoelectric

transducers, J. Appl. Phys. 40, 4357 (1969).

[15] Schroeder, J. B., Bashkin, S., and Nester, J. F., Ionic Polishing of Optical Surfaces, Applied ' Optics 5, 1031 (1966).

[16] Levy, P. W., Shaping or Figuring Ceramic Surfaces of Ion Beam Bombardment, these pro-

'

ceedings.

[17] Rice, R. W., The Effect of Grinding Direction on

Strength of Ceramics, these proceedings.

186

1
[18] Rice, R. W., Becher, P. F., and Schmidt, W. A.,

The strength of gas polished sapphire, these

proceedings.

[19] Becher, P. F., and Rice, R. W., Flame polishing of

flat oxide bars, these proceedings.

[20] Schmidt, P. H., Spencer, E. G., and Walters,

E. M., Ion milling of magnetic oxide platelets

for the removal of surface and near-surface

imperfections and defects, J. Appl. Phys. 41,

4740 (1970).

[21] Reswig, R. D., The use of hydrogen in cathodic

vacuum etching, Microstructures, 15 (1970),

Discussion

Schneider : Was there any evidence of chemical change or did you investigate this?

Rice: No, v^^e haven't had a chance to look

at this on these materials. I don't think there

is much evidence for change but this is an area

that v^ould be good to try and check in more detailed studies.

White: I was vi^ondering, have you checked

for any possible evidence of argon implanted in

the damaged layer?

Rice: Pardon?

White: What is the gas?

Rice: It's argon. The acceleration potential

of the RF sputtering is of the order of 2 kilovolts and about 6 kilovolts in the ion beam sputtering.

White: We have found argon in very sur- prising situations.

Rice: Argon stuffed in the surface v^^ould

probably provide a compressive layer w^hich

v^ould increase the strength. I feel that this

probably is not a major factor. For example,

annealing after sputtering did possibly decrease

some strengths a little, but there v^^ere no clear

changes of significance. Other factors such as

tv^rinning in sapphire may be more important

variables.

Perry: Did you notice any difference in

statistical scatter in the strengths of the glass ? Rice: Yes, I mentioned that particularly in

the case of glass we get some specimens that

are maybe three, even four times stronger than

we would get by mechanically finishing. That

is, you tend to have strength say from about the

mediam of that found for normal mechanically polished glass to substantially greater

strengths. This suggests that there is indeed a

potential for improving the strength of glass,

considerably higher than that achieved by mechanical finishing.

Green : Did you observe an increase of

strength of the glass while removing 1 to 4

mils of sputtering?

Rice: Yes.

Lange: In chemical polishing, the topography depends on the rate of the polishing. Did

you find any change in topograph by rate of

sputtering?

Rice : Really we haven't been able to investigate a significant difference in rate.

Heuer : If you have any dirt in your system,

you can get a very interesting but unwanted

topography. This is probably the most important factor in obtaining uneven topography.

Rice: This can be a factor since sputtering

is essentially an integrating or accumulative

process. However, I clearly observed that the

major factor in the uneven topography was the

prior surface finish and whether the specimen

was polycrystalline or not. The type of sputtering and the sputtering parameters also appear

to affect this.

Unidentified : How wide were the ion traps

that you originally showed in some of your ex- amples?

Rice: You mean the scribe mark? Oh, I don't recall, probably something like 25 microns

or so. It was done with just a hand held diamond scriber. The scale was on the figure, but

I don't recall explicitly the dimensions of the

scribe mark.

national bureau of STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

187

Computer Controlled Ionic Polishing of Optical Surfaces

James W. Douglass ^

Perkin-Elmer Corporation, 77 Danbury Road

So. Wilton, Connecticut 06897

Results of the development of an ionic polishing investigation are described.

During this investigation many controlled polishing experiments were performed. The

results of two of these experiments are presented. In the first of these, a 0.203 m

diameter optical flat of fused silica was polished to a diff'raction limited surface

quality of 95A (9.5 nm) rms. The second experiment described was the conversion of a 0.203 m diameter Cervit sphere into an f/5 paraboloid of diffraction limited quality

by using analytical expressions to generate the removal distribution. The ionic polishing

process was computer controlled for this experiment. The temporal stability and optical scatter performance of ionically polished surfaces has been determined to be as good

as or better than conventional surfaces.

Key words: Computer controlled process; ion beam technology; ionic polishing; optical

fabrication; optical scatter; sputtering.

1. Introduction

The limitations of modern optical systems

stem in many instances from the inordinately

long time required to produce the surface fig- ures necessary for their high performance. Furthermore, the recognized benefit of larger aperture systems makes the fabrication of even the

conventional surfaces more difficult.

For some time, Perkin-Elmer has been directing considerable effort toward developing

optical fabrication processes whose goal is to

fabricate optical components in an automatic

or semiautomatic fashion, and which require

high speed numerical analysis and control. A

goal of these programs is to eliminate much of

the "art" that has always been an important

requirement in conventional optical fabrication

and to replace it with techniques based on sound

technology.

The discovery by Meinel, Bashkin, and

Loomis [1]^ of the University of Arizona that

an ion beam can polish optical glass, indicated

the possibility that a controllable polishing technique might be developed. Perkin-Elmer en- tered into an exclusive agreement with the University of Arizona to explore this technique. As

a result, elforts have been underway in our

Process Development Laboratory to develop a

controlled polishing or figuring process based

on the use of ion beams. Some very early re- sults were published [2, 3] , but since that time

our results have been essentially unpublished.

This article discusses some of our developments

since that time. The approach will be to emphasize the results without going deeply into the

historical development or the extensive theories

which describe the phenomena contained within the ionic polishing process.

' Physicist, Group Leader Advanced Processes.

^ Figures in brackets indicate the literature references at the end of this paper.

2. Technical Discussion

2.1. Definition of Ion Polishing

Ionic polishing is the controlled removal of

surface material from optical elements by sputtering with an ion beam.

This definition does not limit the type of material to be polished or the characteristics, such

as mass, energy, and charge of the ion employed. The implicit point is that the process

must be very accurately controlled.

2.2. Relationship Between Ionic Polishing

and Sputtering

An atomic displacement occurs when an energetic particle strikes an atom or ion in a

solid and transfers sufficient energy to break

the chemical bond and force the atom to move

away from its original site. This is called a

knock-on collision. If the knock-on occurs near

the surface, the displaced atom may leave the

solid completely. Dislodging molecules in this

fashion is the process of sputtering. Sputtering

of dielectrics is generally produced with neutral plasmas created by radio frequency fields so that the alternate bombardment by oppositely charged particles precludes the buildup of a

net charge on the dielectric. The significance of

the discovery by Meinel, et al. is that ion beams

may be used instead of neutral plasmas to produce sputtering, and thus makes available techniques for the control of the process far beyond

that possible for plasmas. Controlled sputtering

by ion beams is referred to as ionic polishing.

Any net charge which may accumulate on the

surface of the dielectric does not produce fields

strong enough to produce any undesirable effects resulting from surface charge accumulation.

189
2.3. Brief Description of the Research and

Development Effort

A prototype ionic polishing laboratory was

set up with a Texas Nuclear linear accelerator

(30 — 300 KeV) with an RF ion source. The

beam obtained from this accelerator was magnetically analyzed and focussed onto the sur- face of the test specimen in a sample chamber.

This beam was electrostatically deflected in a rectangualr pattern, so as to produce rectangular areas of polishing in order to simplify determinations of sputtered volume. In this way,

a parametric study was conducted to determine

the sputter yield, i.e., the ratio of the number

of sputtered atoms to the number of incident

ions, for various optical materials as a function of the species of incident ion, its mass,

charge state, energy, and angle of incidence to

the test sample. This investigation was principally limited to the following optical materials: fused silica, BK-7, ULE, and Cervit. The

ions were primarily those of the noble gases,

although some experiments were performed

with H% HH% and HHH\

These experiments were conducted on oneinch diameter samples, and sputtered volume

measurements were performed interferometrically. A typical interferogram of such a sample is shown in figure 1. Experiments were performed on polished and ground surfaces. The

results of this investigation indicated sputter

yields in the range of 1 — 3 and that these

Figure 1. Interferogram showing uniform removal in a rectangular pattern.

values strongly depend on the species of ion,

but are only weakly dependent on the ion energy. The sputter yield varies approximately

as the inverse cosine of the angle of incidence

in the range of 0 to 60°. Ground surfaces may

be polished, but long polishing times are re- quired for beam currents on the order of lOO^aA.

Consequently, all subsequent work has been on

polished surfaces.

2.4. Description of the Ionic Polishing and

Data Handling Systems

The current ionic polishing system is shown

in figure 2. The performance of this system

differs from the earlier prototype in that the

ion beam is rapidly raster scanned over the test

sample. The effective velocity of the beam spot

is controlled by controlling the beam spot dwell

VABIAN S3tD/t COMPUTER

D/A

CONVERT8RS

DEFUE

AHPL

CTKW

tram

d

C3

QUAD

LENS

lCCKt£RAT(»

SAMPUE CHAMBER

Figure 2. Block diagram of ionic polishing

system.

INTERFEROGRAM

OF SAMPLE

DATA REDUCTION

WITH FLYING SPOT SCANNER

REMOVAL DISTIUBUTION CONVERTED TO

POSITIONAL AND

DWELL TIME PAPER TAPE

VARLVN 620/i COMPUTER

Figure 3. Block diagram of data handling system.

190
time at discrete locations on the piece separated

by intervals which are small compared with the

beam spot diameter. The positional and dwell

time controls are effected by an on-line Varian

620/i computer. Provided the beam spot diameter is small compared with the dimensions of

the highest spatial frequency of interest on the

sample, one may linearly relate the dwell time

at any location to the desired removal at that

location and thus produce controlled figuring.

In practice, the computer core is stored with

coordinate and dwell time information. The

computer cycles through this data as described, as long as is required to figure the sample. Both

rectangular and polar coordinate systems are used to match the geometry of the removal distributions on the sample.

The data handling system is shown in figure

3. Due to extensive use of computers, a great

deal of special data processing is possible ; i.e.,

deconvolution of data and coordinate transformations, as well as the generation of removal

distributions, based on purely analytical expressions.

3. Experimental Results

3.1. Figuring Experiments on Optical Flats

Many experiments have been performed

whose goal was to polish optical flats to a high

degree of perfection. The results of one such

figuring experiment are shown in figure 4. This

experiment was designed to figure a 0.203 m

diameter fused silica flat whose peak-to-peak

and RMS departures are 1100A. (110 nm) and

180A (18 nm) respectively. The final surface is

of diffraction limited quality with peak-to-peak

and RMS departures of 700A (70 nm) and

95A (9,5 nm) respectively. This experiment

was performed with a 120 /xA ion beam of 150

KeV Ar+. The removal profile was determined

from the initial interferogram.

3.2. Figuring Experiments on Spheres

Experiments have also been performed on

non flat samples, i.e., spheres and aspheres.

One particularly notable experiment was the

automatic conversion of a spherical surface into a paraboloid. In this experiment a 0.203 m diameter sample of Cervit was ground and polished into a sphere with a radius of curvature

of 2.03 m. A removal distribution was generated analytically which represented the minimum volume conversion into an f/5 paraboloidal surface. This removal distribution is shown

graphically in figure 5. This figure was made

by time exposing an x-y plot of the removal

distribution with an oscilloscope as it was output by the Varian computer. The degree of ex- posure at each location is a measure of the

dwell time at that location. Figure 6 shows the

Before

/r

V 1 I!

Figure 4. Before and after interferograms of a

0.203 m Diameter. Fused silica optical flat figured by ionic polishing to a final surface

figure of 95A (9.5 nm) RMS.

Figure 5. Time exposure of positional and dwell time

data used to convert a cervit sphere into a diffraction

limited f/5 paraboloid.

191
After

Figure 6. Before and after interferograms of a

cervit sphere converted to a -paraboloid. Computer generation and control of ionic polishing

system data was used. Interferograms were

taken in a spherical wavefront interferometer.

interferograms of this sample before and after

figuring. These interferograms were made in

a spherical wavefront multiple beam interferometer [4]. Consequently, the initial sphere

shows straight and equally spaced fringes while

the paraboloid exhibits the characteristic "hole

and roll" interferogram. The final paraboloid is

also of diffraction limited quality having a maximum departure of 600A (60 nm) from the

desired paraboloid. This amount of departure

corresponds roughly to 130A (13 nm) RMS.

This experiment was performed with a 30 KeV

Ar* ion beam and required approximately 18

hours to perform.

3.3. Stability of lonically Polished Surfaces

Subsequent interferometric measurements of

the optical flat of Section 3.1 over a 12 month

period show that this ionically polished surface

is as stable as conventionally polished surfaces

in that no change in the shape of the surface

was detected. The values of the surface stress

induced by the implantation of the ions have

been determined for ULE and Cervit when irradiated with 30 KeV Ar+. These measurements

indicate stress levels too low to cause any significant strains in modern optical components

and are generally no larger than the bulk

stresses commonly found in optical materials.

3.4. Surface Cosmetic Finish of lonically

Polished Surfaces

Optical scatter measurements have been performed on ionically polished surfaces. These

measurements indicate that these surfaces are

at least as good as conventionally polished surfaces, and in some cases better. These measurements indicate an RMS surface roughness in

the range of 10A — 50A (Inm — 5 nm). Recent unpublished work at Perkin-Elmer indicates that ion beam polishing may be used to

improve the surface smoothness of sapphire

laser optics.

4. Summary and Conclusions

Ionic Polishing has been applied to figure optical elements with automatic control by a computer. In this way, optical flats and image forming elements have been produced which are of

diffraction limited quality. The stability and

optical performance of these surfaces is as good,

or better than, the initial surface.

Ionic Polishing has been shown to be a useful process for the production of conventional

optical components. Since the process is capable

of producing any arbitrary surface, it is ex- pected that a new generation of optical systems

may result due to the manufacturing capabilities which were not available with conventional

polishing techniques.

5. References

[1] Meinel, A. B., Bashkin, S., and Loomis, D. A., Appl.

Opt. 4, 1674 (1965).

[2] Schroeder, J. B., Bashkin, S., and Nester, J. F., Appl. Opt. 5, 1031 (1966).

[3] Schroeder, J. B., and Dieselman, H. D., J. Appl.

Phys. 40, 2559 (1969).

[4] Polster, H. D., Appl. Opt. 8, 522 (1969).

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceram ic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

192
Arc, Laser, and Electron Beam Machining of Ceramics

R. W. Rice

U. S. Naval Research Laboratory

Washington, D.C. 20390

The various types of arc, laser, and electron beam machining, the considerations

that are required in their operations, and some of the applications for which they

have been tried or used, are briefly reviewed. The most general consideration is whether

thermal shock and stress cracking will be a problem. The environmental and material

parameters, which vary substantially with the types of these different machining

processes, can also be quite important. Meeting the requirements of all these considerations does limit the versatility and applicability of these techniques. Nevertheless, they

deserve consideration for a variety of applications.

Key words: Arc machining; cutting; drilling; electron beam machining; laser machining; shaping; thermal shock control.

1 . Introduction

Speakers who have been actively investigating arc, laser, or electron beam machining of

ceramics v^^ere not able to attend this conference. Since these technologies offer a number

of opportunities for machining ceramics, and

since the author has had some experience v^ith

these technologies, it was felt worthwhile to re- view these techniques. This review is not meant

as a detailed analysis of these techniques, nor

an exhaustive compilation of their applications.

Rather, it is a brief introduction to these techniques and a summary of some of the trials or

uses made of them.

2. Types of Arc, Laser, and Electron

Beam Machining

Much of the scope of arc, laser, and electron

beam machining, shown in table 1, can be rec- ognized from the fact that these techniques

can be used for welding. From metals technology, it is clear that welding techniques are

easily converted to cutting techniques, which is

one of the major applications of these techniques. It has a major advantage over abrasive

cutting of hard materials in that it does not depend on hardness. High melting (or vaporization) temperatures, which are usually associated with high hardness, will limit the rate of

such cutting ; but generally not near as much as

in abrasive cutting. Cutting by these techniques

also has the advantage of being nearly as versatile as a jig saw on wood. While the tolerances of such cutting usually restrict it to

rough or general shaping, it nonetheless has

considerable potential for many hard materials.

Oxidation-assisted cutting of some nonoxides

with an impinging oxygen stream as done with

metals (e.g., Ti) may also have some promise.

Drilling is of course a special case of cutting.

The vaporization and fracture behavior of

many ceramics have important effects on these

machining operations. Many ceramics undergo

considerable vaporization near their melting

point, and many vaporize without any melting.

This can increase the tolerances of such machining since the irregular layer of solidified

melt is reduced or eliminated. For example,

running an electron beam over SiC which does

not melt leaves a groove very similar to that

achieved with an air-driven abrasive stream.

This is important in drilling fine holes since

there is little or no melt to close the hole. Vaporization generally absorbs substantially more

energy than melting, and so takes more power

or may reduce the thickness that can be handled, but this does not appear to be a major

limitation for many applications in view of

powers available for some of these techniques.

Minimizing or preventing fracture from therTable 1. Types of arc, laser, and electron beam ceramic machining

Processes involving:

Complete penetration : Incomplete penetration

1. Cutting 1. Cut and fracture

2. Shaping, by cutting pieces out 2. Score for later fracturing

3. Drilling—complete holes 3. Drilling—incomplete holes

4. Shaping—vary surface topography by melting

or vaporizing areas

193
mal shock is a major consideration in these

techniques, which is discussed below. However,

it should be noted that in some cases thermal

shock fracture can be utilized to advantage.

Also, these techniques may be used to score materials for subsequent fracture along the

scored lines.

3. Considerations in Arc, Laser, and

Electron Beam Machining

A number of factors must be evaluated for

each of the different techniques and their variations. The most general and important is fracture due to thermal shock. This will decrease in

importance as the thermal shock resistance of

materials considered increases, and as the size

of the part to be machined decreases. Thus it may be possible to machine fairly large pieces

of SiOo and possibly SisN j, while smaller pieces

of SisNi and SiC can probably be machined

without any precautions against thermal shock.

In some cases, e.g., rough cutting, some cracking along the edges may be tolerable. However,

in many cases, thermal shock will be a problem

to contend with. This can generally be solved

by supplemental heating, as used in welding of

ceramics The procedure is to simply provide additional heating in the region of cutting or drilling (or welding) so that the net thermal

gradient is below that which will cause fracture. Supplementary heating of 1000 °C ±

300 °C is often adequate.

A variety of considerations pertinent to each

of the three techniques is listed in table 2. Lasers can of course be used in essentially any

atmosphere. Arcs are limited in vacuum, but

can operate in a variety of atmospheres, though

penetration is dependent on the atmosphere

(e.g., greater penetration in He than A). The

conventional and glow discharge electron beams

are limited to vacuum, as listed. Electron beams

' Figures in brackets indicate the literature references at the end

of this paper.

that are generated by conventional vacuum

electron guns, but ejected into the air (about

1.8 cm working distance) remove the vacuum

restriction from the part, but at the expense

of capital cost (by a factor of two or more).

Material coupling is most limited with arc techniques, since a conductor is required. However,

heating can make many oxides conductive. With

lasers, the primary question is the absorptivity,

which depends on wavelength. Since CO. lasers

offer the greatest power and efficiency, absorptivity at their wavelength (10.6 /xm) is of greatest importance. Materials with metallic-like reflection (e.g., TaC, TiN, etc.,) will generally

have poor absorbance until some heating has

occurred. Thus, starting may be slow unless the

supplementary heating (if used to prevent

thermal shock) changes this enough. Limited

oxidation, or other coatings of the area of the

beam will start in (or follow) can make important differences. Electron beam coupling is the

most general. However, if a material is not a

conductor, charge can build up and deflect the

beam away. This has not been a problem in

welding where supplementary heating is used.

The elimination of the problem with supplementary heating apparently results for either

enhanced conductivity or enhanced electron

emission (with the supplementary heater winding collecting the emitted electrons), or both.

While the relative roles of melting and vaporization depend most on the material, they also

vary a fair amount with the atmosphere and

amount of surface superheating. Arcs are usually used in an atmosphere (generally inert) and

heat a considerable depth, so melting is favored. (Note this is true of both the work piece

and electrode. To avoid the melting limitations

of W, other electrodes such as C, or TaC, can

be used.) Electron beam penetration favors

melting, but when used in vacuum, this will

enhance vaporization. Recent work indicates

that laser radiation has limited penetration,

Table 2. Arc, laser, and electron beam machining considerations

A. Technique

B. Mode

C. Operational

Considerations

1. Atmosphere

2. Material coupling

parameters

3. Material removal

4. Cut width

Arc

Conventional,

other—e.g. plasma

Various, but

affects

penetration

Conductivity

Favors melting

Wide

Laser

Continuous, pulse Conventional

electron gun

Electron beam

Glow discharge

Any

Absorptivity

(vs X and temp)

Favors

vaporization

Generally fine

Medium vacuum Low vacuum

(10-' -10-^ torr) (0.1 torr)

Electron discharge

Favors melting

Fine

Air

ejected

beam

Any

194
which leads to greater surface heating and re- sultant vaporization than electron beams [2].

Another consideration is the cut width. This

is generally widest with arcs, and finest for

electron beams, with lasers intermediate, and

often substantially finer than arc. In all cases,

there will be a substantial bevel on the front

edge (or conical entry to holes) and some taper

to the edge or hole wall.

4. Applications of Arc, Laser, and

Electron Beam Machining

Almost no effort has been made to utilize arc machining. However, one firm apparently

tried arc drilling of small ALOs parts (heated

to provide conductivity) with some degree of

success. Many laboratories have probably used a laser

to "drill" one or more holes in glass parts out

of curiosity or for some special laboratory need,

as in the author's laboratory. Several have also

probably drilled holes in more refractory materials (e.g., 1 mm x 10 mm holes in ruby

[3] ) . The most extensive application of laser

machining is the scoring of ALO., electronic substrate materials for subsequent fracturing to

size [4] . There may be some adverse effects on

substrate strength in comparison to mechanically cut [5] ones, but this system is currently

used in industry.

Electron beams from glow discharges have

been used to drill 0.006 in (0.152 mm) holes in

diamond, and to "machine" a spiral groove in

0.5 in (1.27 cm) diameter fused SiO, rod [6].

The latter grooving resulted from vaporization

and melt flow from the spiral path of the beam.

These were apparently done without supplementary heating. Schumacher [7] et al., have

demonstrated that pieces of rock over 15 cm

thick and several cm in cross section can be

split by penetrating 5 cm into them with an

air ejected electron beam operating at 9 KW

with 150 KV.

5. Summary and Conclusions

Arc, laser, and electron beam methods of

machining ceramics offer considerable versatility in cutting, drilling, and shaping ceramics.

Some of this versatility is compromised by

problems of thermal shock fracturing and the

heating techniques often needed to minimize

this, as well as by the tolerances and edges

achieved. Nonetheless, these processes offer considerable opportunity, especially for applications where the minimization of thermal shock

problems are not needed, or are not a significant encumberance. They should, therefore, be

considered because of the speed, cost, or size

advantages they can offer.

6. References

[1] Rice, R. W., Welding of Ceramics, U.S. Naval Research Laboratory Report 7085, July (1970).

[2] Pahlitzsch, G., and Visser, A., Materials Removal

Processes in Laser and Electron Beam Machining, p. 335-57, in Electron and Ion Beam Science

and Technology, R. A. Bakish, ed., (The Electrochemical Society, New York, 1968).

[3] Bernal, E. G., and Ready, J. F., Lasers in High

Temperature Research, in Proceedings of the

3rd Int. Symposium on High Temperature Technology (Butterworths, 113, London, 1967).

[4] Lumley, R. M. Controlled separation of brittle

materials using a laser, Am. Ceram. Soc. Bull.

48(9), 850-54, (1969).

[5] Lo, W. C, Edge effect on the modulus of rupture

of ceramic substrates, this proceedings.

[6] Parisi, A. J., Electron-Beam Power Ready to Topple Cost Barrier, Product Eng., pp. 22-24, Nov.

18, (1969).

[7] Schumacher, B. W., Electron Beam Cutting of

Rocks and Concrete, p. 447, in Electron and Ion

Beam Science and Technology, R. A. Bakish ed.,

(The Electrochemical Soc, New York, 1968).

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-^, 1970, Gaithersburg, Md. (Issued May 1972)

195

The Techniques and Mechanisms of Chemical, Electrochemical, and Electrical

Discharge Machining of Ceramic Materials

D. William Lee and George Feick

Arthur D. Little, Inc., Cambridge, Massachusetts 02140

Electrochemical Machining (ECM), Electrical Discharge Machining (EDM), and

Chemical Machining (CHM) techniques have not been extensively used as a material

removal process for ceramics. The fact that these processes are not dependent upon

the hardness of the workpiece should make them attractive for the shaping of ceramics.

However, ECM and EDM techniques are limited to material with reasonably good

electrical conductivity while CHM may be limited by the availability of effective

etchants and the nature of the ceramic itself. The theory and techniques of chemical, electrochemical and electrical discharge

processes will be reviewed. The variables and important parameters of the material

removal processes will be discussed, including etchants and mask techniques for

chemical milling; current-voltage characteristics, dielectric fluids, electrode materials,

etc., for EDM; electrolytes, gap eff'ects, current density, etc., for ECM. The available

information and experiences developed on metals and alloys will be used to examine

the applications and limitations of these processes to ceramic materials. Where possible

the effects of the material removal process on the surface condition, microstructure, and

subsequent material properties will be pointed out. Finally, comments will be made on

the technical and economic feasibility of ECM, EDM, and CHM processes for the

shaping, cutting, and finishing of crystalline ceramics.

Key words: Ceramic materials; chemical machining (CHM); electrical discharge

machining (EDM) ; electrochemical machining (ECM) ; intermetallic compounds.

1 . Introduction

Although electrochemical machining, electrical discharge machining and chemical machining have not been extensively used as material

removal processes on ceramic materials, these

nonconventional machining processes are potentially attractive for three reasons: First,

the processes are not dependent upon the hardness of the workpiece, secondly, little or no mechanical stress is introduced into the workpiece,

and thirdly, there is the possibility that these

processes could machine configurations and with

tolerances not possible with grinding or chipmaking processes.

Basically, one would like to know if these

nonconventional processes as related to ceramic

materials are technically feasible, economically

competitive, and finally, what effect the material removal process has on the properties of

the machined ceramic body. It is not proposed

that this review will provide the answer to

these questions, since it is immediately apparent that there is too little information from

which to base any firm decisions. Rather, an

attempt will be made to describe the process of

electrochemical machining, chemical machining,

and electrical discharge machining in a general

way and to point out previous experience which

may shed some light on technical problems and

the effects of the machining process on the

properties of the material. The analysis of the

economic competitiveness of these processes

can at best be only qualitative. This review fo- cuses very heavily on the material removal

process for crystallized materials and the ways

in which these processes have been utilized.

Machining and finishing of glass will only be

touched on as it is pertinent to the overall machine process.

The practical application of these processes

has in many cases been substantial. ECM has

been used as a special production method for

the past ten years particularly in broaching,

trepanning, hole drilling, and contouring of

hard tough alloys. CHM has been available for

commercial production for thin complex metal

parts, engraving and weight reduction of aircraft parts for 15 years. EDM is a well established process for producing holes and cavities

in tough materials, superalloys, refractory

metals and cemented tungsten carbides since

the mid '50's.

In table 1 some of the parameters of these

processes are compared. ECM is a low voltagehigh current process requiring greater power

than other processes. Both ECM and EDM are

close contact operations with gap distances of a few thousandths of an inch, resulting in close

tolerances. In table 2 a comparison is made of

the effects on machined part. The high power

consumption of the ECM process is reflected in

the high metal removal rate which is comparable to more conventional milling. All processes

can produce good surface finish. Surface damage is of particular concern with EDM, while

the damage incurred by ECM and CHM can be

controlled by proper procedures.

These processes are relatively proven production techniques for metal while for ceramics

439-921 O - 72 - 14 197
Table 1. Physical parameters

ECM CHM EDM

Potential (volts) 10 45

Current (amps) 10,000 DC — 60 Pulsed DC

Power (watts) 100,000 2,700

Gap (meterxlO^) 0.0203 0.0254

Medium Liquid electrolyte Liquid chemical Liquid dielectric

Table 2. Effects on machined part

Electrochemical Chemical Electric Conventional

discharge milling

(ECM) (CHM) (EDM) ^**^ Removal rates mVsec(xlO') 264 0.42 133 111

Dimensional control m(xlO') 0.0051 0.0051 .0051 0.0051

Surface finish m(xlO ') 10-250 50-250 25-1200 50-500

Depth of possible damage 0.005 0.005 0.0125 0.0025

m(xlO')

** Stagger tooth milling of 4340 steel.

no commercial or production processes are used

with the exception of EDM of cemented carbides. Since ECM and EDM are applicable to

materials with relatively high electrical conductivity, these techniques are limited by and

large to intermetallic compounds of the carbides, borides, nitrides, silicides, etc., and these

materials do not contribute any significant

commercial dollar volume. Chemical milling,

although more generally applicable, is not widely used, principally because of the slow material removal rate and the difficulties of masking.

Chemical etching and polishing of semiconductors are not covered in this review. References [39] and [40] provide a good discussion

of the etchant and procedure for II-VI and

III-V semiconductor material.

2. Electric Discharge Machining

2.1. Description of the Process

The process was first developed as a practical method of metal working by Lazarenko and

others in Russia during the second world war.

Since that time it has come into widespread

use for difficult and specialized machining operations. It is used especially for shaping hardened steels and high alloys which are difficult

to work by conventional methods, and for special operations such as fine drilling and forming complex shapes. The metallurgical applications of EDM have been extensively reviewed

[1, 2, 3, 4]\ and will not be considered here.

Its applications to ceramic materials have been

' Figures in brackets indicate the literature references at the end

of this paper.

Control 1 ing Servomechanisin

Tool Holder

and Feed

Workpiece -^Pmp Filter

Figure 1. Schematic of a Typical EDM Unit.

relatively few and are limited to simple shapes

such as bending test specimens.

In electric discharge machining (EDM) a

series of electric sparks are produced between

the workpiece (usually the anode) and a shaped

tool which is separated from the work by a

thin layer of flowing dielectric liquid (fig. 1).

The work (and also the tool) are melted and

vaporized by the intense local heating and the

resulting erosion product or swarf appears as a dispersion of fine metal particles which are

carried away by the flowing dielectric. As the

erosion proceeds, the spark gap is kept constant

by a servomechanism which advances the tool

in response to changes in the gap voltage.

A number of phenomenological descriptions

have been proposed for the material removal

process. Some mechanisms rely extensively on

thermal evaporation, others site shock waves

produced by cavitation of the collapsing vapor

bubble. Most evidence, however, points to a thermochemical approach. A narrow channel of

the dielectric is ionized and the spark discharge

occurs as an electron avalanche toward the

198
anode. The liquid dielectric confines the discharge to a very narrow channel and maintains

a very high current density. The avalanche

heats the anode causing some evaporation and

melting. As the impact area increases and cur- rent density falls, the amount of material re- moved by evaporation reduces and the melting

increases as a result of thermal lag. The high

temperature gradient results in substantial

thermally induced stresses in the subsurface

areas. The collapse of the dielectric vapor bubble probably accounts for ejecting some molten

material from the crater.

2.2. Effect of Process Variables

The choice of power supply is a major factor

in the success of the EDM process. There are

three types of spark generators. The particular

choice is usually a compromise in speed, accuracy, finish, tool wear, damage reliability and

cost. Rotary generators provide long pulses that

erode more like an arc and allow large powers

to be applied and result in high removal rates.

Rotary generators are being replaced by transistor-switched generators. The latter maintain

the high cutting rate of the rotary generator

while yielding much lower tool wear. For ce- ramic applications the resistance-capacitance

(R-C) circuit is most generally applicable (fig.

2). This is a form of relaxation oscillator in

which a capacitor is charged through a resistance from a dc power supply. The working gap

is connected in parallel with the condenser.

When the condenser voltage (Vd) becomes sufficiently high the gap breaks down, and the en- ergy stored in the condenser is released to the

spark.

The capacitor C is recharged and the cycle

is repeated. The time constant RC must be such

that the gap voltage does not become so large

that arcing occurs. Machining speed is partially a function of spark energy and repetition rate.

If the discharge voltage is small compared to

the charging voltage, the repetition rate will

be high but the energy will be low.

This is a simple and low cost circuit which

is at the same time rugged and reliable. In general, R-C circuits operate best at relatively

high frequency and low machining rates which

are the conditions required for ceramics.

The cavity in the work exceeds the size of

the tool by an amount corresponding to the

width of the spark gap. This is known as overcut and depends both on spark energy and voltage. Low voltage, low energy sparks give minimum overcut.

The use of a liquid dielectric is essential in

the EDM process. The liquid not only carries

away the swarf but confines the spark energy

to a small area. In air or gas the spark energy

is spread over a wider area, and little erosion

D.C. Source

'V'Volts

+

oTool i_V Controlled

Gap'd' Average fn Voltage Workpiece "

SPARK GENERATION BY RELAXATION (RC) CIRCUIT

Chargi ng Vol tage

TTHT

SIMPLIFIED FORM OF GAP VOLTAGE DURING RC SPARKING

Figure 2. Simplified form of gap voltage

during RC sparking.

takes place. Hydrocarbon liquids such as kero- sene or transformer oil are the most commonly

used dielectrics, and have the advantage of preventing rust and corrosion. Water and aqueous

solutions of polyglycols have been found [1]

to give improved machining rates and less tool

wear, but they have not yet come into general

use.The erosion rate of both the work and tool

depend not only on the spark characteristics

and on the dielectric liquid but also on the properties of the materials from which they are made. In general, materials of high melting

point and high thermal conductivity are least

eroded and therefore most suitable for tool materials (table 3). Tungsten, graphite, copper

and brass have been recommended. Since pure

tungsten is difllicult to machine, a tungsten-silver composite is sometimes used. For ordinary

shop work, brass, copper and cast iron are preferred because of their availability and easy

machinability. A low-melting alloy of tin and

zinc is sometimes used for complex tools beTable 3. EDM erosion rates

Lead 7.0

Zinc 4.5 Aluminum 1.3

Brass 1.1

Copper 1.0

Iron 0.45

Molybdenum 0.39

Tantalum 0.30

Tungsten carbide 0.22

Tungsten 0.19

199
cause it can be readily cast to shape and is

reusable.

Samsonov and Mukha [5, 6] have studied the

behavior of many electrode materials and have

correlated their results on the basis of electronic structure. Since the electronic structure

also correlates with melting point and thermal

conductivity, their work leads to the same general conclusions as the melting-point correlation [1].

2.3. Applications to Ceramics

a. Advantages and Limitations of E.D.M. for Ceramics

With minor exceptions (see 2.3.c. below) the

EDM process can be apialied only to electrically

conductive materials. There is no well defined

limit of conductance to which spark machining

is limited. Increasing the resistivity of the

work, however, is equivalent to putting a re- sistance in series with the spark gap, which

limits the current and absorbs energy which

would otherwise be available to the spark. Walson et al., [4a] report cutting germanium crystals having a resistance of 31 ohm cm with low

spark energies of 10-^ to 10-* joule. It is probable that resistivities much greater than this

would seriously limit the cutting process. For

instance, it has taken 10 hr to cut 1 in (2.5cm)

boule of 300 ohm cm silicon. This requirement

eliminates the bulk of the conventional ceramic

and vitreous materials and limits the application to conductive intermetallic compounds such

as the borides, carbides, nitrides and silicides.

Examples which will be discussed in further

detail include zirconium carbide and nitride,

chromium carbide, molybdenum silicide and

cemented carbide cutting tools.

A major advantage of the EDM process is

that the removal of material is not influenced

by the hardness of the work. The removal rate

is affected by the melting and boiling point of

the work material, and very refractory materials such as ZrC and ZrBa are appreciably

Table 4.

Material Spark energy Depth of damage Tool

(Joules) (Meters X 10 "*)

Copper 5 X 10-' 1100 Rotating

wheel

5 X 10-* 300 Rotating

wheel

5 X 10^ 80 Wire

Fe-3.2%Si 0.6 150 Wheel

5 X 10-* <2 Wire

Tungsten 5 X 10-* <io Wheel

CO/WC 3.6 X 10-= 200 (Hardness) 7 CO/WC 3.6 X 10-= 10-50 (Cracks) 7

slower cutting than less refractory substances.

We find that removal rates of 10-* -10-= in

Vmin (10-^-10-8 mVs) about the optimum.

The energy available from the spark, however,

is enough to erode any known material including graphite and tungsten compounds.

A second advantage is that the tool does not

contact the work and therefore exerts no force

on it. This is particularly valuable in the cutting

of very brittle materials such as molybdenum

silicide. The dielectric fluid exerts some pres- sure on the work, especially if it is pumped

through a hollow tool. This hydrostatic pressure

is appreciable only for large pieces, however,

and is not ordinarily a factor in shaping cer- amics.

The major drawback to the application of

EDM to ceramic-type materials lies in the nature of the surface produced by the process.

This surface consists of a layer of resolidified

melt whose topography consists of small over- lapping craters left by the spark together with

solidified drops and splashes of liquid. There is

also a region of heat-shocked but unmelted material. In the case of brittle materials, both

these surface zones contain an abundance of

flaws and microcracks which seriously impair

the strength of the machined piece (fig. 3).

The extent of the cracking of the surface of

large grained polycrystalline MoSi, depends on

the spark energy. Energies above lO-^J produced severely rough and extremely cracked

surfaces. We have found that slower cuts at

lower spark energies are required for preparation of rough crystals of borides, carbides and

silicides.

The nature of the spark-eroded surface is

shown in figure 4, which is a scanning electron

micrograph of a surface of zone-melted molybdenum disilicide (MoSij) which was cut off

with a 0.015 in (0.035cm) copper wire tool on

a "Servomet" machine made by Metals Research Ltd., of Cambridge, England. This machine employs an R-C circuit and the surface

200
Figure 3. Macrostructure of MoSii cut with EDM

(a) 1.5X10 'J (b) 2.5X10-'J

(c) 6.5X10-'J

(d) 1.3X10-'J

shown represents a moderately heavy cut (225

V, l|u,f, 108 ohms). The surface shows the pitting

caused by the sparks together with droplets of

resolidified material. The diameter of the piece

was 0.235 in (6mm) and the time for cutolT

was 14 min. The overcut was about 0.0025 in

(0.06mm).

As the energy of the spark discharge is increased, the depth of the damaged layer becomes greater (table 4). This fact limits the

spark energies and thus the cutting rates for

ceramics to relatively low values as compared

to those used for ductile metals.

Spark damage to germanium crystals was

studied by Walson et al., [4a]. At spark energies less than about 10-* joules, no surface or subsurface damage could be detected by polishing and etching. Above 10"^ joules, however,

cracks and dislocations were abundant.

b. Refractory Intermetallic Compounds

Lanin et al., [8] studied the bending strength

of ZrC specimens cut by the EDM process,

some of which were subsequently lapped with

boron carbide powder (fig. 5). They found that

the EDM cut surfaces were extensively cracked,

especially at higher machining rates (corresponding to greater spark energies) , and therefore very weak. Removal of about 50/aM of the

damaged material from the surface by lapping

increased the bending strength by as much as 50 percent in some cases. Similar results were obtained with ZrN by

Yanchur et al., [9] who found that the bending

strength of spark-cut single crystals could be

improved by about 30 percent by lapping with

boron nitride. Even higher strength was ex- hibited by polycrystalline samples with undisturbed surfaces made by pressing and sintering.

In our laboratories we have utilized EDM

techniques for many years to cut and shape

single crystals of carbides and diboride (ZrC

and ZrBo). Specimens for deformation studies

were rough cut using a tinned copper wire and

spark energies of 10^' — lO^V. In order to re- move possible surface damage, specimens were

chemically polished to remove up to 25 microns

of the surface before testing.

c. Cemented Carbides

Artamonov et al., [10] report that surface

damage to cemented chromium carbide was

caused by loss of bonding metal from between

201
Figure 4. Scanning Electron Micrograph of Surface

of MoSit after EDM Cutting at 225 Volts, 108 Ohms,

IfiF. 625X.

the carbide grains. Mechanical weakening can

be minimized by using short, low-energy pulses

with water as the dielectric.

The EDM process has been of some interest

for sharpening and shaping cemented carbide

cutting tools. Matsuyama and Fukatsu [11] re- port that in some cases, the electrically sharpened tools containing WC and TiC will give

better life than those sharpened by diamond

grinding. They attribute the improved life to a

new intermetallic phase, detectable by X-ray

diffraction, which may he a solid solution of

WC, TiC and TiO. Spark energies of 10-' to lO V

are sufficient to form the new phase, however,

if the spark energy is too high the heat-affected

layers become extensively cracked, and tool life

is adversely affected. Based on hardness measurements, the authors report effect of the order

of 200/i, below the surface. (See table 4.)

The surface layers on electrically machined

straight WC-Co were examined by Filimonenko et al., [12] who find that the original hexagonal WC phase is changed to a mixture of

W2C and a face-centered cubic modification of

WC. The relative amount of the phases varied

inversely with the spark energy and the specific

surface reaction depends on the cobalt content.

d. Nonconducting Ceramics

It is usually considered that non-conducting

ceramics are not amenable to the EDM process.

Saito et al. [13], however, report that the sur- face of alumina and similar materials may be

etched by covering the surface with a thin

layer of electrolyte (e.g., NaOH, HCl or HF)

in a container of the same liquid in which an

inert electrode is immersed. Electrical discharges are generated at the edge of a second

S/S„ .8

EDM + MECHANICAL POLISH

(50 M Removed)

— 18t Porosity

3% Porosity

.5 1.7

Figure 5. Cutting rate—meters '/sec X 10 '

202
electrode which is in contact with the wet ce- ramic surface and is moved about to give a uniform effect. However, since the rate is slow

and the material removal is limited to the edges

of the inert electrode, it is doubtful if the process can be utilized as a significant material re- moval process other than to prepare ceramic

samples for microscopic examination.

3. Electrochemical Machining

3.1. Description and Scope of the Process

Controlled electrolytic removal of the metal

was first practiced as a method of preparing

polished surfaces by Jacquet in 1935, although

ECM was first proposed by a Russian, W.

Gusseff in a 1929 patent. It was not until about

1950 that any general use was made of the

technique.

In the electrochemical machining (ECM)

process, the tool and the work are separated by

a thin layer of flowing electrolytes—such as an

aqueous solution of sodium chloride, nitrate or

chlorate—while a heavy DC current is passed

between the two (fig. 6). The workpiece is

made the anode and the tool is the cathode.

Metal is dissolved electrolytically at the anode

and is usually precipitated as a hydroxide

sludge in the electrolyte by hydroxyl ions genCONSTANT

Figure 6. Schematic Diagram of Electrochemical machine.

TOOL

(a) (b)

Figure 7. Electrochemical machining.

erated at the cathode. The electrolyte may be

recirculated after filtering to remove this

sludge. ECM has come into extensive use in the

metalworking industries for the formation of

complex shapes of difficult metals such as one

piece, high alloy turbine blades, small holes and

fragile parts. The process parameters and

metalworking applications have been the subject of extensive recent reviews [1, 14, 15, 16,

17, 18, 19].

Essentially ECM involves passing current

through an electrolyte in the gap between the

workpiece and a suitably shaped tool (fig. 7).

In electropolishing the asperities on the surface

have the effect of concentrating the current to

the workpiece at these points so they are preferentiallv removed—in ECM the tool concentrates the current on these parts of the

workpiece from which preferential removal of

material is required, i.e., the shape of the workpiece then becomes complementary to that of

the tool as the process proceeds. The maximum

amount of material removed depends on the

current flowing and theoretically is given by

Faradays law. In practice, rates are usually

75-100 percent of theoretical.

The gap between the workpiece and tool depends on the material removal rate and how

the tool is advanced. Control of the working

gap is important also in its effect on the flow

of the electrolyte. Adequate flow is necessary

to prevent overheating and boiling as a result

of joule heating since overheating changes the

conductivity of the electrolyte. Depending on

the application gaps vary from 0.005 — 0.030

in (0.0127-0.076 cm). Small gaps give more

accurate shapes but large pressures are re- quired to give adequate electrolyte flow.

Although the main function of the electrolyte

is to provide ions to carry current in the cell,

the electrolyte must be capable of dissolving

the workpiece and avoid the formation of insoluble products that may passivate the surface.

The ECM process has many points of similarity to EDM but there are also some differ- ences of importance to the machining of ce- ramics and other brittle materials. Like EDM,

the use of ECM requires that the work be

electrically conductive, and it is not affected by

the hardness of the work. The tool does not

normally contact the work, and the process is

therefore well suited to the fabrication of brittle materials.

Specific types of material removal processes

applicable to ECM include turning, trepanning,

drilling and milling (fig. 8). By feeding a shaped tool into a rotating workpiece, sequential turned parts can be made. Only small tools are required for even large pieces and thin

sections can be made since there are no distortions. Milling of tungsten carbide tools can be

done without an abrasive using an aluminum

203
TURNING

DEEP HOLE DRILLING

TREPANNING I

MILLING

Figure 8. Applications of Electrochemical machining.

wheel and sodium nitrite. In drilling the tool is a tube with insulated sides and the electrolyte

is forced down the center. Although overcutting is minimized by insulating the tool down to

the tip, some enlargement is always found.

ECM drilling is particularly useful for producing deep small holes, i.e., 0.035 inch diameter x

12-24 in long (0.089 cm by 30-60 cm).

Unlike EDM, however, no melting of the

work takes place, and no stress or damage is

introduced into the machined surface. In many

cases a very smooth or polished surface is obtained. A wide choice of electrolytes is available, and the electrolyte must be chosen to produce the desired surface on the material being

worked. Some factors influencing the choice of

electrolytes are discussed by Hoare, et al., [20],

and by Boden and Evans [21]

3.2. Electrochemical Grinding

Although ECM has not been extensively used

for cemented carbides, a related process—electrochemical grinding or ECG—has been extensively used for shaping and sharpening carbide

tools [1, 15, 25, 26]. In this process the tool is a metal-bonded diamond wheel which rotates

against the work and is fed with an electrolyte

instead of a conventional coolant (fig. 9). CurABRASIVE

PARTICLES

ELECTROLnE

Of0 *

c

c

WORKPIECE +

METAL-BONDED

GRINDING WHEEL

INSULATED

SP-INDLE

Figure 9. Electrolytic grinding.

rent is fed to the wheel through a slip-ring and

passes through the electrolyte and the work, the

latter being the anode. Metallic contact is prevented by the diamonds projecting from the

wheel surface. Most of the stock is removed

electrolytically with only a minor fraction being removed by mechanical grinding. As a result, the consumption of diamond is reduced,

the rate of cutting is increased and the surface

finish is improved. The ECG process is said to

reduce grinding costs by 25-40 percent and to

reduce diamond consumption by 90 percent

[15].

The main function of the abrasive grain is to

provide insulation between the anode and cathode and to determine the gap. The particles also

continually remove any passive layer that

might form. Wheels for electrolytic grinding

differ from conventional wheels in the concentration and size of the abrasive particles and

in the method of manufacture. Since relatively

low metal removal rates are required and because of the mechanical action of the wheel,

the electrolytes are generally chosen to avoid

corrosion of the machine. Solutions of sodium

nitrate, sodium nitrite, sodium phosphate, etc., are common for cemented carbides. Phosphate

dissolves the cobalt binder phase and should be

avoided. Sodium silicate tends to give better

results for WC/Co materials. Current densities are of the order of 100-800 amp/sc[ in (16-150

amp/cm-) at 5-20v,

3.3. Application to Ceramics

Up to the present time ECM processes appear to have seen little application to refractory materials.

Cook and his coworkers [22] attempted to

increase the material removal rates by using

high pressure aqueous electrolytes and molten

salt electrolytes. A number of electrolytes were

investigated (fig. 10) and the effect of fused

NaOH was significant. The material removal

rates for several electrolytes are given in table

5. The use of high pressure and fused salts to

get higher current densities without boiling of

the electrolyte results in substantially increased

removal rates.

Lee and Haggerty reported on electrochemical machining of ZrC [23] and ZrB, [24] . In

these cases the process was used to obtain thin

sections for electron transmission microscopy.

Single crystals were oriented by X-ray back

reflection and the goniometer was transferred

to an EDM setup and sections 0.010 in

(0.0254cm) thick were cut. The sections were

then electrochemically thinned. The apparatus

was a typical laboratory electropolishing setup

(fig. 11). The electrolyte was four parts methyl

alcohol to one part perchloric acid at —35 °C.

The situation was more akin to electropolishing

in that thin flat specimens were required and

204
Table 5. Cemented tungsten carbides [22]

Condition

NaoH (20%)

Electrolytic grinding

Pressurized tool and

electrolyte

Fused salt (NaoH) - 900° C

Current densities

(Amps/m'XlO-')

465

930

1450

2930

1800

Material

removal rate

(m/sec XlO^)

2.54

6.35

33.5

41.5

28.5

is 1

1 1 r

TUNGSTEN CARBIDE

NaOH (Fused)/

NaN02

NaOH (20% Solution)

^used^)

ZnCe2(Fused)

60 120 180 240

t (sec)

300

Figure 10. Material removal

rates for tungsten carbide in

various eletrolytes [22].

the working gap was very large. Sections were

of the order of 500-1000 A thick.

The problem of possible damage to the section

was considered. Etching and X-ray topography were used to show that the dislocations

in as grown crystals were largely in low angle

boundaries and that the area between the cell

boundaries were essentially free of dislocations

(fig. 12). The misorientation calculated for the

X-ray image was of the order of 0.01 degree

and therefore a dislocation spacing of ^1.6

microns. Examination of the ECM prepared

sections in electron transmission revealed few

dislocations and those that were observed were

in row with spacing of approximately 1 micron

(fig. 13). Another example is the Burger vector

determinations made on ZrBa (fig. 14).

We have found the thinning technique to be

applicable to the carbides and diborides of

hafnium, tantalum, zirconium and titanium as

well as MoSio.

COOLING

BATH

SPECIMEN

(Anode)

ELECTROLYTE

Figure 11. Cell for electrolytic

polishing [15].

4. Chemical Machining

4.1. Description of the Process

Chemical machining (CHM) is a term re- cently applied in the metalworking industry to

the removal of material by simple etching processes, without the use of an electric current

(fig. 15). The process shares most of the characteristics of ECM with the additional caution

that some etches may cause hydrogen embrittlement on sensitive materials. Much of the art

involved in CHM lies in the application of the

masking material or resist which blocks the

action of the etch from selected areas. Three

general methods are in common use. In the "cut

and peel" method a thick layer of masking

lacquer is applied to the entire metal surface

and allowed to dry. The areas to be etched are then cut out and peeled olT, thus exposing the

metal. This method is adapted to the production

of large parts with simple patterns such as

aircraft wing sections. Because the masking

agent can be heavily applied, it is capable of

withstanding deep and severe etches. Stepped

work can be made by removing portions of the

resist between etching periods.

The second method of masking involves the use of a photoresist, which is coated on the

surface and allowed to dry. The areas to be

protected are exposed to actinic light through

205
Figure 12. (a) Single crystal of ZrC etched on a (100) plane in 1:1

boiling HSO'' to show both individual dislocations and

dislocation of sub-boundaries, 9X.

(b) X-ray image of same area as in (a) showing the boundaries to be of a tilt character, 5X.

i #11. mmmmmtm

Figure 13. A row of dislocations which form a low

angle boundary in an undeformed ZrC single

crystal.

a suitable mask, which renders the resin insoluble in a developing solvent. The image is then

developed by washing off the unexposed areas with the solvent. The remaining image may be

baked to increase its chemical resistance and

the part is then exposed to the etching liquid.

The photoresists are generally not able to withstand severe etching conditions and they are mainly used for fine work and complex parts.

The third method of applying the masking

agent is by way of the silk screen method. This

method is adapted to moderately severe etches

and moderately complex patterns, but the ac- curacy is not as great as may be obtained by the

photographic method.

The etches applicable to various metals and

a general review of the process are given by

Springborn [1].

For many years simple etching processes

using wax resists and hydrofluoric acid etches

have been used for the decoration of glassware

and polishing of TV face plates. Similar methods have been used for frosting electric light

bulbs, although etched surfaces are now being

replaced by sprayed coatings. The production

of complex and precise parts of vitreous materials, however, became practical only after

the development of the photosensitive glasses.

4.2. Photosensitive Glass

Photosensitive glasses and processes for

chemical machining were described by Stookey

[27, 28, 29] of the Corning Glass Works in

1953 and following years. These glasses contain a metallic photosensitizing agent (such as Cu,

Ag or Au) which will precipitate as fine particles when the glass is exposed to light and

moderately heated. When the exposed glass is

further heated to a temperature below the softening point, these metalic particles serve as

nuclei for the growth of a crystalline silicate

phase, which eventually results in substantially

complete devitrification of the exposed areas. The areas which have not been exposed to

light remain in the vitreous form. If the heattreated glass article is now exposed to hydrofluoric acid, it is found that the exposed and

206
Figure 14. Dark field transmission electron photomicrographs of ZrBi.

(a) is taken with the (0110) plane diffracting. Magnification app.

17,500X.

ETCHANT

\ 1 1 1 1

PROTECTIVE

METAL

Figure 15. Chemical milling.

devitrified areas are much more rapidly attacked than the unexposed areas. In this way,

complex patterns, perforated parts and shaped

electrical insulators can be mass produced. After the parts are etched they can be converted

into a glass-ceramic (i.e., devitrified) by further exposure to light follovi^ed by heat treatment.

4.3. Application to Ceramics

The application of CHM techniques to ce- ramic materials has been restricted by the lack

of suitable etching agents. Politis and Ohtani

[30] give a list of suitable etches for a number

of ceramic materials to be used in preparing

samples for microscopic examination. Alford

et al., [32] give the rate of attack of 85 percent

phosphoric acid at 425 °C on the 0001 face of

sapphire crystals as 0.00025 in (0.00625 mm)

per min. Molten KsAlFg at 1000 °C attacks the

1120 surfaces at the rate of 0.0015 in (0.0375

mm) per min.

Tighe reports the use of hot orthophosphoric

acid to chemically thin single crystal sapphire

at rates of '—I micron per min. The material

removal rate is sensitive to crystal orientation

and temperature.

These low rates indicate that the methods are

suitable mainly for surface preparation, rather

than true shaping operations. A more serious

problem is that of finding suitable resists for

these conditions.

In reviewing the available literature on the

chemical machining of aluminum oxide, it is

quite apparent that the material removal rate

and the effects of the machining on mechanical

properties are dependent on several factors,

such as, the type of etchant used, the processing conditions of the machining operation, and

finally the compositional variations on the types

of alumina investigated. The variables in chemical machining of aluminum oxide has not been

systematically studied. However, there is

enough information in the literature to indicate

that there are major variations in the material

removal process that can have a profound effect on the physical properties of the machined

body. The work of Kirchner and his coworkers

on compositions containing 94-96 percent

aluminum is perhaps a unique application of

CHM. Starting with a two-phase structure in

which both the alumina and intergranular

phases are co-continuous, the action of the

etchant, in this case, hydrofluoric acid, was a highly preferential attack on the intergranular

phases leaving a matrix of high alumina with

207
Table 6. Chemical polishing of high alumina materials

A1203 Etchant Temperature Rate

Initial effect on strengtli

m/sec XIO' in /min X 10"'

94-96% AI2O3 (p = 3.86 gm/cc) H3P04 320° C Decrease

0.25 Si02, 0.25 MgO (p = 3.87) H3P04 320° C 1.6 4.1 Decrease

280° C 0.3 0.8 Decrease

(0.5 MgO) p = 3.98 H,Po4 320° C >0.01 N C

^V/.O iVigW/ p — O.i/O 280° C ^0 01 94-96% (p = 3.86) H=So4 250° C 0.1 0.2 Increase

(0.5 MgU) p = 6.yo "Mo n r\ sySi.2tii\Ji u.o n 0 U.J Increase

850° C 2.0 5.1 Increase

900° C 3.0 7.6 Increase

94-96% AhO, (p = ZM) Na2B407 850° C 2.0 5.1 Increase

94-96% AhO, (p = 3.86) KHS04 650° C Decrease

high pore volume. In this condition the high

alumina ceramic can be drilled, turned, and

otherwise machined. After shaping is complete

the pieces can be refired to restore most of the

original strength.

In a more general context there have been a number of observations on the chemical polishing or machining of high AloOa materials (table

6). This data illustrates the variables that can

affect the process, particularly the composition

of the materials and etchant used.

In general, phosphoric acid is detrimental to

the strength of polycrystalline ALO,, v^hile for

single crystals Kirchner reported strengths after polishing in 85 percent orthophosphoric

acid to be slightly greater than those of the as

received controls [34].

On the other hand, CHM of AI2O3 in fused

sodium borate has resulted in some substantial

increase in the strength of several polycrystalline materials [31, 32] as well as single crystals [34, 35].

The effects that CHM may produce on the

machined properties of AI2O3 material will depend on the dissolution process, i.e., the particular effect the etchant has in producing flaws or change the geometry of existing flaws. Preferential attack of grain boundaries or of impurities or producing etch pits may create new

flaws [35].

In figure 16 the changes in flaw geometry

which might take place are shown [38]. The

flaw may be rounded or blunted (c) in which

case we might expect the strength to improve

as a result of the material removal process. The

existing flaw may be extended or sharpened by

preferential attack at the tip as in (a) and the

strength would be expected to be lowered. In

case (b) the rounding of the crack may be offset by the increase in length and no effect on

the strength would be expected due to the cor- rosion or dissolution process. Thus, one may

find a rationale for strength increases, de- creases and no change as a result of CHM.

In the Charles-Hillig theory, flaw sharpening

is due to preferential attack at the crack tip

(a) (b) (c)

Figure 16. Hypothetical

changes in flaw geometry due to corrosion or

dissolution.

(a) Flaw sharpening.

(b) Flaw growth such

that the rounding

of the tip by stress corrosion balances

the lengthening of

the flaw.

(c) Simple rounding by

corrosion or dissolution ( Joffe effect)

[38]

through a stress corrosion mechanism. In CHM

there is no imposed stress, however, we have

found that internal stresses of a thermal nature are sufficient in some cases to provide that

necessary for enhanced corrosion.

For instance, polycrystalline MoSia undergoes catastrophic failure in Oo or air at temperatures of 400-500 °C—well below the temperature of significant oxidation—with no external

stress applied. Single crystals are not effected.

Under similar conditions, however, with ex- ternal stress, single crystals exhibited static

fatigue in the same manner as polycrystalline

materials (fig. 17). These results are consistent

with the concept of crack sharpening due to

stress corrosion according to the Charles-Hillig

theory.

We could postulate that depending on the

dissolution action of the etchant in producing

new surface flaws or altering existing ones, the

strength of chemically milled AI2O3 may increase, decrease or remain constant (fig. 18).

As far as subsurface damage is concerned,

such as introduction of dislocations, the evi208
LOG V (sec)

LOG r (hours)

Figure 17. Delayed failure experiments for single crystal,

polycrystal and MoSii.

Figure 18. Effect of chemical machining on the transfer rupture strength

of polycrystalline alumina oxide.

209
dence is that no effects have been observed by

Tighe [39] in AUO-, or by Lee and Haggerty in

ZrC or ZrB^.

5. Summary

The process of EDM, ECM, and CHM are attractive techniques for material removal for

ceramics ; however, little use has been made of

these processes. Some part of this is due to the

fact that two of the processes are limited to

material of relatively good conductivity and to

the low material removal ratio of CHM.

It appears that these techniques will continue

to be used as a method of preparing special

shapes of materials in the laboratory.

In those areas where it has been applied, lit- tle information is available on the nature of the

prepared surfaces and the extent of any damage. Perhaps a less impirical base would be

helpful to choosing machine condition to obtain a desired surface condition and limited but well

defined damage.

6. References

[1] Springborn, R. K., Ed., Nontraditional machining

process, (American Society of Tool and Mfg.

Engineers, Dearborn, Michigan, 1967).

[2] Whiteman, A. R. C, and Leadbetter, R., Electrical discharge machining covering the period

1960 to 1967 (Atomic Energy Research Establishment, Harwell, Berkshire, 1968), (Also

available from National Technical Information

Service, Springfield, Va. 22151; Accession No.

N68-28165).

[3] Krabacher, E. T., Mod. mach. shop 41, (8), 130-

137 (1969).

[4] Bucklow, I. A., and Cole, M., Metallurgical Reviews 14, (135), 103-18 (June 1969).

[4a] Walson, R. P., Haworth, W. L., and Birnbaum,

H. K., Spark Erosion cutting of Germanium,

ONR Technical Report (July 1, 1970).

[5] Samsonov, G. V., and Mukha, I. M., Poroshkovaya

Metallurgiya, 7 (67), 81-87 (July 1968).

[6] Ibid, 8 (68) 84-90 (Aug. 1968).

[7] Lee, D. W. et al.. Investigation of the Mechanical

Properties of Transition Metal Borides and

Carbides, Final Report to U. S. Atomic Energy

Commission under contract No. AT(80-1)-3411

(14 October 1967).

[8] Lanin, A. G., Fedotov, M. A., and Glagolev, V. V.,

Poroshkovaya Metallurgiya 5 (65), 97-101,

(May 1968).

[9] Yanchur, V. P., Andrievtkii, R. A., Spivak, I. I., and Fedotov, M. A., Neorganicheskie Materials

5 (No. 6), 1012-17 (June 1969).

[10] Artamonov, A., Ya., Lemeshko, A. M., and Murzin, L. M., Poroshkovaya, Metallurgiya 9, 95-

100, (Sept. 1969).

[11] Matsumaya, Y., and Fukatsu, T., in Powder

Metallurgy, Leszynski, W., ed., 683-700, (Interscience Publishers Inc., N. Y., 1961).

[12] Filimonenko, V. N., Gurland, J., and Richman,

M. H., Metallography 2, 125-36 (1969).

[13] Saito, S., Katagiri, M., and Hasegawa, S., J. Ceram. Assoc. Japan 70, (6), 195-9 (1962).

[14] Olofson, C. T., Gurklis, J. A., and Boulger, F. W.,

in Machining and Grinding of UltrahighStrength Steels and Stainless Steel Alloys, Report on Contract No. DA-01-021-AMC11651 (Z), NASA-SP-5084.

[15] DeBarr, A. E., and Oliver, D. A., editors. Electrochemical Machining (Elsevier, N. Y., 1968).

[16] ThrooT), J. W., A Mathematical analysis of Electrochemical Machining Phenomena. Paner No.

WES7-09 presented to ASM-ASTME 1967

Western Metal and Tool Conference, Los

Angeles, (March 13-17, 1967).

[17] Bellows, G., ECM Machinability Data and Ratings, Paper No. WES7-10, presented to ASMASTME 1967 Western Metal and Tool Conference, Los Angeles, (March 13-17, 1967).

[18] HoDenfeld, J., and Cole, R. R., Electrochemical

Machining-Prediction and Correlation of Process Variables, Paper No. 66 -Prod 5 presented

to ASME Production Engineering Conference,

Detroit. (April 18-22, 1966).

[19] Weingartner, C. J., Electrochemical Machining

During Eight Years of Operation, Paper

680659 presented to SAE Aeronautic and

Space Engineering and Manufacturing Meeting, Los Angeles, (October 7-11, 1968).

[20] Hoare. J. P., LaBoda, M. A., McMillan, J. L., and

Wallace, A. J., J. Electrochem. Soc. 116, 199-

203 (1969).

[21] Boden, P. J., and Evans, J. M., Nature 222, 377-

8 (1969).

[22] Cook, N. H. Loutrel, S. P., and Meslink, M. C,

Increasing Electrochemical Machining Rates,

MIT Final report on contract DA-19-066-AMC268 (WK AD659, 004.

[23] Lee, D. W., and Haggerty, J. S., Jour. Am. Ceram.

Soc. 52, 642-7 (1969).

[24] Leombruno, W. J., Haggerty, J. S., and O'Brien,

J. L., Matl. Res. Bull. 3, 361 (1968).

[251 Benz, H. G.. Abrasive Eng. 15, 41-42 (1969).

[26] Marten, K. H., and Heck, K., Werkstatt u Betrieb

101, 615-20 (1968), in German.

[271 Stookey, S. D., Ind. Eng. Chem. 45, 115-8 (1953).

[28] Stookey, S. D., U. S. Patent 2,684,911, July 27,

1954.

[29] Corning Glass Works, British Patent 752,243, July

11, 1956.

[30] Politis, C, and Ohtani, S., Praktische Metallogranhie 4, (8), 401-16 (1967).

[31] Gruszka, R. F., Mistier, R. E., and Runk, R. B., Ceramic Bull. 49 (6), 575-9 (1970).

[32] King, A. G., Materials Science Research, Vol. 3,

edited by W. W. Kreegel and H. Palmaur III, Plenum Press, New York, (1963).

[33] Bortz, S. A. Ihid.

[34] Alford, W. J., Bauer, W. H., and Motolka, R. W.,

Perfection of Single Crystal Alumina Grown

in R. F. Plasmas and by Other Techniques,

School of Cermaics, Rutgers, the State University, New Brunswick, N. J., June 14, 1966,

AFCRL-66-524.

[35] Richel, P. A., Infield, J. M., and Kirchner, H. P., Ceramic Bull. 47, 702-6, (1968).

[36] Kirchner, H. P., Gruner, R. M., and Walter,

R. E., J. App. Phys. 40, 3445, (1969).

[37] Malliner, F. D., and Proctor, B. A., Proc. Brit.

Cer. Soc. 6, 9, 1966.

[38] HuUig, W. R., and Charles, R. F., in High

Strength Materials Surfaces, Stress-Dependent Surface Reaction and Strength, V. F. Zackay, ed. John Wiley and Sons, New York

1965).

[39] Streilow, W. H., J. of Appl. Phys. 40 (7), 2928-

2932, (1969).

[40] Gatos, H. C, and Levine, M. C, in Progress in

Semiconductors, Vol. 9, Edited by A. F. Gabron

and R. E. Burgess, Temple Press, London,

(1965).

210
Discussion

Rice: Would you expect a higher removal

rate and possibly less complex surfaces in the

machining of materials such as silicon carbides

which normally vaporize w^ithout melting?

Lee: Yes. I v^ould expect that you v^^ould still

get some thermal shock in the material. Thermal shock seems to be the major factor causing

the poor surface and/or the cracking. I did not

find any references to silicon carbide but it would seem to me that silicon carbide is a commercially available refractory material that

could be cut and machined quite easily by EDM

type techniques.

Lasson: Pertaining to the photosensitive

glass that you mentioned, do you know what

dimensional tolerance range you would be able

to hold? Also, do you get any tapering when

you do the etching?

Lee: The tolerance is ± 10 percent. Yes, you

do get a 2° taper.

Martis: a word of caution concerning EDM

and its effect on the removal of cobalt on a

wide variety of tungsten carbides. They propagate brittle fracture. This is a major cause of

failure.

Lee: Yes, the work done at Brown I believe

showed there were extensive areas of cracking

of the material, although they didn't make any

physical measurements afterwards.

Martis: We found you have to remove that

particular layer using conventional diamond

abrasives.

Lee: That would correlate with the Russian

work where they did have to remove at least 15

microns of the material before the strength of

the material was restored.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Sc.ence of Ceramic Machining and Surface Finishing Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).

211

III. TECHNIQUES AND MECHANISMS OF

SURFACE FINISHING

Session Chairman

P. J. Gielisse, University of Rhode Island

Improvements in the Surface Finish of Ceramics by Flame Polishing and

Annealing Techniques

M. J. Noone

General Electric Company, Philadelphia, Pennsylvania 19101

A. H. Heuer

Case Western Reserve University, Cleveland, Ohio 44106

Improvements in the surface perfection (i.e. "smoothness") of ceramics can be

achieved by thermal treatments at temperatures where material transport may be

induced at the surface. Such treatments include annealing in controlled environments,

where surface diffusion, vapor transport mechanisms, etc., may be active, and flame

polishing, where a thin layer at the surface of the material is melted, allowed to flow freely, and then resolidified.

The effects of these treatments on the surface of single crystals of a-AhO, are reviewed in this paper. The most sensitive assessment of the degree to which surface

perfection is attained is the measurement of the strength of the treated crystals;

for this reason, strength data are used extensively to characterize the thermal treat- ments described. In the case of flame polishing, it found that a sufficient degree of

surface perfection can be attained so that the strength is no longer limited by

surface defects (the usual experience with ceramics) but by defects within the material. The technological significance of high strength AI2O3 filaments for reinforcement of metals and the unfortunate deterioration of the surface in these applications

is briefly described.

Key words: a-ALOs; annealing; flame-polishing; reinforcements; ruby; sapphire;

strength; surface perfection.

1 . Introduction

Improvements in the surface finish of singlecrystal specimens of a-AloOa can be achieved

by thermal treatments at temperatures up to

and including the melting point. Two regimes

of temperature are considered separately in

this paper: (a) annealing at temperatures up to

1800 °C; and, (b) flame polishing in a steep

temperature gradient which includes melting

of a thin layer of material at the surface. Improvements in the surface perfection or smoothness arise by movement and redistribution of

material at the surface; in the former case (a)

by surface diffusion and/or by evaporation and

subsequent recondensation, and in the latter case (b) by rapid material transport in a liquid

film. The improvement in the surface perfection which arises from any particular treatment can best be measured by determining the

changes in strength that are produced in the

material, since the strength of single-crystal

brittle materials is determined almost entirely

Since the strength of polycrystalline ceramics invariably de- creases with increasing grain size, thermal treatments which in- duce appreciable grain growth will also lead to a loss in strength

by the condition of their surfaces. Flame polishing in particular produces striking changes

in strength, since it completely melts all preexisting surface damage and leads to a "liquid

smooth" surface.

Although this paper is devoted exclusively

to single crystal « — AI2O3 (sapphire) and/or

ALO.TrCr (ruby), the techniques and mechanisms discussed are applicable to a wide range

of ceramic crystals. In the case of brittle ce- ramic polycrystals however, thermal treatments

generally do not produce such profound changes

in strength as are found for single crystals;

microstructural features absent in single crystals—grain boundaries, second phases, and

voids—may limit the surface perfection achievable. Furthermore, such "defects" may be preferentially etched or otherwise accentuated by

thermal treatments,^ thus leading to a degradation of strength. The low strength of brittle

single crystals is a result of an abundance of

surface defects microcracks and notches

which are introduced during the cutting and

grinding operations used to shape a specimen

from a bulk crystal. Since such materials are unable to undergo appreciable plastic deforma439-921 0 - 72 - 15 213
tion at room temperature, stress concentrations

arise at the tips of these cracks (or notches)

when the crystals are stressed and result in low

strengths. The flaw theory of brittle fracture

advanced by Griffith completely accounts

for this Dhenomenon. (Griffith was able to show

that high strengths could be attained in brittle

materials (glass fibers) by careful preparation

of flaw-free surfaces.)

The existance of a surface perfection/

strength relationship in bulk « — ALOh crystals

was demonstrated most spectacularly by Morley and Proctor [2] , who measured strengths in excess of 7 GN/mMlO'^ psi) in carefully flamepolished rods. The strength of these rods was

similar to that of a-ALO., whisker crystals fS]

which were considered to have highly perfect

internal as well as surface structures. However, the surfaces of a-ALO., whiskers are not

atomically smooth ; on the contrary, Mehan and

Feingold [4] reported growth steps and other

surface nrotruberances to be common features.

Marsh [5] showed that a step on a crystal sur- face could cause a stress concentration of the

same magnitude as that around a crack of similar dimensions, which accounts for the fact

that, even with whiskers, the measured

strength, although often high, rarely approaches the theoretical estimates of the

strength. ( For alumina, the theoretical strength

E

is of the order of —-p— [6] )

15

Although annealing can also increase the

strength of ground sapphire, the strength in- creases achieved are not as dramatic as for

flame polishing ; presumably, the degree of sur- face perfection achieved is greater for the latter process. Nevertheless, annealing may have

some technological importance, inasmuch as flame polishing appears to be restricted to

simple shapes.

In view of the fact that there appears to be

unfamiliarity with the techniques of flame polishing, the nature and experimental details of

the process are discussed at some length in this

paper (sections 2 and 3). By way of contrast,

the literature on annealing is reviewed only

briefly, with major emphasis on the degree of

strength improvements that have been re- ported (section 4).

2. Previous Work on Flame Polishing

The earliest reported exploitation of flamepolishing on sapphire was in 1952 by Popov [7]

and by Barnes and McCandless [8] who independently developed the process to improve the

surface smoothness of crystals for use as thread guides in the processing of textile fibers.

Flame polishing was found to be superior to

chemical and mechanical polishing for this pur-

' Figures in brackets indicate the literature references at the end

of this paper.

pose and has been in general use for many

years as a means to reduce friction in thread

guides and to improve the surface appearance

of simple cylindrical shapes.

Barnes and McCandless [81 described a process for flame polishing sapphire and ruby rods

which produced scratch-free, glossy, and wearresistant surfaces suitable for low friction applications. They melted the surface of a rotating rod using a flame which traversed the

length of the rod, and left a resolidified

smoothed surface in its wake. Only rods less

than 5mm diameter could be successfully polished without fracturing and the best quality

surface was achieved on rods in which the c- axis made an angle <65° with the rod axis

(growth direction). Flame-polished a-ALO,,

rods had an average tensile strength of 0.7

GN/m^ compared to 0.46 GN/m- for the as- grown material.

Popov was also the first to report plastic

bending of sapphire rods and developed an

apparatus for bending flame-polished rods to

form curved textile guides [7]. Although the

conditions were not well defined in his work

("temperature near to the melting point") a

relationship was noted between resistance to

bending and the crystallographic orientation in

the plane of bending, with almost a factor of

two difference in the bending force required between diff'erent orientations. The rods had

"slip-plane traces" on their surfaces after

bending [7] and no doubt also contained extensive deformation twins. The strength of the

flame polished rods was not determined during

Popov's work.

Popov used as-grown Verneuil sapphire rods,

4—4.5mm in diameter, and polished these by

rotating them in an oxy-hydrogen flame at

110-115 rpm in a vertical plane: the polishing

was performed within a muffle furnace, which

allowed preheating of the rods and slow cooling after polishing. The flame gases were

metered separately and mixed roughly in the

ratio 102:2H,, by volume; they were fed

through a burner orifice approximately 4 mm

in diameter at a distance of 75—80 mm from

the rod. The rotating rod was withdrawn downwards through the flame at a rate of 20—25

mm/min. by an automatic screw mechanism.

Popov recommended that the polishing operation be repeated at least twice on each rod and

that for uniform polishing, it was essential

that the rod be concentric with the furnace and

for the rotation rate, withdrawal rate, and gas

flows to be maintained constant during the operation. No optimum conditions could be es- tablished and each polishing unit had to be set

up individually, by qualitative assessment of

the polished surface in reflected light microscopy at low magnification.

The surface features observed on these rods

214
were similar to results obtained by the present

authors using "over-polishino:" conditions (section 3.1 below). Although Popov states that

ideally only a "thin" layer of material was melted, it is likely that he had to fuse up to 0.5

mm depth of material to accommodate irregularities in diameter and the slightly elliptical

nature of the cross section of the as-grown rods.

A pronounced ripple was evident on the surface

of the rods and many entrapped bubbles were

found, together with parabolic markings which

probably represent a tendency for the development of crystal facets (see section 3.1) ; the

ripples and thermal etching effects were found

to vary with, and relate to, the crystall ographic

orientation of the rod. Popov also made the observation that contamination of the surface by

refractory dust (from abrasion of furnace

bricks, etc.) during polishing would result in

fusion of particles into the surface, rendering

the rod unserviceable as a textile guide. As will

be shown later, the presence of contaminants

or even coarse ALO-, particles fused into the

surface severely limits the strength that can

be achieved in flame-polished rods.

Sapphire and ruby rods are available in

flame-polished form from a number of commercial suppliers ; however, the extent of the polish

is generally limited to the point where the material surface is improved aesthetically by appearing glass-smooth on a macro-scale (underpolished in the context of the rods described in

this paper). These materials have been used

for studies of the mechanical behavior of

a-ALOs (e.g., ref. 9) and the fracture stress

is generally only modestly higher than centerless ground rods. A significant advance in the

state of the art of flame polishing was made by

Morley and Proctor [2] who showed in 1962

that the technique could be used to produce

surfaces on sapphire rods with such a high degree of perfection that the strength of the re- sulting crystals was not limited by the surface

condition but by internal defects. Carefully

polished rods were bent to over 2 percent elastic strain, and for the first time, stresses in excess of 7 GN/m^ were measured in bulk ceramic

crystals [2] Proctor, in discussing these results, [10] compared them with results from

fused silica and silicate-glass rods with highly

perfect surfaces prepared with similar care, and

concluded that a strength /size relationship was

not valid in sapphire or glasses ; a strength/surface-condition relationship dominated the mechanical behavior. Further work by Mallinder

and Proctor [H] showed that internal defects,

inclusions, bubbles, etc., became significant in

limiting the strength of rods when the surface

was highly perfect and that the highest

strengths were attained in rods with a high degree of both surface and internal perfection.

Mallinder and Proctor described the flamepolishing of centerless-ground sapphire rods

3/4, 1, or 2-1/4 mm diameter and 6-8 cm long

[11, 12]. These rods were machined from

boules grown by the Verneuil technique and

were obtained from several sources (Linde

Co. USA, Salford Electrical Co.. England, and

Hrand Djevahirdjian SA, Switzerland): the

crystallographic orientation was random with

the c-axis generally making an angle of about

60° with the rod axis. The rods were cleaned

for 10 min in an ultrasonic bath containing an

aqueous solution of a wetting agent, rinsed

in distilled water, and allowed to drain

dry. Polishing was accomplished by rotating

the rods at about 200 rpm while an oxygen/

coal-gas flame traversed along with their length

at a rate of about 20 mm/min; the flame size

and temperature were varied by adjusting the

gas pressure. The rod rotation speed, traversing rate, and jet-to-rod distance were chosen

arbitrarily and then held essentially constant

during the polishing studies, so that the degree

of polishing and quality of the resultant sur- face was controlled by carefully regulating the

flow of gases to the burner jet.

The quality of the polished surface was as- sessed using incident-light microscopy at high

magnification. Optimum flame conditions produced a clear glassy transparent surface, although a faint spiral marking corresponding

to the pitch of the flame traverse (pitch = 0.1

mm under the conditions used)' was observed.

Using a cooler flame, all traces of grinding

marks in the surface were not removed; with

a more intense flame, a cloudy surface appearance resulted from the generation of voids

within the resolidified layer. Over-polished

crystals were found to have low strength, similar to that of poorly-polished material, so that

precise control of the flame conditions was nec- essary to ensure high strengths.

By examining a necked area which formed

when the flame was extinguished during polishing with an intense flame, Mallinder and

Proctor [11] deduced that a "neck" of molten

material existed throughout the polishing process and traveled along the rod with the flame.

In well-polished rods, the molten neck resolidified epitaxially from the inner core outwards,

the surface thus reflecting the structure of

the underlying crystal. With an excessive depth

of liquid, solidification also began at the outer

surface and proceeded inwards; the change in

volume as the entrapped liquid subsequently

solidified led to the formation of the subsurface voids. Careful examination of the crystallographic orientation of polished rods by optical microscopy and x-ray techniques showed

that the resolidified surface crystallized with

identical structure to the underlying crystal

and that subgrain boundaries at the crystal

surface persisted after polishing.

215
Well-polished rods were carefully selected

for freedom from internal defects—bubbles

and inclusions, which were rendered visible

during the polishing operation by the intense

light generated at the molten surface—and

were tested in bending with a gauge length of

3 mm or 6 mm selected from the most defectfree portion of the rod. Extreme care was taken

to prevent any mechanical contact with the

rods after polishing. Only the extreme ends

were handled and the bending load was applied

through epoxy sleeves cast onto the ends of

the rods and then fixed into an elegant testing

device, which applied a pure bending moment

to the rod without offering any other constraints [11]. This test arrangement was operated in a manner which allowed the rods to

be broken while immersed in a small bath of

glycerine to reduce the degree of shattering

and confine the fragments after fracture. By

patient examination of the fragments and careful matching of fracture faces, it was possible to positively identify and locate the sources

of fracture initiation. In well-polished rods, the

origin of fracture could invariably be traced

to internal (growth) defects. Curiously. Mallinder and Proctor [11] state that although

dirt on the rod surface gave cloudiness and

complex patterns when fused into the surface

during polishing, this did not appear to affect

the strength. (This is contrary to the experience of the present authors, who found that

prepolish cleaning reduced the scatter in

strength and increased the average strength

of the rods by eliminating many potential flaws

("dirt").)

No relationship between strength and crystallographic orientation was found in the above

work. However, an important practical observation was that the perfect surfaces achieved

by careful flame polishing were extremely sensitive to damage by mechanical abrasion. The

strength of the rods was reduced by a factor

of between 2 and 14 by merely allowing the

rods to roll under their ov^ weight over a

surface sprinkled with 90 mesh silicon carbide

grains. Proctor noted that the sapphire rods

were markedly less susceptible to abrasive

damage than similar high strength silicate glass or fused silica rods and it was suggested that

the dependence of strength on crystallographic

orientation found in sapphire by other investigators might have been a reflection of the dependence of susceptibility to damage on orientation [10].

An interesting recent development in highstrength alumina single crystals is the con- tinuous sapphire rods and filaments produced

by pulling crystals of uniform cross-section

directly from the melt [13, 14] these have

surfaces sufl^ciently perfect to permit tensile

stresses of about 3 GN/m^ to be attained before

fracture. Nevertheless, flame polishing techniques were found to improve the strength of

these filaments (section 3.4.d) and tensile

strengths up to 3.5 GN/m^ have recently been

reported in an effort to polish the filaments for

commercial exDloitation of the process [15].

Flame-polishing techniques have thus been

used to produce alumina specimens whose sur- face perfection is such that the fracture stress

is no longer limited by surface features but

by internal defects within the crystals. It now

appears that continuous high strength filaments will soon be available with strength close

to the high levels previously found only in

laboratory specimens and with good uniformity (small scatter). The greatest potential for

exploitation of these filaments is as a reinforcement in a metal matrix at high temperatures. Unfortunately, most useful matrices are chemically incompatible with the filaments and

the surface perfection may be degraded during

composite fabrication and subsequent use. Preservation of the highly perfect flame-polished

surface has proven to be a difficult technical

feat in such matrices [16]. and useful exploitation of these high strength filaments depends

on the extent to which these highly perfect

surfaces can be retained.

3. Experimental

3.1. Flame Polishing

The aim of the work described in this section

was to explore the possibility of producing

high strength a-AljOg rods on a routine reproducible basis by a simple flame-polishing process. The work of Mallinder and Proctor [12]

was used as a guide in establishing the initial

polishing technique.

Single crystal a-AlaO,, rods were obtained for use in these polishing studies in various forms.

Centerless-ground rods, 0.5 mm and 0.25 mm

diameter sapphire and 0.5 mm diameter ruby

(0.5 w/o CroOs nominal doping level) machined from Verneuil-grown boules, were obtained from ("Linde") Union Carbide Corp.,

Crystal Products Division, Union, N.J. ; 0.5

mm diameter commercially flame polished sapphire and ruby rods were obtained from the

same source. In addition, 0.5 mm diameter

centerless-ground sapphire and 0.5 mm diameter commercially flame polished sapphire rods

were obtained from Insaco Inc., Quakertown,

Pa. A small quantity of 0.5 mm diameter centerless ground sapphire rods machined from

boules grown bv the Czochralski technique was

kindly provided by the Union Carbide Corp.

for comparison with the Verneuil material.

Finally, continuous filaments of melt-drawn

sapphire, 0.5 mm and 0.25 mm diameter produced by Tyco Inc., Waltham, Mass., were included in the studies.

216
Figure 1. Apparatus used for flame-polishing the sur- face of a AhOi rods.

A flame-polishing rig, shown in figure 1, was

constructed which consisted of a small "hobby"

lathe ("Unimat-SL", Model DB-200, American

Edelstaal Inc., N.Y.) arranged so that a sapphire rod, when held in a pin-vice in the lathe

chuck, was rotated with its long axis vertical.

The lathe drive was modified to provide con- tinuously variable rotation speeds from 0 to

~600 rpm ; the synchronously driven tool hold- er moved up and down in a vertical plane at

traversing speeds related to the chuck rotation.

The tool holder carried an oxygen/hydrogen

microtorch, which was used with a range of

fine burner jets cut from hypodermic needles, as shown in figure 2. This rig was operated under a wide range of

Figure 2. Detailed view of the burner jet and the impingement of the flame (barely visible) on the rod

during polishing. The apparent double-image of the

rod is a reflection from the metallic heat shield behind the rod.)

conditions so that the effects of each variable

could be assessed and a suitable set of operating

conditions chosen. The rods to be polished were

usually 10-12 cm long and 0.5 mm diameter,

but rods up to 25 cm long and 3 mm diameter

have been successfully polished in this same

rig.

a. Flame Temperature

The gas supply to the flame consisted of a

stoichiometric mixture of hydrogen and oxygen generated by electrolysis of water in a

small generator adjacent to the lathe ("Water

Welder (R)", Henes Manufacturing Company,

Phoenix, Arizona) ; the two gases were never

separated and were delivered directly to the

burner in one line. The combustion temperature

was about 3300 °C, but could be varied by the

deliberate introduction of vapors into the gas

stream. Initially, a hot flame, with the gases

dried by a CaCL tower, was used; however, it was found that a greater degree of control

could be obtained by using a cooler flame (about

2500 °C) produced by bubbling the gas stream

through methyl alcohol. The flame was also

rendered more visible by this procedure and

had a clearly defined "inner cone" which facilitated positioning the flame with respect to

the rod. No diff"erence in appearance of the

rods or in their strength was detected between

rods polished in the dry flame or in the slightly

reducing "alcohol" flame.

b. Flame Size

The size of the flame was determined by

the size of the hypodermic needle jet and by

the gas pressure ; the addition of alcohol vapor

increased the size of any given flame. The gas

generator had a sensitive pressure gauge which

was used to maintain a constant supply pres- sure of 5kN/m^ (controlled by a Variac applying voltage to the electrolytic cell). A range

of jet sizes from No. 16 to 27 (1.05—0.19 mm

bore) was used during the work. The larger

jets provided a flame which was too large to

be conveniently accommodated in the space

available, and provided too large an excess of

thermal energy compared with that required

to flame polish a slender rod: flames from the

smaller jets were only capable of heating a

small volume of the rod with a steep temperature gradient on either side of the hot spot. A

No. 23 jet (0.32 mm bore) was found to provide a flame ^:il2 mm long X 1 mm diameter,

inner cone 2.5 mm long, which was suitable

for polishing the 0.5 mm diameter rods. The

flame was large enough to envelop the rod, and

produced an apparently molten length of about

0.5 mm of the surface. A steep temperature

gradient existed above and below the molten

zone.

217
c. Temperature of the Rod Surface

The temperature at the surface of the rods

was determined by the distance of the burner

jet from the rod and by the rate of movement

of the flame over the rod. The distance from

the rod was the most critical parameter in the

production of well-polished rods and was controlled by the drive screw of the lathe toolpost

on which the burner jet was mounted. The

position of this screw could be read to ±0.05

mm and noticeable differences in surface appearance could be detected with changes of 0.2

mm in the rod-to-jet distance.

The rod surface temperature and /or the

volume of molten material changed appreciably

with changes in the rate of movement of the

rod through the flame. For this reason, the

rod-to-jet distance necessary to produce a good

polish was adjusted to suit the speed of ro- tation of the rod ; larger distances were needed

for slower speeds (or smaller diameter rods)

and vice versa. Attempts to measure the temperature at the

rod surface by optical pyrometry and to use

this as a means of control were unsuccessful.

It was soon realized that the simplest, practical means of assessing the optimum temperature and the quality of the polished surface

was by visual observation of the rod during

polishing through dark-glass "welders" goggles.

d. Lathe Spindle Rotation Speed and Flame

Traversing Speed

The speed of rotation of the lathe spindle

was initially varied from about 6 to 600 rpm

using Variac-controlled motors and a range

of pulleys and belts. The lathe chuck was

aligned such that a 10—12 cm long rod rotated

within 0.05 mm out-of-true about its long axis.

Longer rods had to be supported at both ends

to ensure stability of rotation and satisfactory

polishing. The rate of traverse of the flame was

directly linked via pulleys to the lathe spindle

and initially was 25 mm of travel per 1000

revolutions of the rod. While it was realized

that a high output rate would be desirable for

ultimate exploitation of the polishing process,

it was thought in the early experiments that

better quality surfaces would be produced at a low polishing rate. Initial work was done

on rods revolving at 1-200 rpm with automatic traversing which produced a prohibitively low rate af about 0.5 mm/min of polished rod. The spindle speed was progressively

increased up to 600 rpm, and the traversing

mechanism modified to produce 16 mm/min

of polished rod.

The predominant surface feature visible following this polishing procedure was a pronounced helical ripple, coincident with the path

traversed by the flame across the surface of

the rod. This feature was recognized in earlier

work on flame-polished sapphire but was not

found to prohibit the achievement of high

strength. Barnes and McCandless [8] used a

multiple jet burner to produce a broad flame

which heated a substantial length of rod and

effectively removed the surface ripple; this

expedient was not used in the current work.

Nevertheless, it was considered desirable to re- duce the effect as far as practicable and the

traversing pulleys were changed to reduce the

pitch of the helix to 0.013 mm. Reduction of

the pitch of the helix also reduced the depth

of the furrows between the spiral traces and

was therefore expected to increase the strength

of the rods. The rate of production of polished

rod was thus reduced to 8 mm/min (about

half of the rods produced in this work were

polished at this rate and half at 16 mm/min)

experience has indicated, however, that if necessary, higher polishing-rates (over 25 mm/

min using higher spindle speeds) are practicable.

e. "Optimum" Polishing Conditions

Good quality polished surfaces were obtained

from a wide range of polishing conditions so that the concept of an optimum set of conditions was not appropriate to the process [17]

It was evident that a balance had to be es- tablished between the amount of heat supplied

to the rod and the various sources of heat ex- penditure such that an optimum amount of

molten ALOs was maintained at the surface.

The depth of the molten pool during polishing

was not known with certainty but was estimated to be about 0.05 mm. The importance

of heat conduction through the rod was demonstrated by the fact that it was found necessary

to start polishing at a short distance (~2 mm)

below the top of the rod so that heat was con- ducted away from the molten zone equally in

both directions. When starting from the top

of the rod, a flame that would normally produce

a polished surface would melt the whole cross- section of the rod and form a progressively

larger droplet as the flame traveled downward,

because of heat conduction away from the

flame in only one direction.

f. Summary of Conditions Used for Flame Polishing

of Rods and the Operating Procedure

The following conditions were adopted for

the production of the majority of the polished

rods used in this work as a result of the ex- perience described above.

(i) Gas pressure SkN/m^ with methyl alcohol vapor addition.

(ii) Jet size 0.32 mm bore (hypodermic

Jet No. 23).

(iii) Jet-to-rod distance about 6 mm: the

218
exact position being determined by

visual observation of the rod surface.

(iv) Rod rotation speed 600 rpm; flame

traversing rate 0.013 or 0.026 mm per

revolution.

The molten ALOs formed a collar round the

rod about 0.5 mm wide by about 0.05 mm deep

under these conditions. In operation, conditions (i), (ii), and (iv) were established, the

flame was then brought to impinge on the rod

2 mm from the top, the traverse mechanism

engaged, and the jet-to-rod distance adjusted

to produce the desired degree of polish.

3.2. General Observations on the Polishing

Process and Comparison of the Results

Obtained on Different Rod Materials

During polishing, the rod acted as a lightpipe for the intense illumination generated by

the molten AI2O3 and light scattering reflections

were produced at any discontinuity within the

rod or at its surface; "flaws" were thus readily visible. Examinations of rods during polishing was therefore the most sensitive means

of assessing surface perfection and freedom

from gross flaws, both internally and at the

surface. It was found that if too much AI2O3

was melted during the polishing process, coarse

ripples and "necking" of the rod occurred

in the extreme, the rod would resemble a string

of beads. There was therefore a tendency to err on the side of under-polishing especially for

the 0.25 mm diameter specimens, to avoid loss

of the rod by complete melting of the cross section. In fact it was not possible to "overpolish"

these small rods; (i.e., produce sub-surface

voids), as is possible for larger diameter rods,

which would tolerate greater depths of molten

material without failure.

Examination of the rods from the initial polishing work showed evidence of "foreign" particles in the surface which often corresponded

to the edge of fracture faces in rods subjected

to strength tests. An ultrasonic cleaning procedure (following Mallinder and Proctor) was

subsequently used prior to polishing; a detergent solution, a water rinse, and finally an alcohol rinse was used and a marked improvement in the surface appearance and in the

strength of the polished rods was observed.

The "foreign" particles had obviously originated in dust and loose grinding debris present

on the surface of the as-received ground rods.

Both centerless-ground and commercially

flame-polished rods were used in the initial

work, and it was found that it was easier to

judge the correct polishing condition when us- ing a centerless-ground rod than when using a

previously polished one. It was thus more difficult to produce a high-quality polish on previously commercially-polished rods. Also, some

debris appeared to be sealed into the surface

in the commercial polishing operation and was

not easily removed in further polishing. Experience showed that the best quality polish, and

hence the highest strength rods, was attained

by arranging that the amount of liquid material

was sufficient to melt through all pre-existing

surface irregularities in a single pass of the

flame; multiple polishing rarely improved the

strength of the rods unless they were greatly

under-polished on the earlier polishing attempt.

Many rods were found to contain internal

flaws (bubbles, inclusions, etc.,) which could

not be removed by polishing. These internal

flaws were suspected as sources for fracture

initiation in the rods [11] and therefore the position of the flaws (rendered visible by the

illumination of the flame during polishing) was

mapped out for several rods which were subsequently broken in bending. Reassembly of the

broken fragments showed that in most cases the fracture faces corresponded to the previously located and "mapped" internal flaws.

When polishing the few Czochralski-grown

rods included in the studies, it was immediately

obvious that it was a much better quality

ALOs material—no internal defects could be detected during polishing. Unfortunately, because

of the scarcity of this material at the time, the

polishing was performed somewhat cautiously

to avoid loss of material or void formation by

overpolishing. The measured strength of these

rods (Section 3.4.e) was not especially high,

despite the high internal perfection, and further examination showed that their surfaces

had been slightly underpolished so that all

traces of machining damage had not been re- moved.

When polishing the Tyco c-axis sapphire filaments, it was found to be relatively easy to

improve on the as-grown surface (and hence

on the strength). However, it was not possible

to achieve strengths as high as those measured

in other material because the Tyco material

tested contained many internal voids [13], some

of which appeared to intersect the filament

surface. In some samples, there was also a ripple in the surface due presumably to a growth

instability. It was therefore easy to melt and

reform a more perfect surface but the internal

structure, containing many defects, was un- changed. In subsequent testing, only a relatively

moderate but nevertheless significant increase

in the average strength of the filaments was

measured. As a result of this work, a flame polishing process has been studied by Tyco as part

of the efforts to improve these filaments with

considerable success [15] and may soon become

part of the normal production process.

Interesting and highly significant differences

in behavior were observed between sapphire

and ruby during flame polishing, which were

219
not completely understood at the time. In sapphire and ruby obtained from the same source

and grown by identical techniques (except for

chromium addition in ruby) [18], there were

significantly less internal defects in the ruby

crystals. Also, because of the difference in optical properties, the thermal gradient on either

side of the flame was more evident with ruby, so that correct positioning of the flame and observation of the molten zone was greatly facilitated. In practical terms, this meant that

high quality polished ruby rods could be produced by unpracticed personnel after only a few

minutes instruction, whereas to produce a good

polish on sapphire material required considerably more experience. Therefore, a significantly

higher average strength was invariably found

in polished ruby compared with polished sapphire; even if the surfaces were similar, the

scatter in strength was usually less for ruby

because of the reduction in internal defects.

Study of ruby was initiated because of re- ports that the high temperature creep resistance was higher in ruby than in sapphire [9,

19] and because of the possible differences in

compatibility behavior between the two materials and Cr-containing metallic alloys. Apart

from the differences in strength arising from

internal defects referred to above, no significant

difference in the mechanical behavior between

flame-polished ruby and sapphire were found.

Furthermore, there was no detectable difference

in behavior in compatibility studies. Electronprobe analysis indicated that the surface layer

of the ruby rods was in fact depleted in Cr to

a depth which corresponded closely to the depth

of molten liquid during polishing. It appears

therefore that the surfaces of flame polished

"ruby" rods may differ only slightly from those

of sapphire rods.

Over a three-year period of studies with ruby

rods, it was observed that the red color asso- ciated with the chromium doping could vary

noticeably from one rod to another, even within

a given batch. All the material had a nominal

doping level of 0.05 w/o CrgOs but the color

could vary from a pale pink to quite a deep red.

No systematic difference in behavior which

could be related to this variation was detectable

either in the polishing procedure or in the re- sulting mechanical property measurements.

3.3. Surface Features of a-ALOs Rods with

Varying Degrees of Flame Polish

Figure 3 shows the surface of a centerlessground sapphire rod and illustrates the amount

of surface roughness which has to be melted

and recrystallized during flame polishing. It can

be seen that damage fragments have typical dimensions of 0.025 mm in the direction of the

surface as well as in depth. Figure 4 shows the

Figure 3. The surface of a centerless-ground sapphire

rod.

surface of a centerless-ground rod after one

pass of the flame under grossly "under-polished" conditions. The remnants of the grinding

damage, although smoothed and rounded, are

still evident and only slight increases in

strength occur at this stage. The right-hand

end of the crystal in figure 5 has been given a

second pass of the flame and is well polished;

the surface at the left where the flame had been

rapidly removed, contains many bubbles which

seem to remain in the molten ALOs collar and

are swept down the rod by the liquid.

The surfaces of typical commercially flamepolished rods "as-received" are shown in figure

6. A large amount of grinding damage and

remnants of "dust" particles are still evident

but some improvements in strength would be

expected over that of a ground surface and

was indeed found. No pronounced ripple is vis-

» 4

I 50/im

Figure 4. The surface of a centerless-grownd sapphire

rod after one pass of a flame under grossly "underpolished" conditions.

220
Figure 5. The surface of a rod similar to figure U but

after the right hand end of the rod had been given a second pass of the flame under ideal polishing conditions. The left hand end of the rod is where the

flame was removed and contained bubbles apparently

swept down the rod by the pool of liquid.

ible on the commercially-polished rods, which

indicates that only a small amount of AlaO,,

was liquid during polishing (using a relatively

large flame [8] ) which appears to have been

insufficient to produce uniformly polished rods.

The surface of a typical rod flame-polished in

the present work under conditions described

earlier is shown in figure 7. The helical ripple

is the only visible feature on such a surface.

The varying width of the pitch of the helical

path is caused by slight changes in friction and

slipping in the pulley-driven flame-traversing

mechanism.

Other features which have been seen on the

surface of apparently well-polished rods are

believed to be related to the crystallographic

orientation of the underlying material. For ex- ample, distinct lines parallel to the rod axis oc- casionally form and disturb the normal helical

ripple; Mallinder and Proctor [17], and Popov

[7] observed similar disturbances suggesting

development of crystal facets. This behavior

was occasionally pronounced on Tyco sapphire

filaments, which often had a rounded approximately triangular or even hexagonal cross-sec- tion which could be accentuated by polishing. A

tendency toward the development of a noncircular cross-section was noted also for many

Verneuil-grown rods. The surfaces of a few

rods were replicated by standard techniques so that surface features could be studied at high

magnification in an electron microscope. One

of these rods clearly illustrate the development

of crystal features (fig. 8). These were not

common features on polished rods and they

probably arose due to an over-polishing condition, perhaps also with some thermal etching.

Figure 6. (a) & (b) The surfaces of typical commercially flame polished sapphire rods showing remnants

of grinding damage and surface contamination.

Figure 7. The surface of a sapphire rod flame-polished

under ideal conditions in this work.

221
Figure 8(a) shows hexagonal outlines of basal

planes normal to the emergency of the c-axis

from the surface of the rod, while figure 8(b)

is from the surface of the same rod at a position roughly 90° around the circumference and

shows an edge-on view of the intersection of

basal planes with the surface. The bend

strength of this particular rod was 4 GN/m^

but it is not known whether any of these features were present in the regions of maximum

stress during the test.

3.4. Results of Strength Measurement on

Flame-Polished a-AloO,, Rods

a. Bend Test Procedure and Fixtures

The degree of improvement in surface

smoothness of the flame polished rods was as- sessed by measuring the strength of polished

rods in bending and in tension. Since it was a major object of this work to develop simple

routine procedures to produce and evaluate

rods, a simple bend test was used to measure

their strength. Selection of the test specimen

and crystallographic orientation of rods was

essentially random (except for a few tests alFiGURE 8. Features on the surface of a sapphire rod

following flame-polishing, (a) shows outlines of a

faceted basal plane at the emergence of the c-axis

from the surface of the rod and (b) the edges of

similar features (i.e., intersection of basal planes

with the rod surface) at a position on the same sur- face approximately 90° around the rod circumference.

ready mentioned in which the position of internal flaws seen during polishing was noted so as to relate them to fracture origins). Rough

handling and contact between rods was avoided

after polishing but normal handling to transfer

rods to the test machine, etc., including occasional deliberate contact of the gauge section

with steel tweezers, did not cause noticeable

deterioration in the surface or catastrophic re- duction in the measured strength at room temperature.

It is well known that three-point bend tests

should be avoided for ceramic materials, since

only a single point on a rod specimen is subjected to the maximum stress. The strength of

a ceramic is usually governed by discrete isolated flaws, which are unlikely to occur at the

point of maximum stress unless there are many

closely spaced flaws. In the case of flame-polished crystals with only a few well isolated

flaws, it was expected that a three point bending test would give a wide range of results,

with some very high values. This behavior was

indeed found in such tests (12.5 mm span and

steel knife edges used) of sapphire rods, where

strength over 9 GN/m^ was measured. The

elastic deflection was very large in these tests

and the validity of the simple elastic equations

used to calculate stress may be in doubt [20].

With high-quality flame polished ruby, a strength of about 9 GN/m^ could be measured

frequently. This was clearly an unrealistic es- timate of the useful strength of these rods.

Testing in four point bending was therefore

adopted for routine assessment of the quality

of the surfaces produced in this work. All results quoted were obtained on a test fixture with

25.5 mm outer span and 12.5 mm inner span

(i.e., a line on the rod surface 12.5 mm long

was subjected to maximum stress) with 3 mm

diameter steel outer "knife edsres" and 0.25

mm radius inner "knife edges". The fixture was

carefully and precisely aligned to ensure that

each of the rigid center knife edges made contact with the rod at the same time. A small box

and plastic screens surrounded the fixture in

order to collect the broken fragments. Small

pieces of tissue paper were often draped over

the free ends of the rods to prevent excessive

fragmentation, which was usually caused by

impact of pieces of the rod against the fixture

following the initial fracture. The test fixture

was mounted between the crosshead and a compression load cell of an Instron testing machine,

testing being accomplished by moving the cross- head at a rate of 0.5 mm/min.

Large elastic deflections were imparted to

the high strength rods before fracture, so that

several minutes elapsed before failure in each

test. The possibility of error arising from nonlinear behavior at large deflections was considered in calculating the strength of the rods us222
ing only simple elastic bending theory, but it appeared that such errors, even with the strongest rods, were less than 10 percent and meant

that the calculated strength was in fact less

than the true maximum stress in the rods at

fracture [20] . The test fixture is shown in figure 9 with a flame-polished sapphire rod in

position under a load corresponding to an outerfiber stress of 3.5 GN/m^ (The camera angle

was slightly off" center in figure 9 to show the

arrangement of the knife edges.) Figure 10

shows a flame-polished ruby rod under test

when the outer-fiber stress was 4.4 GN/m^ the

average strength for all the flame-polished

Verneuil-grown material (sapphire and ruby

combined). Light reflection from the surface

ripples can be seen on the rod in figure 10, particularly at the left hand side; the unpolished

ends of the rods are also evident in figures 9

and 10.

b. Results: Verneuil Material

The results of room temperature bend tests

on Verneuil-grown sapphire and ruby and having different surface conditions, as described in

section 3.1, are summarized in figure 11. No

difference in behavior could be detected between

sapphire materials obtained from different suppliers and therefore no such distinction is

drawn in figure 11. It can be seen that, for any

given surface condition, the scatter in the data

is approximately the same, i.e., the strongest

rod is 2 to 3 times as strong as the weakest

rod in each set of data. However, the average

strength of the rods changes significantly with

changes in the surface smoothness ; the average

strength of the flame-polished rods being 10

times that of the rods with centerless-ground

Figure 9. A flame-polished sapphire rod undergoing a bend test at room temperature in the steel test fix- ture. The outer fiber stress in the rod at this time was 3.5 GN/m^

I 1

I cm

Figure 10. A flame-polished ruby rod undergoing a

bend test at roojn temperature showing the elastic

deflection in a rod at a calculated outer fiber stress

of U-U GN/m', which corresponds to the average

strength of flame polished Verneuil material {both

ruby and sapphire) produced in this work.

>

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•

AVERAGE STREI^jGTH

•

RESULTS) ^ • - 1 en •

•

11

RE

(13 LU • S r.

JLTSI

SAPPh < i

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1

1

(20

RESULTS)

;SULTS)

SAPPHIRE

(24

RESULTS

RUBY

(10

RESULTS)

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SAPPHIRE

110

RESl •

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••

•

•

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• 1 Q- >- <: DQ

t/^ ZD — en ?!

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CEI^TERLESS 1 COMMERCIAL UNDER POLISHED POLISHED GROUND 1 FLAME POLISHED BUT NOT ARER 1 POLISH VISIBLE PRE-CLEANED CLfANING SURFACE

, DEFECTS

> ^, —^ \. ^

RODS AS RECEIVED RODS FUME-POLI^HED IN THIS WORK

SURFACE CONDITION

Figure 11. Summary of data obtained in room temperature bend tests on a-AkOs crystals (ruby and

sapphire material) grown by the Verneuil technique

and having various surface conditions.

surfaces. Therefore, although flame polishing

does not change one of the disturbing features

of ceramic crystals, namely the wide scatter in

strength results from any given material and

surface condition, it does appear to provide a

223
means to almost guarantee a minimum strenprth

which is determined by the internal perfection

of the crystal. In other words, flame polishing

provides a means of attaining stress levels cor- responding to the intrinsic strength of the bulk

crystal.

As was mentioned earlier, a detailed examination of fracture faces to determine the origin of fracture (surface or internal) was not

made; however, the indications obtained from

rods in which the position of observable defects

was noted during polishing showed a close

correspondence with fracture fragments after

testing. These observations suijport the earlier

conclusions of Mallinder and Proctor fll] con- cerning fracture origins in well-polished rods.

It is significant that strengths of the order

of 7 GN/m- reported by Mallinder and Proctor

[11] were not achieved in 4-point bend testing

of sapphire rods in this work. This was initially

attributed to the possibility of handling damage and the relatively crude testing technique

used in the present work. However, since it was

subsequently found that high strengths could

frequently be attained in ruby single crystals

polished, handled, and tested in the same manner as for sapphire, it was clear that a careful

choice of test section was needed to attain freedom from internal flaws for the sapphire sDecimens. The relative freedom from internal flaws

in the as-received ruby rods resulted in more

perfect test sections being obtained, even with

random rod selection.

It must be emphasized that the "random" se- lection of rods referred to above does not mean

that all the polished rods were acceptable for

testing. Experience showed that gross flaws

near the surface of the rods could be easily

picked out during polishing and these rods were

then set aside for other studies. From the

strength data obtained (figure 11). it is clear

that this visual selection could readily remove

all defects corresponding to strengths below

about 2.5 GN/m^

Casual handling of rods, including contact

with steel, did not catastrophically affect their

strength at room temperature but contact with

materials of equal or greater hardness (other

rods or silicon carbide abrasive) did markedly

reduce the strength. Even casual handling, however, could seriously affect the strength of the

rods which were heated at 1000 °C for 1 hr

in air before testing. Presumably salts, dust, or other contamination stuck to the surface

during handling, and then fused or otherwise

interacted with the surface during the heat

treatment, thus reducing the strength. Rods

protected from contamination could be given

similar heat treatments without loss in

strength. Similar observations were made by

Proctor [10] on sapphire rods and also on high

strength glass surfaces ; the latter were particularly susceptible to damage by even the slightest contact or contamination of the surface.

c. Results: Czochralski Material

The crystallographic orientation of the few

Czochralski-grown sapphire rods was similar

to the Verneuil rods used previously. Unfortunately, as mentioned earlier, these rods were

slightly under polished but the freedom from

internal defects was obvious during polishing.

Two rods were long enough for testing in the

bending fixture and had strength of 2.98 and

2.93 GN/m=. Smaller fragments of these rods

were then tested in 3 point bending with a 12.5

mm span and strengths of 8.16, 4.00, and 3.75

GN/m^ were measured. It would be interesting

to study this material further, particularly in

tensile tests, because its high internal perfection should permit the achievement of ex- tremely high strengths, possibly as high as the

best AloO, whiskers.

d. Results: Tyco Sapphire Material

The bend strength of 0.5 mm diameter Tyco

sapphire rods (cut from essentially continuous

filament) was determined on 2 batches of material obtained at different times. Batch 1 was

from the early production material (1968) and

the strength was low but reasonably uniform.

This material contained many internal voids,

some of which appeared to have faceted sur- faces and many were close to, or exposed at, the

rod surface; a typical rod is shown in figure

12. Batch 2 was specially selected from production material about 12 months later (1969) and

had a higher average strength than batch 1,

but also had a very large scatter in strength.

This material also had a high void content, but

the distribution was different than batch 1, the

voids being mostly in a zone just below the

surface but away from the core of the rod.

Figure 12. The surface of a typical Tyco savphire fila- ment (hatch 1) in the as-grown condition.

224
7X 10

10'

10»

A AVERAGE STRENGTH

AS GROWNFLAME

POLISHED,

AS FLAME GROWN POLISHED,

BATCH 1

SURFACE CONDITION

BATCH 2

Figure 13. Room temperature bend strength

of Tyco sapphire filaments from 2 different

hatches, with both as-grown and flame-polished surfaces.

Flame-polishing improved the average

strength of both batches of rods, but also increased the scatter in the strength data (fig.

13). A few large flaws in batch 1 were not

completely eliminated by polishing but otherwise the surfaces of the Tyco rods appeared to

be identical to those obtained on the centerlessground rods (fig. 7). However, because of voids

immediately beneath the polished zone, the lowest strength measured in all 4 sets of data was

similar (about 1.3 GN/m^). The highest

strengths measured in batch 1 compares favorably with the average strength of the flamepolished Verneuil sapphire rods (fig. 11), and

was presumably determined on a relatively

void-free length of rod (detailed examination

of the void distribution was not made). Also,

the highest strength in batch 2 is comparable

with a good flame-polished ruby rod and is indicative of freedom from serious flaws, at least

in a small volume of the rod. It is noteworthy

that a strength of 5 GN/m- was measured on

a rod in the as-grown condition, again indicating the potential for the achievement of high

strength in these filaments if growth flaws can

be avoided.

Because of the large scatter in the bend data

and the many internal flaws in the Tyco rods,

no tensile testing of flame-polished Tyco filaments was performed. With the material as described above, it was expected that flamepolishing would have an insignificant eff'ect on

the tensile strength, since tensile failure is dominated by internal flaws (Sec. 3.4.f). However,

recently produced Tyco filaments are much improved in internal perfection and it has been

shown that for such crystals, flame-polishing

can significantly improve the tensile strength

[15].

e. Elevated Temperature Bend Testing of

Flame-Polished Rods: Verneuil and Tyco Material

Plastic deformation in AI2O3 single crystals

at elevated temperatures has been known and

studied in both sapphire and ruby for several

years [9, 21]. Basal slip has been observed

above 900 °C [9] and rhombohedral slip about

1200 °C [22], in suitably oriented specimens,

and deformation twinning, both basal and

rhombohedral, has been observed over a wide

range of temperatures [21, 23]. Wachtman and

Maxwell found that creep deformation by basal

slip could be induced in commercially flame polished sapphire and ruby rods at temperatures

above 900° under low tensile stresses (< 50

MN/m^) [9]. There was no apparent difference

between the creep yield stress for flame polished or unpolished rods [9] but since the fracture stress at room temperature was low for

both types of rod, the quality of the commercial

flame polish was evidently poor. Since highquality flame-polished surfaces resulted in large

increases in the fracture stress at room temperature in the present work, it was of interest

to determine the mechanical properties at elevated temperature.

The bend strength of flame-polished ruby

rods and of as-grown Tyco filament (both 0.5

mm diameter) was therefore determined using

a test fixture identical in geometry to that used

for the room-temperature tests but with sintered alumina knife edges and superalloy supporting fixtures. The fixture was surrounded by

a Nichrome-wound split-tube furnace capable

of operation up to 1100 °C in air. Tests were

conducted at temperatures from 900 to 1100 °C

(measured by a thermocouple fixed between the

center knife edges of the fixture) in an atmosphere of argon/air (an argon flow was maintained through the furnace in an effort to re- duce oxidation damage to the metallic parts

of the rig). The crystallographic orientation of

the ruby rods was random (c-axis in each case was about 60° to rod axis) and the Tyco rods

had the c-axis parallel to their length.

The results of the elevated temperature testing are shown in table 1 ; the room temperature

strength quoted for each set of data is the

average strength of 10 rods from the same

batch as that used for the elevated temperature

tests. The onset of plastic deformation was observed in the ruby rods at 1000 and 1100 °C as

non-linearity in the stress strain curve. One

ruby rod fractured at 1000 °C without evidence

for deformation, no doubt because of unfavorable orientation of the slip direction in this

crystal. All the Tyco filaments were completely

elastic up to fracture, and no evidence for plastic deformation could be detected in the fragments. This was expected because of the un- favorable orientation of the basal planes in

these filaments and the previously reported ob225
Table 1. Bend strength of flame-polished ruby rods (Vemeuil) and as-grown Tyco

sapphire filaments (both 0.5 mm diameter) tested in an air/argon atmosphere

at various temperatures

Temperature

Bend strength (GN/m^)

Flame polished

ruby

Plastic*

deformation

As-grown

Tyco sapphire

Plastic*

deformation

20 4.13 0 3.31 0

900 2.07 0 1.43 0

900 2.07 0 1.12 0

1000 1.54 + 1.27 0

1000 1.49 0 0.76 0

1100 1.12 + 0.85 0

1100 0.79 + 0.55 0

*+= slip and/or deformation twins observed "1 optical examination

0 = no slip or twinning observed ^at. 100 X m polarized

J light

servations of rhombohedral slip only at 1200 °C

[22] . Two of the ruby rods were plastically deformed under essentially constant load (one

at 1000 °C and 1.5 x GN/m^ the other at

1100 °C and 1 GN/m^) ; there was a slight load

reduction as deformation proceeded until a permanent deformation of several per cent was

induced into the rods. The load was then re- moved and the bent but unbroken rods were

recovered from the furnace and examined for

evidence of basal slip and deformation twinning.

The data in table 1 shows that the fracture

(or yield) stress for the flame-polished rods is

at least an order of magnitude greater than

that measured on rods with imperfect surfaces

[9] ; the strength ratio between perfect and imperfect surfaces is greater even at high temperature than at room temperature. The high

stresses withstood by these rods before and during deformation suggests that dislocations nec- essary for plastic deformation are nucleated at

the surface, presumably at flaws or stress concentrations. Therefore, the same mechanism

which is responsible for the high fracture

stresses in the rods at room temperature

(namely the freedom from surface defects) is

also responsible for the high yield stress for

plastic deformation at high temperatures. The

strength/ temperature relationship for the ruby

rods is greater than that found for sapphire

whiskers with highly perfect (thermally

etched) surfaces [3], but not as great as found

for rods with imperfect surfaces [24, 26].

Within the scatter in data, the strength/temperature relationship is the same for the Tyco

sapphire as for the ruby; the strength at any

given temperature is lower for the Tyco material, in proportion to the lower room temperature strength of these rods as compared to

the ruby crystals.

f. Tensile Strength Measurements on Flame-Polished

Sapphire Rods: Verneuil Material

A cursory examination of fracture surfaces

suggested that fracture frequently originated

at internal flaws, as was found by Mallinder

and Proctor [11, 12]. The maximum stress in a bend test is only at the surface of the rod,

and the bend strength of flame-polished rods

may not be a realistic indication of the strength

of the rods measured in tension, where the

entire cross-section is uniformly stressed. Se- vere internal defects close to the center of the

rod would be relatively ineffective flaws in a bend test. Since many potential applications

for these rods would require them to carry a

direct tensile load, tensile tests were performed

on several rods as discussed below.

The apparatus used for tensile testing of

short sapphire rods is shown in figure 14. An

Instron machine with carefully aligned rigid

cell-couplings (developed for whisker testing

[27] ) was laid horizontally and rods were ce- mented to carefully aligned rigid anvils with

an epoxy resin. (Epon resin 815 (Shell) and

Versamid 140 polyimide resin (General Mills)

50/50 by weight were used and cured at 90

°C in about 1 hr under an ultra-violet lamp.)

The horizontal arrangement greatly facilitated

alignment of the specimen and enabled this

alignment to be maintained during curing of

the cement. An attempt was made to use as large a gauge length as possible but it was

also necessary to have an adequate length

(—'25mm) immersed in the epoxy in order to

prevent 'pull-out' from the grip.

The results obtained using this arrangement

226
Figure 14. Apparatus used for tensile testing of

flame-polished rods 0.5 mm diameter and up to 10 cm long.

are shown in table 2. It is evident that the

rods indeed possess high tensile strength, but

examination of fracture surfaces showed that

internal defects were limiting that strength.

All test results are included in table 2, although good alignment was not achieved in

every case; this presumably caused some low

strengths to be recorded because high bending

forces are developed in this situation [28].

With the exception of tests numbers 4 and 7,

in which fracture occurred at visible surface

irregularities, all fractures initiated internally.

The rods used in these tests were taken from

the same batch of sapphire rods used to obtain

Table 2. Tensile strength of flame-polished sapphire

rods

Gauge length Strength

Test (mm) GN/m' Alignment

1 26 2 85 Good

2 6 > 2.2 (pull-out) Good

3 15 2.15 Good

4 25 1.93 Good

5 23 3.66 Good

6 5 2.35 Good

7 28 1.89 Good

8 8 1.78 Poor

9 20 2.50 Poor

10 18 3 52 Good

11 10 >3'.50 (pull-out) Good

12 4 3.02 Good

13 30 2.37 Poor

14 10 3.18 Good

the bend strength data in figure 2, in which

strengths between 2.12 and 5.51 GN/m^ were

measured. It is significant that the average

strength measured in tension was only 65 percent of the average strength measured in bending. These data gave further support to the

contention that the strength of the flamepolished material is limited by internal defects

and not by the surface perfection.

The strength measured in test no. 5, table

2, is probably the highest strength yet recorded

in tension in a ceramic crystal of this size

(23 mm gauge length, 0.5 mm diameter, breaking load 76Kg). This strength, and indeed the

average tensile strength, compares favorably

with the tensile strength of sapphire whisker

crystals of much smaller diameter and length

[29] . Furthermore, higher tensile strength has

recently been reported in flame-polished Tyco

sapphire filaments [15] of smaller diameter

(0.25 mm) than those used in this work.

Tensile strength measurements were also

made on centerless-ground and commerciallypolished sapphire rods of similar dimensions to

those used above. The centerless-ground rods

had an average strength of 0.21 GN/m- (with

a range 0.08 and 0.32 GN/m- for 9 results)

and the commercially polished rods had an

average strength of 0.36 GN/m^ (a range of

between 0.26 and 0.45 GN/m^ for 4 results).

The average tensile strength of the centerlessground rods was 50 percent of that measured

in 4 point bending on similar rods, and that

of the commercially polished material was 35

percent of the strength measured in bending

(fig. 11).

The tensile strength described above was performed in the early phases of the work before

ruby rods were being used. Because of the

greater internal perfection of this material

it is thought that flame-polished ruby rods

would have higher tensile strength than sapphire but this has not yet been determined.

3.5. Further Comments

The ability by flame-polishing to produce

highly perfect surfaces on ceramic crystals has

been known for several years and yet has re- ceived little attention. For example, only re- cently have such studies been directed toward

exploiting the high strength of such crystals

in structural applications. Furthermore, it is

surprising that the technique has not been more

widely used in research studies of the mechanical behavior of ALOg and other ceramic crystals. It is again emphasized that, with reasonable care, the production of high strength rods

is a simple technique—certainly easier and

more straightforward than most mechanical

polishing procedures which never achieve comparable results.

227
On the other hand, a drawback for widespread industrial application of the flamepolishing process is that it is best suited to

simple rod or cylinders of small diameter,

although it has been claimed that the process

is applicable to slabs and other shapes [8].

With larger diameter specimens, there is a

severe thermal gradient throughout the cross- section of the specimen during polishing, caus- ing high internal strains, which in turn can

produce fracture. For diameters above about

3 mm, a large flame with a smaller inner hot

zone is desirable, so that the whole cross-section

and an appreciable length (say 3 to 5 diameters) of the rod on each side of the hot zone

is heated, thus avoiding severe thermal stress.

One solution is to flame-polish large diameter

(e.g., 5 mm) rods within a muffle furnace.

Attempts have also been made to polish a

flat surface on 10 mm diameter x 1.5 mm thick

discs of single crystal Al^Or; using the smaller

flames used for rods. High thermal gradients

were inevitable; thermal strains caused visible

basal deformation twins to form before surface

melting had occurred and spread rapidly across the disc at the point of impingment of the

flame. By preheating similar discs to 1200-

1500 °C during polishing, it was possible to

achieve some surface improvement but the

technique has not yet been perfected.

4. The Effect of Annealing

Klassen-Neklyudova [30] first demonstrated

that the strength of machined sapphire specimens could be increased by annealing at elevated temperatures. Since that time, similar

results have been obtained by Tomilevski [31]

,

Charles and Shaw [25], Davies [32], and by

Heuer and Roberts [26] , all of whom employed

Verneuil-grown crystals. This phenomena will

first be summarized and then discussed in

detail:

1. There is a threshhold annealing temperature below which strength increases

cannot be observed. Since the strength

increase results primarily from a reduction in severity of surface damage [26],

this threshhold temperature undoubtedly

coincides with the onset of gross material

transport, either by surface diffusion or by evaporation-condensation. (It is implicitly assumed in this paper that room

temperature fracture in sapphire is a purely brittle process and does not involve plastic deformation in any critical or important sense.)

2. The magnitude of the strength increase

is a function of the annealing temper-

' Davies employed tensile testing, all other workers [24-26, 30, 31] employed the bend test.

ature, the annealing atmosphere, and the

crystallographic orientation. (The as- machined strength also varies with

orientation, due to hardness and wear

anisotropy affecting the severity of the

surface damage imparted during machining of the specimens. [25]

3. Annealing affects the temperature dependence of the strength between room

temperature and 1000 °C. The temperature dependence is anomalous in sapphire, in that a strength minimum is

observed between 300 and 600 °C in as- ground specimens (as first reported by

Jackman and Roberts [24] ) but occurs

at '-SOO °C in annealed samples [25, 26,

32] . Inasmuch as the explanation for this

anomalous temperature dependence is

still controversial, it will not be discussed

any further here.

4. Although annealing can lead to substantial strength increases, strengths do not

approach those obtained for flame-polished specimens with near-perfect surfaces. On the other hand, annealing is a simpler process than flame-polishing and

is much more suitable for complex-shaped

parts.

In general, results of the several investigators who have studied the effects of annealing

in sapphire are in reasonable, although not perfect, agreement. For example, Klassen-Neklyudova [30] reported 30 percent increase in the

strength of 90 ° sapphire (c-axis normal to the

rod axis) following annealing for 5 hours at

1000 °C, while 0° specimens, which were about

twice as strong as the 90° specimens as- machined, showed little, if any, strength in- crease following this treatment. Tomilevskii

[31] reported a 30-40 percent strength increase

for 0° and 90° samples annealed at 1300 °C,

while Charles and Shaw found a similar

strength increase after 1200 °C; furthermore,

0° rods were about 1.8 times as strong as 60°

rods but 0° rods did not increase their strength

after this annealing. On the other hand, Davies

[32] found a strength increase of —'250 percent at 1200 °C for 30° rods, while Heuer and

Roberts [26], using 45° rods, obtained '-'30

percent strength increase for annealing temperatures of 1325-1525 °C but little strengthening at 1250 °C. These latter workers, however,

employed annealing temperatures up to 1800

°C, and found the strength increases to be

greater as the annealing temperature increased.

Since Heuer and Roberts also studied the effects of surface finish [machined versus mechanical polishing] and annealing ambient,

their results will be presented in more detail.

Figures 15 and 16 [taken from ref. 26] show

the effects of annealing temperature and atmos228
• ANNEALING ATMOSPHERE-DRY OXYGEN. • TWO SPECIMENS.

900 -

800

700

600

g 500

100-

= 300 QO

RANGE OF UNANNEALED

CRYSTALS.

200-

100

AIR {

-AIR

1000 1200 1400 1600 ANNEALING TEMPERATURE "C

DRY

HYDROGEN

1800

900

800

700

E 600 z

500 -

(C 400

300 -

O 200

100

A ANNEALING ATMOSPHERE-j DRY OXYGEN. '

1,035

t

RANGE OF UNANNEALED

CRYSTALS .

1000 1200 1400 1600 ANNEALING TEMPERATURE 'C.

DRY ^mOROGEN

18 00

Figure 15. Dependence of the room-temperature modulus of rupture of machined corundum crystals on

annealing temperature. The atmosphere for the an- nealing was dried oxygen except where otherwise

stated, 24 hours were used for all temperatures but

1250 °C, where the annealing time was 6 hours.

phere on the room temperature strength of

machined and mechanically-polished sapphire

respectively. With machined specimens (fig.

15), quite high strengths were achieved, and

the maximum strength obtained increased with

increased annealing temperature. Annealing

in hydrogen at 1800 °C did not produce as much strengthening as annealing in oxygen

at the same temperature. Much greater scatter

was observed with mechanically-polished and

annealed crystals (fig. 16) but strengthening

could certainly be achieved on occasion. (Although the mechanically-polished specimens

before annealing also showed greater scatter,

the room temperature strengths had usually

been less than 0.6 GN/m^ [26].) Heuer and

Roberts also determined the temperature dependence of strength of a group of machined

and mechanically-polished crystals annealed

together at 1800 °C for 24 hours in dry oxygen.

* Heuer and Roberts [26] also investigated the influence of pre- machining (i.e., boule) annealing (at 1800 °C) and found it not

to have any e(iect on the strength of either machined or mechanically-polished samples. This is an important piece of information

relevant to item 1 above, i.e., the strength increase achieved during

annealing results from a reduction in the surface damage introduced during preparation of samples from the larger boules. However,

Besson [33] has found annealed sapphire windows to be less sus- ceptible to failure in high pressure equipment than unannealed windows, and has ascribed this to a reduction of internal strains, rather than to reduction in damage introduced during preparation

of the windows.

Figure 16. Dependence of the room-temperature modulus of rupture of mechanically polished corundum

crystals on annealing temperature. The atmosphere

during annealing was dried oxygen except where

otherwise stated; 24 hours were used for all temperatures but 1250 °C, where the annealing time was 6 hours.

For machined specimens, strengthening was

observed at all temperatures below 1200 °C,

for mechanically-polished crystals, strength in- creases at liquid nitrogen and room temperature were realized but both the maximum

strength and the scatter of strength seemed to

decrease with increase in testing temperature.

Annealing increased the strength of machined crystals more than it did those with

mechanically-polished surfaces. For example,

strengths at -196 °C were about the same for

the two types of specimens after annealing,

whereas before annealing, the mechanicallypolished crystals were stronger.

Heuer and Roberts concluded that the main

strengthening effect of annealing was to modify

surface damage, as shown in figures 17 and

18. Although the detail observed on machined

specimens after annealing was difficult to interpret, changes in gross surface roughness

were obvious and were confirmed using a Talysurf machine (fig. 19). Decrease in surface

roughness was first detected following annealing at 1400 °C, and the surface became significantly more smooth as the annealing temperature was increased ; subgrain boundaries

could be seen after annealing at temperatures

439-921 O - 72 - 16 229
Figure 17. Surface of machined specimen as received

(reflected light, normal incidence)

of 1600 °C and above. This "threshold temperature" for surface improvement is consistent

with measurements of surface diffusion kinetics

(by grain boundary grooving) in alumina single crystals, which yield very low values

(<10-" MVsec.) for D^, the surface diffusion

coefficient below 1400 °C [34] ; also, LEED

measurements, show structural transformations on (0001) surfaces of alumina to occur

above 1250 °C in vacuum [35].

The annealing in hydrogen (1800 °C) was

more effective in decreasing the gross surface

roughness than was annealing in oxygen. Furthermore, machined specimens as-received were

initially translucent, and annealing had the

effect of increasing the light transmission. The

I fj. m

/i m

lO.ZSmm I

Figure 19. Profiles of surface roughness as determined

by a Talysurf machine.

a. Machined specimen, as received.

b. Machined specimen annealed for 24 hours at

1800 °C in oxygen.

c. Machined specimen annealed for 24 hours at 1800 °C in hydrogen.

d. Mechanically-polished specimen.

Figure 18. Surface of machined specimen after 2U hour

annealing in dry oxygen at 1800 "C (reflected light,

normal incidence).

specimens annealed in hydrogen (1800 °C),

thus, became almost transparent. In view of

these results, the weakening effect of hydrogen

annealing is not understood, but may be as- sociated with the chemical etching known to

occur when alumina is heated in hydrogen

above 1300 °C [36].

The changes that occur on mechanicallypolished sapphire surfaces after annealing have

been discussed extensively by Heuer and Roberts in a separate publication [37]

It may be concluded that air or oxygen an- nealing is an effective means of strengthening

machined alumina single crystals.

5. Conclusions

Both flame-polishing and annealing can be

used to improve the surface finish of alumina

single crystals, and thereby improve the

strength. Flame-polishing in particular is a

process using relatively simple techniques and

230
equipment (available in most laboratories) for

routinely producing highly perfect surfaces.

The strengths of such crystals with their nearperfect surfaces is generally limited by flaws

in the interior of the material and not at the

surface—the general case for ceramic materials. It is thus possible to obtain specimens

for measurement of the "intrinsic" mechanical properties of the bulk crystals.

Annealing, while yielding much more modest

strength increases, may be more suitable for

structural parts of complex shape.

The authors express their thanks to A. P.

Levitt of the Army Materials and Mechanics

Research Center, and to R. A. Schmidt of the

Naval Air Systems Command, whose support

of research projects enabled much of the work

described in this paper to be accomplished.

They are also grateful to Dr. D. J. Barbour

of Univ. of Essex (England), for providing

them with a translation of reference 30.

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[3] Brenner, S. S., Mechanical behavior of sapphire

whiskers at elevated temperatures, J. Appl.

Phys. 33, 33-39, (1962).

[4] Mehan, R. L., and Feingold, E., Room and ele- vated temperature strength of a-AUOj whiskers

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[5] Marsh, D. M., Stress concentrations at steps on

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[11] Mallinder, F. P., and Proctor, B. A., The

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[12] Mallinder, F. P., and Proctor, B. A., Preparation

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133 (John Wiley and Sons, Inc., 1970).

[151 Hurley, G. F., and Pollock, J. T. A., Tyco, Inc., (a) Progress Report to XVIIth Refractory Composites Working Group Meeting, Williamsburg,

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(b) Pollock, J. T. A., Continuous Flame-Polishing of

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[16] Noone, M. J., Feingold, E., and Sutton, W. H.,

The Importance of coatings and the preparation

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ASTM, STP No. 452, pp. 58-89, (1969).

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[18] Keig, G. A., (Union Carbide Corporation), Private Communication, (1968).

[19] Chang, R., Creep of AhO, single crystals, J. Appl.

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[20] Vrooman, D. L., and Ritter, J. E., Non-linear

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184, (1970).

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(1967).

[23] Heuer, A. H., Deformation twinning in corun- dum, Phil. Mag. 13, (122), 379-393, (1966).

[24] Jackman, E. A., and Roberts, J. P., The Strength

of single crystal and polycrystalline corun- dum, Phil. Mag., Ser. 7, 46, 809-811, (1955).

[25] Charles, R. J., and Shaw, R. R., Delayed failure

of polycrystalline and single-crystal alumina.

Report No. 62-RL-3081M, General Electric

Company, Research Laboratory, Schenectady,

N.Y., (1962).

[26] Heuer, A. H., and Roberts, J. P., The influence

of annealing on the strength of corundum crystals, Proc. Brit. Ceram. Soc. 6, 17-27, (1966).

[27] Mehan, R. L., Testing Fine Filaments and Whiskers, in Ceramic Fibers and Fibrous Composite

Materials by H. W. Ranch, Sr., W. H. Sutton,

and L. R. McCreight, pp. 27-48, (Academic

Press Inc., New York, N.Y., 1968).

[28] Roberts, J. P., and Watt, W., Selected Government Research Reports (British), Vol. 10

"Ceramics and glass", pp. 31-50, (1952).

[29] Mehan, R. L., and Hertzog, J. A., Mechanical

properties of whiskers, in Whisker Technology,

Levitt. A. P., ed., pp. 157-195, (John Wiley &

Sons, Inc, 1970).

[30] Klassen-Neklyudova, M. V., The mechanical properties of corundum crystals, J. Tech. Phys.

(USSR) 12, 519-551 (1942).

[31] Tomilevskii, G. E., Effects of heat treating synthetic corundum crystals at 1300 °C on mechanical properties, Trudy Inst. Krist. Akad.,

Nauk USSR, 8, 341-354 (1953).

[32] Davies, L. M., The effect of heat treatment on the

tensile strength of sapphire, Proc. Brit. Ceram.

Soc 6, 29-35, (1966).

[33] Besson, J. M., Private Communication, Lab. de

Phys. de Solid, de la Fac. des Sci. de Paris,

France, Equipe de Rech. CNRS, 1970.

231
[34] Robertson, W. M., and Ekstrom, F. E., Impurity

effects in surface diffusion on aluminum oxide

in Kinetics of Reactions in Ionic Systems, Vol.

4, Materials Science Research, Gray, T. J. and

Frechette, V. D.. eds., (Plenum Press, Inc., New

York, N.Y., 1970).

[35] French, T. M., and Somorjai, G. A., Composition

and surface structure of the (0001) face of

^-Alumina by low-energy electron diffraction,

J. Phys. Chem. 74, 2489-2495, (1970).

[36] Mercier, J., Dopant transfer in heteroepitaxial Si

layers on sapphire substrates, J. Electrochem.

Soc. 117, 812-814, (1970).

[37] Heuer, A. H. and Roberts, J. P., Thermal etching

of single-crystal corundum. Trans. Brit. Ceram.

Soc. 65, 219-232, (1966).

Discussion

Gielisse: You have promised us a few brief

comments on the difference between flame polishing sapphire and ruby.

Noone: For instance, we found that in our

lab scale tests the ruby material which we used

was much easier to polish than sapphire. A

technician, after only 10 minutes of instruction,

could produce a million psi by flame polishing

of ruby material; with sapphire this was not

the case. We always found much less internal

defects in the ruby material than was present

in the sapphire. IDifferences between sapphire

and ruby were expected in compatibility studies

with metals but on closer examination by electron probe techniques, it turned oiit that the

surface of a flame polished ruby rod was depleted of chromium. Chromium apparently

evaporated from the liquid zone during the

flame-polishing process so that the "ruby" rod

was actually identical to sapphire within maybe

2 mils of the surface. Overall, in the 300

results that I discussed (fig. 11) the average

strength of the best batch of ruby was around

790,000 psi. This illustrates, I think, more than

anything else the gross effect of internal defects

on the strength of these materials.

WiEDERHORN: I am interested in the strength

decrease at elevated temperatures for flame

polished rods. I noticed that there was no mention made of the temperature region between

room temperature and 900 °C.

Noone: That's right. We haven's yet covered

the whole temperature range with flame-polished material unfortunately. We tested at

room temperature and above 900° where we

expected to see the onset of basal slip which

was critical for the proposed applications of

these rods as reinforcement at high temperatures.

K. Smith: I think that you perhaps were

forming a suboxide in your hydrogen furnace.

Were you getting a fully oxidized alumina?

Noone: What do you think Arthur? I call

on my coauthor to answer.

Heuer: If one goes by the color of the alumina, there is no evidence of reduction.

Wachtman: Do you know or care to speculate what the state of stress is in your flame

polished rods? Do you expect to have a compressive or tensile stress?

Noone: We certainly do not know in detail.

I don't expect to have a high stress, but the

surface must be in tension since it solidifies

and cools over an already solid and relatively

cold core. Wachtman: One would think that the point

at which you could no longer relieve stress

would be reached not far below the melting

temperature. At this point the surface might

still be hotter than the interior and you would

have some tension.

Noone: I suspect you could probably relieve

stresses all the way down to about 1200 °C

in these small diameter rods. We haven't done

the kind of experiments which John Pollock

will describe a little later in which he split the

rod down a center axis and looked for resultant

curvature. We've looked at a lot of rods under

polarized light and seen evidence of growth or machine induced strain which appears to be

removed by polishing. I think the correct an- swer is that we don't really know what's going

on inside.

Lange: What are your thoughts either

about dislocation motion or twinning prior to

fracture, or at fracture in plain polished specimens?

Noone: At room temperature I see no evidence for plastic flow or deformation twinning

in the rods.

Lange: So fracture more than likely was not

nucleated by either dislocations or twinning.

Noone: Fracture at room temperature we

think is nucleated by the internal defects, predominantly bubbles.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machimnc and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).

232
Healing of Surface Cracks in Ceramics

F. F. Lange

Westinghouse Research Laboratories, Pittsburgh, Pennsylvania 15235

Grinding, cutting, thermal shocking, impacting and rough handling all introduce surface cracks that degrade the strength of ceramics. Within the last year, it has

been shown that, once introduced, these cracks can be eliminated by either resintering

the damaged ceramic component for the case of oxides or oxidizing the component for

the case of materials such as SiC.

Results of crack healing experiments, performed on AUOa, ZnO and SiC will be

reviewed. The technical implications of these results will be discussed as related to abrasive machining.

Key words: Abrasive machining; ceramics; crack healing; sintering; surface damage.

1. Introduction

Shaping of most ceramic materials into en- gineering components requiring strict tolerances are currently carried out by what can be

most suitably described as abrasive machining.

This process involves the surface removal of

one brittle material by another, harder, brittle

material. Surface removal by abrasive machining leaves a damaged surface which, for many

cases, impairs the usefulness of the machined

ceramic component. Although the nature of the

damage can be described by many property

measurements, e.g., residual strain, dislocation,

density, electrical conductivity, etc., its appearance can best be described as a surface containing numerous microcracks and notches. These

flaws can be directly observed using both optical and electron microscopy. Indirect observations of these flaws can be made using both an

oil dye penetrant and strength measurements.

Elimination of this damaged surface can be

carried out by several techniques: (1) chemical

polishing; (2) surface stressing; (3) lapping

(polishing) ; and (4) heat treatment. Chemical

polishing has been found successful for removing surface damage from some single crystals,

[1]^ polycrystals, [1] and glasses [2]. Surface

stressing (chemical strengthening, glazing,

etc.) alters the damaged surface to a surface

that supports a compressive residual stress,

thus strengthening the material. [3, 4] Lapping

is an abrasive machining technique that re- moves the severely damaged surface and leaves

a less damaged surface. Heat treatment, the last

of the techniques listed above, is the subject of

this review.

Prior to 1970, several investigators [5, 6]

reported that the strength of several ceramics

could be increased by a heat treatment. One of

these heat treating techniques was flame polishing. [7] In this report, it was suggested that

' Figures in brackets indicate the literature references at the end

of this paper.

' Hot-pressed ZnO, Lucalox AI2O3 (General Electric Company)

and 80 percent dense Crystar SiC (Norton Company).

surface damage (e.g., microcracks) could be

eliminated by heat treatment. During the past

year, it was shown [8, 9, 10] that deep surface

cracks can be eliminated from several polycrystalline ceramic materials. The object of this

paper is to review these recent results and to

suggest a simple technique for removing sur- face damage left by abrasive machining.

2. Healing of Deep Surface Cracks

A large density of surface cracks can be con- veniently introduced into ceramic specimens by

thermal shocking them from an appropriate

quenching temperature into water. Cracks introduced in this manner usually transverse the

tensile stress zone formed by the thermal gradients [11] and, therefore, their depth is approximately 1/4 of the specimen width. Both

the pattern and depth of these surface cracks

can be revealed by an oil dye penetrant. The

crack pattern appeared similar to that which

appears on a crazed glaze. These deep cracks

drastically decrease the material's strength.

This thermal shocking technique was used to

introduce deep surface cracks in specimens of

ZnO^ (1/8 by 1/8 by 1 1/2 in), [8] AW,' (1/8

by 1/4 by 2 in) [10] and SiC ' (1/4 by 1/4 by

2 in) [9] which were quenched from 400 °C,

400 °C and 600 °C, respectively into water at

room temperature. Different sets of specimens

were heated at 1100 °C for the ZnO, 1700 °C for

the AI2O3 and 1400 °C for the SiC; each different set was heated for a different period. The

first two temperatures were chosen because

powders of the respective materials can be sintered at these temperatures. [12] The temperature for SiC was chosen because it was known

that an oxide surface layer rapidly forms at

1400 °C. [13] All ambients were air (the ZnO

specimens were contained in covered ZrOa

boats).

After cooling, the flexural strength of each

specimen was determined by four point loading. Flexural strengths were also determined

233
Table 1. Flexural strengths of 'as-cut' and thermal shocked specimens

Material

ZnO

AI2O3

SiC

Strength

'as-cut'

psi {MN/rn^)

9600±12% (66.2)

35,800 ±3% (246.8)

19,000 ±7% (131.0)

Strength

thermal shocked

psi (MN/rn^)

3900±12% (26.9)

12,800± 7% (88.3)

9000±11% (62.0)

— A _

.2

ZnO, Heated

to 1100°C

_L

10 20 30 40 Time at Temperature, h50

J \ L

J Vacuum 1^*^ Strength

Oxidization Behavior

J L

40

30

20

10

0 10 20 30 40 50 60 7 0 80 90 100 Time (hr) at 1400°C

Figure 1. Fractional recovery of initial flexural

strength of both thermal shocked ZnO speci- mens heated to 1100 "C and thermal shocked

AUOz specimens heated to 1700 °C. (Normalized to as-cut strength (no thermal shock, no

heat treatment).

for both "as-cut" specimens (no thermal shock,

no heat treatment) and thermal shocked specimens (no heat treatment). These are summarized in table 1. One set of SiC specimens was

also heated in vacuum to 1400 °C for 56 hours.

The oxidation rate for the SiC material was

obtained from weight measurements before and

after the specimens were heated. Average grain

size measurements were made for all sets of

ZnO and AI2O3 specimens.

Figure 1 summarizes the normalized flexural

strengths of the thermally shocked, reheated

oxide specimens. Figure 2 shows both the flexural strengths and weight gain (oxidation rate)

of the thermally shocked, reheated SiC specimens. Table 2 lists the grain size measurements.

Both figures 1 and 2 show that all three materials regained all (or a significant portion for

AI2O3) of their strength after a particular heat

treatment. These data suggested that the deep

surface cracks were eliminated. For both of the

oxides, this was confirmed by an oil dye penetrant, i.e., surface cracks that could clearly be

Figure 2. Static oxidation behavior and strength* re- covery of 80 percent dense, thermal shocked SiC

specimens heated to 14.OO °C. * (Normalized to as-cut strength)

seen after thermal shock could no longer be

observed after a certain period of heat treatment. By comparing the data represented in fig- ure 1 with the grain size measurements presented in table 2, it can be seen that complete

recovery in strength occurred for the ZnO despite considerable grain growth, whereas grain

growth due to prolonged heating affected the

AlaOj strength recovery. This latter result has

been discussed by Lange and Radford [10].

For the SiC, the surface cracks could still

be seen (several surfaces were polished prior

to thermal shock) under an oxide surface layer.

Fracture surface observations revealed some

traces of an oxide surface layer for SiC specimens heated for short periods. This suggested

that the oxide filled and healed the deep cracks.

This fact was confirmed by the low strength

found for the specimens heated in vacuum.

Lange [9] has shown that such an oxide bond

has little affect on the strength recovery of

hot-pressed, high strength SiC.

These data and observations clearly show

that surface cracks can be either eliminated

(for the case of the oxides) or made inopera234
Table 2. Grain size* of heated oxide specimens

Time at temperature AU.O3(1700 °C) ZnO(1100 °C)

hr

0 43 /xm 6 Atm

7 44

10 -- 10

20 11

25 52

35 14

50 54

*In micrometers, determined by linear analysis and multiplied by 1.5.

tive (for the case of SiC). Qualitative oil dye

penetrant experiments show that heat treatment can eliminate surface cracks in polycrystalline MgO. Evans and Davidge [14] have

recently shown that crack-like surface flaws on

reaction-sintered SigNi can be made inoperative

by the formation of an oxide surface layer. For

certain conditions, they have shown that this

results in a strengthening effect. Thus, this

author believes that crack healing by heat

treatment is a phenomenon general to most

ceramics.

3. Heat Treatment of Abrasively

Machined Surfaces

Since deep surface cracks can be healed, microcracks introduced by abrasive machining

should also heal with heat treatment. To test

this hypothesis for the materials reported

above, viz., ZnO, AI2O3 and SiC, two sets of

specimens of each of these three materials were

diamond cut; one set was heated for the re- spective maximum period and temperature re- ported above. Four point flexural strengths

were then determined for each set of specimens. The results are shown in table 3.

These results show that for both ZnO and

SiC, a strengthening was observed. The reason

for the decrease in strength for the AI2O3

is most likely the increased grain size. This

has been discussed elsewhere. [10]

Using scanning electron microscopy, both asTable 3. Strength change for diamond cut specimens

after heat treatment

Material Heat

treatment

Percent average

strength change*

ZnO 35h, 1100 °C + 17

AI2O, 50h, 1700 °C -18

SiC llOh, 1400 °C + 10

*Table 1. Contains diamond cut strengths before heat

treatment.

cut and heat treated AI2O3 surfaces were ex- amined. Representative SEM micrographs are shown in figure 3. Figure 3(a) shows the rough

'as-cut' surface. The arrows on this micrograph show microcracks, grain-pullouts and

Figure 3. Scanning electron micrographs showing diamond cut surface of AUO3 prior to heat

treatment (a) and after heat treatment (b).

Heat treatment was 50h at 1700 °C.

235
sharp notches. The heat treated surface, shown

in figure 3(b), is much smoother, i.e., no transgranular microcracks were observed and all

edges and notches were rounded relative to the

'as-cut' surface. Thermal etching appears to

have occurred, thus revealing most grain boundaries.

These qualitative experiments and observations show that heat treatment can improve the

quality of an abrasively machined surface. It has also been shown here that the temperature

and time required of this heat treatment can be

chosen as either that required to sinter powders of the oxide or that necessary to form a coherent oxide film for the case of non-oxides.

For strength requirements, factors such as grain growth must also be considered. As it was

pointed out in the introduction, heat treatment

is not the only way of changing degraded sur- face properties, but it does represent a simple

technique for approaching this problem.

The author would like to acknowledge the

helpful discussions with both T. K. Gupta and

R. C. Radford, both of whom participated in

the original work. Support by the Office of

Naval Research, Contract NOOO14-68-0323,

NR 032-507 was appreciated.

4. References

[1] Stokes, R. J. and Li, C. H., Dislocations and the

tensile strength of magnesium oxide, J. Am.

Ceram. Soc. 46, 423-34 (1963).

[2] Hillig, W. B. and Charles, R. J., Surfaces, stress- dependent surface reactions and strength, p. 683 in High- Strength Materials, Ed. by V. F. Zackay, Wiley & Sons, New York (1965).

[3] Stookey, S. D., Strengthening glass and glass-ceramics by built-in surface compression, ibid, pp. 669-81.

[4] Kirchner, H. P., Gruver, R. M., and Walker, R.

E., Strengthening sapphire by compressive sur- face layers, J. Appl. Phys. 40, 3445-52 (1969).

[5] Heuer, A. H. and Roberts, J. P., Influence of an- nealing on strength of corandum crystals, Proc.

Brit. Ceram. Soc. 6, 17-27 (1966).

[6] Kirchner, H. P., Graver, R. M., Platts, D. R.,

Richel, P. A., and Walker, R. E., Chemical

strengthening of ceramic materials, Summary

Report, Contract NOOO19-68-C-0142, Naval

Air Systems (1969).

[7] Mallinder, F. P. and Proctor, B. A., Preparation

of high-strength sapphire crystals, Proc. Brit.

Ceram. Soc. 6, 9-16 (1966).

[8] Lange, F. F. and Gupta, T. K., Crack heating by

heat treatment, J. Am. Ceram. Soc. 53, 54

(1970).

[9] Lange, F. F., Healing of surface cracks in SiC by

oxidation, ibid, pp. 290.

[10] Lange, F. F. and Radford, K. C, Healing of

surface cracks in polycrystalline AUOs, ibid. pp. 420-21.

[11] Davidge, R. W. and Tappin, G., Thermal shock

and fracture in ceramics, Trans. Brit. Ceram.

Soc. 66, 405-22 (1967).

[12] Gupta, T. K. and Coble, R. L., Sintering of ZnO:I,

J. Am. Ceram. Soc. 51, 521-25 (1968).

[13] Lange, F. F., (unpublished).

[14] Evans, A. G. and Davidge, R. W., The strength

and oxidation of reaction-sintered silicon ni- tride, J. Mat. Sc. 5, 314-25 (1970).

Discussion

Rice: Two questions: First, do you feel the

small increase indicated in silicon carbide was

real? If so, do you feel this might possibly be

due to the fact that SiOa is formed in the

cracks? This might be a way of putting a compressive stress on the surface layer.

Lange: Could be. Whether or not it's real,

I really don't know. As discussed in the paper,

Si02 did form within the cracks. I also have

tried to heal deep surface cracks in hot-pressed

SiC. It was impossible to strengthen this thermally shocked material. Apparently the silica

bond was not strong enough relative to the degraded strength of the thermally shocked, highdensity silicon carbide.

Hockey: I'd just like to add that I've looked

at cracks associated with hardness indentations

in single crystal aluminum oxide in both the

unannealed and annealed state by transmission

electron microscopy. In the unannealed case,

the cracks always tended to propagate during

the thinning of the specimen. In the annealed

case, for example at 1000 °C for one hour, I often found dislocations associated with the

cracks created during indentation. These cracks

remained arrested during thinning.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

236
Flame Polishing of Flat Oxide Bars

P. F. Becher and R. W. Rice

U. S. Naval Research Laboratory, Washington, D.C. 20390

Some of the problems and limitations of flame polishing flat bars are discussed.

Results are presented for single crystal, a-AhO,, as well as for more limited trials of

MgAUOn, Ti02 and soda lime glass. The wide variability of strength is partially related

to variations in surface, but twinning also appears to be important in sapphire and

TiO:. Preliminary results on twin sources and their eflS'ect on strength of sapphire are

discussed.

Key words: Flame polishing; glass; rutile; sapphire; spinel; strength; surface characterization ; twinning.

1 . Introduction

High bend strengths have previously been

achieved by flame polishing materials such as ALO. [l-4y and SiO. [1]. However, these studies have typically been conducted on round

specimens. This paper presents results of an

initial study of flame polishing of flat single

crystal bars of a-ALOa, MgAL04, TiO, and

glass. These tests show more clearly the problems of flame polishing flat bars and the advantages of using round bars. Flat bars have

some advantages in revealing microstructure

and fracture mechanisms, and some reasons for

the differences in the flame polishes achieved

with these and initial results on fracture mechanisms are presented.

2. Experimental Techniques

Single crystal bend bars of a-ALO,, both

undoped and Cr-doped, having {1010}- and/or

{1123} tension surfaces with the tension axes

within 20° of <1120> and 20° of <4261>

(normal to (2025)), respectively, were flame

polished. In addition, bars of MgAl.Os and TiOa

single crystals and of soda lime glass were also

cursorily investigated. The flame polishing

utilizing an oxy-hydrogen torch consisted of

two techniques related to the amount of melt

formed. The first, "pool," method required a slow increase in temperature by moving the

torch back and forth over the length of the bar

until melting was initiated at the edges. Then

the torch was brought to close proximity to

the surface to form a definite melt pool, which

was then "pushed" along the surface.

The second "clearing" technique involved

prolonged heating of the bars to obtain a general clearing over the entire surface. A thin

layer of melt was apparently formed but was

barely, if at all, evident to the unaided eye.

' Figures in brackets indicate the literature references at the end of this paper.

' Indices of planes and directions based on c/a — 2.73 : 1 for

«-Al,Oj.

The "clearing" polish gave much smoother finishes, generally free of surface artifacts. Prior

to flame polishing, the a-ALOs crystal tension

surfaces were prepared by diamond grinding

parallel to or perpendicular to the tensile axis,

mechanically polishing (1 /j.m diamond paste)

or gas polishing [5, 6]. The other ceramic

specimens had as air-polished and/or ground

surfaces. The flame polished specimens

(--1.8x3.6 mm in cross section) were subsequently tested in 3 point bending at room

temperature over 1.27 cm. spans in an Instron

test frame employing a cross speed of 2.1 X

10"= mm/s. In addition, specimens were ex- amined optically before and after testing, and

in some instances, hot prosphoric acid was

utilized to delineate the microstructure (especially in the case of the a-AloO;, bars). The

sapphire samples were oriented with standard

laue back reflection x-ray techniques together

with the use of indexing tables [7]. Similar

methods were used for the analysis of twins.

3. Results and Discussion

3.1 Flame Polishing of a—AI2O3

The room temperature strength of the flame

polished sapphire and ruby bars having various

prior surface finishes are shown in table 1, together with data for these bars prior to flame

polishing. In general, flame polish strengths are greater than machined surfaces (all of which

have comparable strengths except for those

ground across the tensile axis). The improvements obtained by flame polishing, however,

are less than expected from previous studies

[1-4] ; this is particularly evident in the results

with the "pool" technique. The best overall

strengths are obviously obtained using the

"clearing" method which was most successfully

applied to ground surfaces rather than mechanically or gas polished ones. Cursory flame polishing trials by the authors,

using circular rods, demonstrated that greater

strengths were achieved in rods versus flat bars

237
Table 1. Room temperature strengths of a-AUO^ single crystal bars

[Modulus of Rupture, lO'psi {10''N/m^)'\

Sapphire Ruby

orientation orientation

Surface finish [1010]'=:<1120> ' [1123]':=; <4261> = [1123]';=<4261>'

Ave. M. R.

Range

Ave. M. R. Range

Ave. M. R. Range

I. Mechanical finishing

A. As-ground

1. T. A.' 25±7

14-36 (17)

24±6

14-35

(16)

2. T. A.' 53±2

35-73 (37)

55±6

49-66 (38)

69±3

62-78

B. Mechanically polished

59±4

42-72 (40)

II. Flame polishing

A. "Pool" technique

1. Ground i T. A.' 33±8

20-77 (23)

2. Ground || T. A.' 63±7(43)

50±14

(34) ^^-^

^38) 42-80 ^^^^ 41-83 34-73

3. Mechanically polished

47±8

33-67 (33)

4. Gas polished

B. "Clearing" technique

1. Ground || T. A.'

71 ±10

57-87

109 + 9

(49)

(75)

140 ±7

(97)

126±21

97-156 ^^^^ 92-150 135-145

2. Mechanically polished

160 ±90

92-232 (110)

3. Gas polished

109±20

89-140 (75)

' Orientation of tension surface.

^ Orientation of tensile axis.

' Tensile axis.

(compare values in table 1 with references

[1-4]. Although a difference in the crystal

stock characteristics might explain the lower

strength values for the bars (as compared to

rods), gas polishing of identical flat bars re- sulted in greatly improved strengths [6] as compared to flame polished flat bars. This indicates that the flame polishing process itself,

and not the initial bulk material, limited

strengths. A qualitative understanding of the

effects of flame polishing can be obtained by

observing its effects (microstructural and

strengthening) on these specimens. Some of

these are more evident in flat bars than in

circular rods.

a. "Pool" Technique

Three types of surface features characterized the flnishes obtained by the "pool" technique of flame polishing. First of these were

surface "waves," of the order of 0.5 cm. in

length, which resulted from non-uniform motion of the flame tip to avoid excessive melting.

Secondly, surface doming was observed which

occurred when a uniform but extensive melt

was formed. This thicker melt yielded cross sec- tions having curved polished surfaces as a re- sult of liquid surface tension (fig. la). Surface

facets generally formed in the region where

the melt was deepest, which was usually in the

center of the domed region or near "wave"

crest. This crystallographic faceting was attributed to solid state diffusion brought about

by the slow cooling which results from the melt

thickness and slower torch motion in the pool

method.

Specimens having "wave"-type surfaces tended to have somewhat higher strengths than the

238
Figure 1. Effect of specimen shape on melt

configuration during flame polishing by

the pool technique. Curvature in the sur- face is a result of liquid surface tension

during melting. The greater the amount,

and hence extent, of the melt at any one

time, the greater the degree of doming.

Flame polishing of flat surfaces accentuates this doming as shown in (a), while

circular cross sections tend to mask the

effect (h) where the surface before and

after polishing was approximately coincident as depicted here. In these two

cases, the maximum depth of the melt

(dotted line) and resultant contact angles

are illustrated as being equal.

other types produced by the "pool" technique.

In part, this resulted from the discontinuous

nature of these surface crests and the increased

likelihood of the maximum stress in three point

bending to be located away from these regions

which are likely to behave similar to domed

surfaces.

The domed surfaces and their associated

thicker melts were more likely to be subject

to subsurface bubble formation as a result of

solidification from the melt surface rather than

the underlying solid surface. As discussed by

Mallinder and Proctor [1, 2], bubbles or voids

act as a source of weakness; however, bubbles

were not generally observed in the present

specimens and few fractures originated at or near bubbles. Thus, bubbles did not appear to

be a primary cause of the low strengths obtained in these specimens. The domed surface

contour and surface faceting would be expected

to lead to stress concentrations ; however, their

radii of curvature was such that they did not

generally appear to explain the low strengths

observed. In addition, these surface features

were not continuous along the bar samples, and

failure often took place in sections devoid of

them. Furthermore, gas polishing [6, 8] was

utilized to remove surface irregularities from

a limited number of these flame polished

("pool" method) bars, yet the strength of these

bars was unaltered.

Twins were observed in these specimens, suggesting another source of structural weakness.

During cooling of these bars, surfaces are subjected to considerable thermal stresses as a re- sult of heating one surface. Twins have previously been reported to be introduced in

a-AlgOg during rapid heating and cooling [9,

10], and flame polishing-induced stresses apparently served to initiate twins especially at

regions of stress concentration, i.e., bar edges

and surface facets (fig. 2). As noted in figure 3,

faceted regions are particularly active twin

sources. Both basal and rhombohedral twins

were detected using optical and x-ray analysis.

Furthermore, there was evidence for slip at

basal twin tips (fig. 4) induced by flame polishing. The population of these twins (basal and

rhombohedral) was higher in ruby specimens

than sapphire bars, consistent with the observations of Levengood [11]. The results of his

study at room temperature indicated that the

stresses necessary to initiate twins decreased

with chromium additions. However, the present

results suggested that intersecting twin systems were a more significant influence on

strength than twin populations, as seen by the

similar strengths of ruby and sapphire flame

polished bars. Although twin-twin and twinsubboundary intersections previously reported

to result in crack initiation [12-14] were not

explicity associated with fracture origins in

the present study, cracks were often associated

with twins (figure 5), and twins were noted at

fracture surfaces. The fact that similar low

strengths occurred when nonparallel twins

were observed in fracture regions devoid of

any surface irregularities, further suggested

that twins were a definite factor. Also in support of this, the previously mentioned flame

polished bars that were gas polished but had

unaltered strengths were found to contain intersecting twins.

b. "Clearing" Techniques

The finest quality (i.e., flattest and most

featureless) surfaces were produced by the

"clearing" techniques and with moderate effort.

Seldom were comparable surfaces achieved

with optimum operation of the "pool" method.

This was attributed to the much smaller quantity of melt produced, which rapidly solidified,

inhibiting surface doming and facet formation.

The fact that a melt did form, even though

difficult to detect, is seen in that clearing

239
Figure 2. Twinning in a-AltO, as a result of flame polishing.

a. As flame polished surface exhibiting wave crests (dark, faceted areas) and linear surface features paralleling the interference fringe bands. Polarized transmitted light.

b. Chemical etching of surfaces similar to (a) delineates intersecting twins which correspond to linear struc- tures in (a). Interference phase contrast, reflected light.

Figure 3. Twin source on flame polished surface. At

slightly higher magnification, it is seen that heavily

faceted surface regions serve as an origin of tapered

twins occurring on parallel planes. Transmitted light.

Figure 4. Evidence of slip preceding basal twinning in

ruby. Etching of flame polished surface with hot

phosphoric acid reveals etch pits in twin interface

and a dense row of dislocation etch pits in advance

of twin path. These are suggestive of basal twinning

involving dislocation glide processes on the basal

plane. Interference phase contrast, reflected light.

240
Figure 5. Fractures associated with flame polishing

twiris. Cracks are observed in twin regions or along

twin boundaries after testing and the crack at the

twin tip above extends to the fracture origin. Chemically etched; interference phase contrast, transmitted

light.

"against the grain" of grinding striations

could not be accomplished. It was evident that

the grinding stria halted the flow of the thin

melt layer and good results were achieved only

when the melt was allowed to flow along the

grinding grooves. In bars ground perpendicular

to their lengths, this resulted in problems of

melt overlap, partial polishing, and increased

thermal stresses. Further problems were en- countered with mechanically or gas polished

surfaces where the lack of difference of clarity or contrast of the surfaces resulted in poor

control of this technique. This lead to formation of thick melts and related problems ; however, when properly applied, clearing of

mechanically polished surface yielded the highest strength of any flame polished bar (230,000

psi, 155x10^ N/m'').

The high quality finishes of the cleared sur- faces were consistent with the improvements

in strengths of these as compared to the previous "pool" polished bars (table 1). In addition, the cleared bars were generally free of

twins and when twins were observed, they oc- curred only on parallel planes (fig. 6). This

observation further indicated that formation of

intersecting twins was a factor in the low

strengths of the pool polished bars.

3.2. Flame Polishing Other Oxides

a. MgALOs

Several attempts were made to flame polish

AlaOs-rich spinel single crystal bars, primarily

by the pool technique. Here again, the resultant

strengths were widely scattered (table 2).

Cracks from thermal stress, facetting and subsurface bubbles (which were more frequently

associated with surface doming in these bars

than in a-ALOs) were also major factors in

the low strengths. Nevertheless, the strengths

of the flame polished bars were substantially

improved over those of as-ground or mechanically polished samples. However, the higher

strengths associated with the best finishes

appeared to be low as compared to the improvements obtained with a-Al.Os. Precipitation of excess AI2O3, which has been shown to substantially reduce the strength of spinel [17, 18]

,

may be an important factor in these bars. This

was suggested by the cloudy nature of the more

domed specimen surfaces of the lower strength

bars. Subsequent bars that were annealed and

slowly cooled prior to or after flame polishing

exhibited even lower strengths (table 2), further indicating that precipitation was a factor

in the flame polished strengths. Hence, partial

precipitation may have been limiting strengths, so that greater improvements might be obtained in stoichiometric spinel.

Table 2. Strengths of various ceramic oxides [Modulus of Rupture, lO'psi (lO'N/m')']

Mechanically Flame

Material As-ground polished polished

MgAU04 (Single crystal) 34±2 (23) —- 42±6' (29) 7-122'

85 ±23' (59) 60-122

26±6' (18) 19-35

38 ±13° (26) 28-47

TiOi (Single crystal) 13±1 (9) — 2.7± (1.9) 1.7-4.4

Soda lime glass 14±1 (10) 17±3 (11) 15±2 (10) 12-18

' All spinel specimens

' Range of values.

' Spinel specimens devoid of bubbles, cracks, or other obvious defects.

* Spinel which was air annealed at T > 1100 °C prior to flame polishing.

'Flame polished spinel with subsequent air anneal, T > 1100 °C.

241
Figure 6. Parallel twins in high strength, flame

polished sapphire bar. These twins, accentuated by

etching, are observed to be appi'oximately perpendicular to the twin axis. The two twins at far right

intersect fracture surface in region of fracture

origin. Interference phase contrast, reflected light.

b. TiOj

Flame polishing of a limited number of

ground rutile bars was attempted, essentially

by the pool technique. In comparison to similar as-ground samples, the flame polished

strengths were quite low (table 2). In polishing TiOo, the melt was more difficult to control,

resulting in part to a tendency for the melt to

skip along the surface. This yielded surfaces

having a mixture of wave, dome, clear, and unpolished surface regions. The air-annealed

(yellow) TiOn specimens also became a dark

blue or black, indicating that considerable re- duction had occurred during flame polishing. In

addition, most specimens cracked, some to the

extent that testing was meaningless. During

examination of fractured test specimens intersecting, narrow bands, whose optical behavior

and etch structure was indicative of twinning,

were often found to be associated with fracture

origins (fig. 7). These results indicate that

TiO. was not a likely candidate for further

attempts at flame polishing and suggested that

flame polishing induced twins which were a source of fracture initiation.

c. Soda Lime Glass

For cursory comparison, a few ground soda

lime glass bars were flame polished primarily

using the clearing technique. The strengths of

these were only slightly higher than those

Figure 7. Fracture surface of flame polished TiOi

crystal. Fracture origin (to left of center, upper

edge) is associated with linear configurations suggestive of twins. Note further that these twin

features occur on intersecting planes. Chemical etch

(fuming H3SO4), interference phase contrast, re- flected light.

242
obtained with mechanical polishing, table 2. It was noted that the flame polished surfaces

frequently contained bubbles which generally

were sites of fracture initiation. Improvements

in flame polishing strengths were seen to be dependent on the quality of previous surface

finishing. Smoother prior surfaces might have

eliminated bubbles which apparently resulted

from gases trapped in grinding stria and

cracks.

3.3. Applications of a—AI2O3

Briefly, the two flame polishing techniques

employed with the flat sapphire bars resulted

in significant strength variations. The detection of intersecting twins solely on samples

prepared by the pool technique indicates that

twinning as well as surface irregularities

played an important role in the lower strengths.

As compared to flat bars, circular rods polished

by the pool method would not be expected to

be subjected to such pronounced doming. Although the surface tension would still tend to

draw the liquid up, the radius of curvature of

the rod would substantially reduce the effect

in comparison to flat surfaces (fig. 1). The

ability to rotate circular rods (and the melt)

is also felt to be influential in maintaining a more uniform shape and this would also be

conducive to forming thin melt layers, thus

further minimizing the problems associated

with thick melts. In this respect, it was previously found that the best quality finishes on

circular rods was achieved with thin melt

layers [15]. Further, flame polishing was

undertaken in that study in a furnace designed

to maintain the rod at a fixed elevated temperature prior to and after passing the rod through

the torch flame [15]. This feature, as well as control of withdrawal rates, apparently substantially reduced axial thermal stress and

eliminated fracture or cracking problems. The

rotation of the rods during polishing would

also greatly reduce cross sectional (or radical)

stresses which were a factor in the polishing

of flat bars which are polished on one surface

only. The polishing controls for circular rods

generally eliminate "wave crest" features

found in flat bars ; however, under poor polishing conditions, the circular rods exhibited similar structures [15]. Surface features resembling twins were also produced in circular rods,

but were avoided by better control afforded by

the polishing process [15]. These twins did not

appear to occur on intersecting planes consistent with the lack of stress concentration

regions (i.e., edges and also facets) in circular

rods. The curved surface also reduces the probability of faceting or other surface irregularities being favorably situated for a favorable

twin orientation.

It is worthwhile to briefly consider application of these flame polishing techniques to

flat surfaces as disks and plates. While the

clearing technique of flame polishing is definitely preferred and is applicable to narrow

flat surfaces, it would be difficult to apply on

large surfaces because of the requirement for

more uniform surface heating in comparison

to the pool technique. Use of a multiple flame

torch or torches with flame dimensions (i.e.,

width of hot portion of flame) comparable to

surface dimensions, might reduce this problem.

In addition, initiation of polishing at the center

of the surface and traversing out (e.g., in a

spiral path) may also reduce thermal stresses

in large pieces in comparison to starting from

the outer portion or linear motion of the flame

along the surface. However, control of the

clearing technique is easier on a ground sur- face but does not work well when the flame

motion traverses across grinding stria, thus

making spiral-type paths difficult.

The pool technique could be applied, but besides the associated surface finish problems,

overlapping of the melt could cause difficulties.

This problem would be accentuated by larger

radii of torch travel, as the previous melt

would have sufficient time to solidify, but might

be compensated by increasing rotation speeds

with increasing radii of travel. Thermal stressing would also become a greater problem as the

piece size increases, but as in circular rods,

may be substantially relaxed by supplemental

heating. Similar techniques have also been applied during welding of ceramics with success and indicate that maintaining the workpiece at

1000°-1500 °C was necessary to prevent stres- ses induced during the process [16]. In this re- gard, the success of supplementary heating will

be limited by the problems of eliminatincr thick

melts and surface faceting. Thus, while this

work gives some guides to flame polishing flat

objects, it is regarded as a difficult task, usually

with limited benefits.

4. Summary and Conclusions

The findings illustrate that flame polishing

can yield improved strengths of flat surfaces,

particularly those with narrow, rectangular

cross sections. However, it has been shown that

extreme care must be taken to avoid surface

irregularities (facets, etc.) and other defects

(bubbles and cracks). By employing a technique which forms only a thin melt layer, intersecting twin formation and/or faceting can

be avoided in a-ALO., and strengths are substantially improved over conventionally machined bars. Less perfect flame polish finishes are usually obtained when thick melts are formed resulting in little or even negative

improvement of strengthening.

243
The utility of flame polishing- flat bar-shaped

samples is not so much in achieving strength

improvements; much greater strengths are us- ually achieved in flame polished round rods

[1-4] and gas polished flat bars [6]. The

above technique does, hov^^ever, present some

advantages in studying fracture and twin behavior (especially in a-AloO:, and TiO,,). It also

illustrates that flame polishing is not necessarily a simple process of melting out surface

flaws. Recent work by Hockey [19], showing

that many twins and dislocations are introduced in the surface of ALO;, by mechanical

machining, indicates that flame polishing is

also melting out these defects. Hence, the high

strengths achieved by flame polishing are not

unequivocal evidence for a flaw mechanism

of failure. Further, this work shows that during and after flame polishing twins and dislocations can be reintroduced, depending on

specimen and polishing parameters. Results,

while not conclusive, clearly suggest that twinning may be an important factor in the room

temperature strength of sapphire, as can precipitation in ALOg-rich MgAlaO,.

5. References

[1] Mallinder, F. P., and Proctor, B. A., The strengths

of flame polished sapphire crystals, Phil. Mg.

13, 197 (1966).

[2] Mallinder, F. P., and Proctor, B. A., Preparation

of high-strength sapphire crystals, Proc. Brit.

Ceram. Soc. 6, 9 (1966).

[3] Kirchner, H. P., Gruver, R. M., Platts, D. R., Rishel, P. A. and Walker, R. E., Chemical

strengthening of ceramic materials, summary

report. Contract N00019-68-C-0142, Ceramic

Finishing Company (1969).

[4] Noone, M. J., and Heuer, A. H., Improvements in

surface finish of ceramics by flame polishing

and annealing techniques, this volume.

[5] Witter, D. E., and Palmour III, H., PIMAX

tables for alpha alumina. Engineering School

Bulletin 84, North Carolina State University

(1967).

[6] Rice, R. W., Becher, P. F., and Schmidt, W. A.,

The strength of gas polished sapphire and

rutile, this volume.

[7] Rice, R. W., The eff^ect of sputtering on surface

topography and strength of ceramics, this

volume.

[8] Schmidt, W. A., and Davey, J. E., Preparation of

smooth, crystalline, damage-free, sapphire su''- faces by gaseous etching, this volume.

[9] Barber, D. J., and Tighe, N. J., Electron microscopy and diff'raction of synthetic corundum

crystals. I. pure aluminum oxide grown by the

verneuil nrocess, Phil. Mag. 11, 495 (1965).

[10] Klassen-Neklyudova, M. V., Mechanical twinning

of crystals. Consultants Bureau, New York

(1964).

[11] Levengood. W. C, Quantitative Aspects of Twinning Deformation on Sannhire and Ruby Single

Crystals, nn. 93-97 in Crystal Growth, H. S. Reiser (editor), Pergamon Press, New York

(1967).

[12] a. Heuer, A. H., and Roberts, J. P., The influence

of annealing on the strenerth of corundum

crystals, Proc. Brit. Ceram. Soc. 6. 17 (1966).

b. Heuer, A. H., Deformation twinning in co- rumdum, Phil. Mag. 13, 379 (1966).

[13] Parikh, N. M., Fracture mechanisms in polycrystalline nonmetallic materials, Final Report,

DA-19-066-AMC-288(x), IIT Research Institute (1966).

[14] Becher, P. F., and Palmour III, H., High-temperature deformation of alumina double bi- crystals, J. Am. Ceram. Soc. 53, 119 (1970).

[15] Popov, S. K., The growth and uses of gem-grade

corundum crystals, Part II. The use of corun- dum in making artificial fibers, pp. 135-152 in

Growth of Crystals, Vol. 2, A. V. Shubnikov

and N. M. Sheftal (editors). Consultants

Bureau, New York (1959).

[16] Rice, R. W., Welding of Ceramics, NRL Report

7085, U. S. Naval Research Laboratory, Washington, D.C. (1970).

[17] Grabmaier, J. G., and Falckenberg, H. R.,

Strength of flame-fusion-grown magnesiumaluminum spinel, J. Am. Ceram. Soc. 52, 648

(1969).

[18] Rice, R. W., Press forging and recrystallization

of ceramic crystals, submitted for publication.

[19] a. Hockey, B. J., Observation of plastic deformation in alumina due to mechanical abrasion.

Am. Ceram. Soc. Bull. 49, 498 (1970).

b. Hockey, B. J., Observations on mechanically

abraded aluminum oxide crystals by trans- mission electron microscopy, this volume.

Discussion

Heuer : You were saying that during flame

polishing of sapphire, twins were formed.

These were interacting with or intersecting

other twins and formed small cracks. Was it these cracks that gave rise to the fracture at

low stresses?

Becker: Yes, flame polishing (especially

when thick melts occurred) frequently re- sulted in the formation of twins. Cracks were

observed at some twin-twin interactions and

some fracture origins were closely associated

with intersecting twins or the associated

cracks, but the correlation of twinning and

specific fracture origins was not always clear.

However, it was clear that the strengths were

always low (comparable to as-ground samples)

when nonparallel twins were produced during

flame polishing. When only parallel twins were

formed during polishing, both low and high

strengths could be obtained. The possibility of

twin-sub-structure interactions yielding cracks

could account for the lower strengths, and evidence of this type of interaction was observed.

The presence of dislocations which appeared

to emanate from tips of basal twin tips,

somewhat complicates the picture of twinning

in fracture. Thus, although the exact mechanism (s) of failure are not fully resolved,

twinning is clearly important. This is shown by

gas polishing which greatly improves the sur244
face finish of sapphire and in the absence of

twins, substantially increases its strength.

However, gas polishing of flame polished

bars containing nonparallel twins results in

strengths comparable to as-flame polished

values.

Noone: I'm not sure that I understand the

terminology that you are using. In fact, it appears to me that you are not flame polishing.

I am particularly confused with this idea of

"clearing." Isn't that nothing more than what

we describe as annealing?

Becher: No, this isn't an anneal. I should

have brought out that a very thin melt is

formed using the clearing technique. This is

most obvious when one tries to use this method

in a direction normal to the grain of the grinding striations. In this case, the flow of the melt

is impeded and good polishes are difficult to

obtain.

Noone: The strength certainly doesn't reflect the kind of values which are accepted for

flame-polished materials. Strengths about 100,-

000 psi, for instance are normal values which

can be achieved by merely annealing.

Becher: This work in essence points out

that flame polishing of flat surfaces is much

more difficult. One is not able to control the

process nearly as well as in the case of a circular cross sectioned rod.

White: The fissures which you call twins

look very much like the compositionally induced twinning which we see as the cause of

similar effect?

Becher: Well, I can answer it this way.

These are definitely rhombohedral and basal

twins that were induced during flame polishing.

White: Is there a compositional difference

in this twinning?

Becher: From the matrix? I haven't looked

at it, but I don't believe so. H. J. Green : Were the bars polished on all

four sides?

Becher: No. We just flame polished the top

side. You get a little flame polishing on the

edges, however.

K. Smith : What were the dimensions of

those bars?

Rice: About 70 mils thick and 150 mils wide.

In addition I just wanted to amplify a little

bit on Mike Noone's previous question about

why we didn't see higher strengths for flame

polished flat bars. One of the things we discussed in more detail in the paper is the fact

that when you are flame polishing a flat bar

you are only flame polishing one side. Even if

you flame polish all sides, it still is normally

one side at a time. This generates thermal

stresses essentially in three dimensions. When

you flame polish a round bar, the thermal

stresses are primarily along the length of the

bar. The remaining ones are radial. I think it's

important to bring this out because these things

do play a role in causing lower strengths.

Heuer: One difficulty is that you are comparing three point strengths with Mike Noone's

four point tests. When Noone uses three point

testing he obtains 1.3 million psi reproducibly.

I think that you should use four point testing

in your work as this has become almost standard in many laboratories.

Rice: I think that four point bending has a

number of advantages and three point bending

also has its advantages. I feel that three point

bending is adequate or superior for most of the

tests we make, but it depends a lot on what you

are looking for, neither should be used exclusively for all purposes. The fact that Noone

made three point bend tests does give us one

very definite basis for comparison.' However,

the comparison also depends on the volume of

stressed materials to near the fracture stress.

Thus, in our three point bending of flat bars,

we are stressing a volume approaching that

stressed in Noone's four point bending of

round bars. Generally, I have found results

from three point bending in good agreement

with others testing the same or similar bodies.

K. Smith : Have you ever flame polished both

sides?

Becher: No, we haven't done that.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

439-921 O - 72 - 17 245

Continuous Flame Polishing of Sapphire Filament

J. T. A. Pollock*

Tyco Laboratories, Inc., Bear Hill

Waltham, Massachusetts 02154

Continuous flame polishing of nominal 2.5 X 10"^ m diameter single crystal sapphire filament oriented with the c-direction parallel to the filament axis has resulted in

considerable enhancement of the tensile fracture strength. Optimum increases of

approximately 1.1 X 10° N/m^ (1.6 X 10' psi) have been obtained on flame polishing

many lengths of filament exhibiting as-grown tensile strength of 2.2 - 2.8 X 10° N/m'

(3.2 - 4.0 X 10° psi). Flame polishing was carried out in a continuous manner at

6.3 X 10"^ m/sec using an oxygen/hydrogen flame. Data are presented which suggest

that that the enhancement is not entirely dependent on the production of a more

perfect AI2O3 surface. Increases in strength of approximately 5.5 X 10' N/m^ (8 X 10' psi) are reported for filament which has passed through flames not sufficiently hot

for surface melting to occur. Maximum strength enhancement is obtained when the

oxygen/hydrogen flame temperature is such that an axial molten zone two to three

times the filament diameter is produced at the filament surface. When the flame

temperature is too high, the geometric integrity of the filament is lost and an apparent

fall in tensile strength from the peak value is observed.

Metallographic evidence is presented indicating that the molten zone has a radial

depth of less than 6.5 X 10"" m. Scanning electron microscopy and related surface

analysis studies are reported which confirm that optimum polishing results in a smoother filament surface. Experiments to determine the state of stress in the filament

before and after polishing are described.

The relative contributions to the reported tensile strength enhancement of thermal

strain relieving, thermally activated atomic diffusion leading to blunting of possible

fracture sources and the creation of a more perfect AI2O3 surface are discussed.

Key words: Characterization; continuous fiame polishing; sapphire filament; strength

enhancements.

1 . Introduction

It is a well attested fact that the mechanical behavior of brittle materials is greatly

influenced by surface imperfections and also

to a varying extent by their state of internal

strain. Mere handling of freshly formed glass

results in a severe degradation of the mechanical strength Thus, in order to obtain

meaningful values of the strength of glasses, it

has been necessary to devise special handling

and measuring techniques [2] . The realization

that glass could exhibit high strengths has en- couraged studies of the high strength properties

of other brittle materials. Sapphire, a high

melting point oxide with high modulus and

hardness values [4.65 x 10" N/m^ (67.5

108 pgj) 2.93 X 10" N/m^ (2.8 x 10«

psi), respectively] has received increasing attention in recent years with a view to its use

in composite materials, where its thermodynamic compatibility with many metals at

high temperatures would be advantageous.

Sapphire whiskers (4 to 100 x 10-" m diameter) were investigated by Brenner [3] and

found to have strengths in tension in excess of

6.9 X 10« N/m^ (10« psi). The strength values

obtained were inversely related to the whisker

•Present address: Australian Atomic Energy Commission Research

Establishmen, New South Wales, Australia. ^Figures in brackets indicate the literature references at the end of this paper.

diameter, indicating that surface area and/or

volume was playing a role in determining the

sapphire strength. Reported strength to bulk

sapphire is considerably lower than those for

whiskers, bend strength values generally being

less than 6.9 x 10^ N/m= (10= psi) [4, 5].

The knowledge that surface condition can

play an important role in determining the

strength of brittle materials has led to the development of techniques for improving the sur- face condition of bulk sapphire. Morley and

Proctor [6] and Mallinder and Proctor [7] re- ported that considerable strength enhancement

was obtained on flame polishing sapphire rods

in the diameter range 0.75 to 2.75 X 10"' m.

Selected samples, free from internal flaws, ex- hibited strength in pure bending approaching

7.0 X 10" N/m^ (10« psi), and demonstrated

that the strength of bulk sapphire could approach that measured in whiskers. Noone and

his co-workers [8] repeated this work at a later

date and reported similar increases in strength.

The flame polishing technique used by both

groups was simplification of the method devised by Popov [9] who spent a large part of

his scientific life investigating methods for producing useable sapphire parts from verneuil

boules. His rather complex methods included

heating a centerless ground 5 x 10'^ m diameter sapphire rod, in a muflle furnace while held

in a rotating holder. An oxygen-hydrogen flame

247
is inclined and focussed at the sapphire rod, A

molten surface bead is produced and the rod is

withdrawn while rotating, thus allowing the

molten bead to circularly traverse the entire

surface area. The rods were then bent at elevated temperatures for use as guides and no

strength data were reported. Proctor and his

co-workers [6, 7] and Noone, et al., [8] discarded the muffle furnace, probably because

they used smaller rod diameters. The former

used a coal gas-oxygen flame and the latter an

oxygen-hydrogen flame with a methyl alcohol

vapor addition. Further details of the results

of these investigations will be included in the

discussion of the present work.

Recently, long lengths of high strength, single crystal sapphire filament have been produced by a melt-growth technique [10] . The as- grown tensile strength measured for 2.54 x

10"^ m diameter lengths of this material is in

the range 2 - 2.75 x 10« N/m^ (3 -4 X 10^

psi ) . Although substantial for a material in this

form, these values fall somewhat short of the

data reported for whiskers and flame-polished

rods. The aim of this study was to flame polish

lengths of this sapphire filament in an attempt

to raise its strength,

2. Experimental Procedure

2.1. Flame Polishing Technique

C-axis sapphire filament nominally 2.54 X

10"^ m diameter was available in long lengths

and was chosen for flame polishing. The problems associated with flame polishing material

of this geometry are radically different from

those found in polishing larger diameter, short

length rods. The rods have a rigidity which enables them to be accurately aligned and a substantial thermal mass in comparison to the filament. Also, the sapphire filament is in long

lengths and ideally the technique devised should

be one which would allow flame polishing in a continuous fashion. With the latter criterion in

mind, the simple setup shown schematically in

figure 1 was used.

A coiled length of filament, previously coated

with paraffin or starch to prevent self rubbing

degradation is supported on a circular rod at

the bottom of the apparatus. The filament is

fed through two Teflon bearings, which ensure

a degree of rigidity, and then through a small

pulley belt at the top. An oxygen-hydrogen microtorch is positioned with the nozzle tip (orifice size 7.5 X 10-*m) about 3 x lO ^m from

the filament. Gas flow to the torch is controlled

by two separate flow meters. The speed at which

the filament can move is continuously variable

and all the data reported here are for treatment

carried out at 6.3 X 10-*m/s. The filament is

not rotated in any way while being treated. No

Sapphire

filament

Figure 1. Schematic diagram of apparatus used to flame polish sapphire filament.

attempt was made to remove the paraffin or starch coating prior to flame polishing. It was

simply allowed to burn off in the cooler part

of the flame.

Flame polishing is carried out in the following manner. The belt puller is set in motion and

the flame started with only hydrogen (about

4 X 10" mVs) flowing into the microtorch.

Throughout the procedure the hydrogen flow

rate is held constant and the flame temperature

increased by allowing increasing flow of oxygen

into the mixing chamber of the torch. The ox- ygen/hydrogen flow ratios used, i.e., flame temperatures, vary depending on the filament diameter, but are generally in the range 2.75 to

1.5. Flow readings are taken from meters calibrated in standard cubic feet per hour (scfh)

of air and values reported here represent con- versions from this scale. They are not particularly accurate since allowance has not been

made for viscosity differences. This is especially

true of the hydrogen values. Since, however,

visual recognition of the optimum flame condition was achieved and changed with variation

in filament diameter, it was not considered

worthwhile making the readings more accurate.

They do, in fact, represent a qualitative measure of the flame temperature within each ex- periment and are useful as such.

Forty to fifty centimeters of filament were

treated at each flame temperature. The general

flame/filament area was observed continuously

using a stereomicroscope capable of magnification up to 40X. Usually the flame temperature

248
was increased until necking-down of the filament, the result of excessive surface melting,

was observed. The range of the gas flow ratios

used to treat the filament depended on the diameter of each length, which varied between

2.2 to 2.6 X 10-^ m. The larger the filament diameter, the hotter the flame needed to produce

optimum flame-polishing conditions.

The beginning and end of each length of filament, treated at a given flame temperature,

was marked with a wax pencil, and the length

carefully broken off as it cleared the puller belt.

These lengths were very carefully handled and

placed on foam rubber pads prior to layups

into tensile samples.

The tensile fracture stress of 2.5 x 10"^ m

gauge length samples was determined at a

strain rate of 8.3 x 10"^ sec-\ Three of four

samples were tested to fracture from each

piece of treated filament and an untreated

standard length. Additional samples of the filament were carefully set aside and used for

metallographic examinations with a view to

determining structural changes occurring as a

result of flame polishing.

Six flame polishing runs were carried out

and characterized in the above fashion, and

will be used as the basis for this paper.

2.2. Characterization

a. Internal and Surface Structure

The filament was examined for possible

changes in internal structure using transmitted

light optical microscopy. Samples were immersed in oil of refractive index close to sapphire (nn""^ = 1.760) and studied at magnification up to 750X. The primary aim of these

investigations was to determine the depth of

the molten zone during flame polishing.

Changes in surface roughness as a result of

flame polishing were detected by surface analysis using a Bendix profilimeter. The method

employed a chisel point probe capable of detecting surface roughness as fine as 2.5 x 10"^ m.

While it is realized that the device does not

absolutely profile the discontinuities detected

in a manner which would permit quantitative

conclusions about changes in surface microcracks, it was felt that a qualitative measure

of variation in the surface condition could be

obtained and would be evidence for changes in

surface features affecting the strength. Five

samples taken from flame polishing series VI,

were used. These samples, over a 5 x 10-^ m

long, were tabbed at one end in the form of a

flag. Each sample was placed in a groove on a

Teflon block and tested at 2.094 radian (120°)

intervals along a 3.18 x 10"^ m length using a

7.6 X 10-3 m length using a 7.6 x 10-^ m cutoff.

The filament surface was also studied at

various stages of treatment using Scanning

Electron Microscopy (SEM). A thin layer of

gold-platinum alloy was first vapor deposited

to nrevent charge buildup.

The single crystallinity of the filament, before and after polishing, was confirmed by obtaining x-ray precession patterns. Filtered

MoKa radiation, a precessing angle of 30°, and

a crystal to film distance of 6 x 10"^ m were

employed.

b. Residual Stress Since the filament is produced by a meltgrowth process [10] with its concomittant

solid-liquid interface temperature profile, the

thought was considered that the as-grown filament might be in a state of residual stress similar to that exhibited by tempered glass [11],

i.e., a compressively stressed outer shell con- taining an inner core in mild tension. This

stress pattern is advantageous since fracture

in tension, which in a brittle material is generally nucleated at the surface, requires that

the residual surface compressive force be first

overcome. It is possible that the flame-polishing

process causes a more extreme stress pattern

and thus contributes to the strength enhancement observed.

Petrographic examination of sections taken

perpendicular to the growth axis i.e., the optic

axis <0001> was difficult, due to problems in

sample preparation arising from the hardness

and small size of the filament. Samples 3.5 X

lO"'' m thick were prepared and examined for

evidence of radial stress patterns under polarized light in both conoscopic and orthoscopic

conditions.

Another method of determining the filament

stress condition was considered. Samples 7.5 x

10"^ m long were cemented lengthwise to glass

slides and polished to half their thickness, using

silicon carbide papers and diamond paste. On

dissolving the cement, the samples were then

examined using a stereomicroscope for evidence

of curvature, which would indicate the presence

of a nonsymmetric stressed state containing

longitudinal compression and tension components.

3. Results

3.1. Tensile Fracture Strength and

Relationship to Flame Temperature

Average tensile fracture stress and hydrogen/oxygen ratios are listed in table 1, for six

flame-polishing experiments. The as-grown

strength of the material varied between 2.2 to

2.80 X 10« N/m^ (3.2 to 4.0 x 10^ psi). With

the exception of series I and V, a maximum in

tensile strength was measured with increasing

flame temperature. This maximum strength

varied but always represented an increase of

249
Table 1. Average tensile fracture stress and Kt/Ot flame ratios for flame polishing experiments

Sample N"

Average

fracture stress X 10", N/m'H=/0.

Ratio Sample N°

Average

fracture stress X 10 N/m ' Ratio

I -1 2.19 2.5 IV-1 2.64 No flame

2 2.76 2.3 2 9 7Q 9 91 3 2.90 2.1 3 3.10 2.14

4 2.95 2.04 4 3.27 2.12

5 3.05 2.0 5 3.34 2.09

6 3.35 1.95 6 3.64 2.04

7 2.24 No flame 7 3.25 1.99

II -1 2.73 1.85 V -1 2.64 No flame

2 1.94 1.78 2 2.46 2.4 3 2.64 1.82 3 2.80 2.19

4 2.91 1.90 4 3.08 2.1 5 3.22 1.94 5 3.19 1.99

6 2.98 2.08 6 3.41 1.91

7 2.64 2.12

8 2.38 No flame VI-1 2.54 No flame

2 2.14 2.4 III-l 2.80 No flame 3 2.29 2.2 2 3.14 2.17 4 2.61 2.06

3 3.43 2.10 5 2.88 1.94

4 3.80 2.01 6 3.20 1.85

5 3.65 1.95 7 3.65 1.75

8 3.28 1.68

0.9 to 1.2 X 10" N/m- (1.3 to 1.7 x 10= psi)

over the strength exhibited by the standard fil- ament. A decrease in strength was observed for

filament treated in cool flames during runs I, V, and VI. This is thought to be a general oc- currence and absent from runs II, III and IV

because the filament vi^as not treated in suitably

cool flames.

The data given in table 1 for each flame-polishing run all follow the same trend and the

variation in tensile strength with flame temperature may be represented by the curve

drawn in figure 2. During each run, the filament-flame area was observed using a steromicroscopic and variations in the state of the

filament as it passed through the flame could

be observed. These observations have been used

to divide the curve in figure 2 into four sections as shown.

For flame temperatures in range A, little

change was observed in the filament as it passed

through the flame. The paraffin or starch coating was easily removed and none was retained

on the filament surface.

Ts the flame temperature was increased,

small localized areas of intense light, perhaps

Flame Temperature

Figure 2. Idealized plot of tensile fracture stress versus flame

temperature.

250
refracted from microscopic imperfections, could

be observed to disappear within the flame. This

was not a general phenomenon and occurred

infrequently, even in those lengths in which it was noted. An increase in tensile fracture stress over that exhibited by the standard sample was

always measured. This was often as much as

5.5 X 10« N/m^ (8 x 10* psi).

At slightly higher flame temperatures (range

C), a flickering white intense zone was observed. This marked the beginning of surface

melting of the sapphire filament. When first

observed, the zone is unstable and often fades

away, reappearing a few seconds later. Also,

it tends to move vertically up and down the filament within the flame. As the flame temperature is further increased, the overall length of

the zone increases and it becomes essentially

stable, remaining fixed in the flame with respect

to the moving filament. A photograph of a stable zone is shown in figure 3. Some distortion

due to thermal currents is present and makes

photographic recording difficult. Optimum polishing conditions and maximum strength en- hancements are obtained when the zone is 6.2

to 8.8 X 10"^ m long, i.e., roughly two to three

times the filament diameter.

At higher flame temperatures, the filament

begins to neck-down and, if this is severe, it

assumes an hourglass appearance. A fall in tensile fracture stress from the optimum value is

measured. Initially at least, this is not due to a weakening in the filament through the introduction of defects but rather to local reductions

in cross-sectional area which are not measured

with a micrometer in diameter determinations

after testing to fracture.

The above observations of the filament sur- face conditions are made with respect to the

hottest zone in the flame, of which there are

two. The flame width at the filament is about

3 X 10"^ m long and the two hot zones are observed at the periphery of the flame. This inFiGURE 3. Optimum flame polishing conditions

showing the stable molten sheath at the

filament surface.

dicates that the flame is oxygen deficient and

maximum temperatures are reached where a supply of ambient oxygen is available. The hottest zone is always at the top, where the filament leaves the flame, and the two zones tend

to extend inwards as the hydrogen/oxygen ratio

decreases.

3.2. Condition of Flame Polished

Sapphire Filament

a. Depth of Molten Zone During Flame Polishing

Visual observation while flame polishing has

indicated that, under optimum polishing conditions, the molten zone has a length in the range

6.2 to 8.8 X 10"* m. Optical observations were

carried out in order to obtain a measure of the

radial depth of this molten zone. PhotomicroFiGURE 4. Photomicrographs of as-grown and optimum polished filament. Mid-plane section, (a) as- grown (b) optimum flame polished

251
graphs are displayed in figure 4 and are typical

of those obtained for standard and optimum

polished filament. The microvoids present,

formed when the material is grown, are the result of the diiference in density between molten

and solid sapphire. While the actual void distribution varies somewhat between the lengths

treated, the characteristic void-free outer shell

is always present. The thickness of this shell

varies from less than 6.5 x 10"^ m to as much

as 2.5 X 10"= m. If the molten zone during flame

polishing was deeper than these thicknesses, a

detectable reorganization of the void layers

t

*

t «

1 ' «

f ' *

f

1 1 *

tu. ................ -- —-J

Figure 5. Photomicrograph of slightly overpolished

filament showing that the necked-down region has

not penetrated the void layers. Mid-plane section.

effected might be expected. Detailed examination of the void distribution before and after

flame polishing reveals no evidence that any

changes in void distribution have occurred. Further evidence for the depth of the zone being

less than about 6.5 x 10 ® m is obtained when

filament which has been marginally overpolished (range D in fig. 2) is examined. A photomicrograph is shown in figure 5. Evidently, the

radial temperature distribution around the filament is asymmetric and excessive surface melting and loss of geometric integrity has occurred

for one side. However, the start of the neckingdown has not reduced the cross-sectional area

by more than 2 X 10-^ m and the molten zone

has not penetrated the void layers. When severe necking-down in a hotter flame is obtained and

the filament assumes an hourglass shape, the

void distribution is effected and resolidification

results in additional voids which are considerably larger. Examples of these microvoids are observed in figure 6. Defects of this type would

be expected to weaken the filament.

b. Surface Roughness

The average and high roughness values obtained by surface profilimeter analysis are plotted versus tensile fracture stress in figure 7. It is immediately apparent that despite some

scatter, particularly at the low strength end, a marked trend associating high strength with

small values of surface roughness is obtained.

The average surface roughness for a sample

exhibiting tensile strength close to 3.8 X 10^

N/M^ was 2.5 X 10"* m as compared with

6.25 X 10-8 the standard material with

a tensile strength of 2.55 x 10« N/m^

Figure 6. Photomicrograph showing overpolished filament and

macroscopic voids associated with solidification occurring both

from the center outwards and the surface inwards. Mid-plane

section.

252
Figure 7. Surface roughness versus tensile fracture stress for flame polished filament from Series VI.

SEM photomicrographs showing the filament

surface at three stages (as-grown, underpolished, and optimum polished) are displayed in

figure 8 through 10. Samples in the as-grown

condition showed evidence of surface waviness

as shown in figure 8. This rippling which is at

0.762 radians (43°) to the growth plane of the

filament could be interpreted as microfacetting

on the {1013} rhomohedral planes. The central

area where the striations are not obtained represents a change in curvature of the filament

toward a larger radius and indicates that the

material is not perfectly circular.

Figure 8. SEM of as-grown sapphire filament surface.

Figure 9 shows a micrograph of an underpolished surface. The relief shown is not found

over the entire surface to the extent of the

waviness exhibited by the as-grown material

but is typical of surface roughness observed in

underpolished samples. The nonuniformity of

the surface is considered to be the result of the

flickering nature of the molten zone during

underpolishing. The surface of the optimum

polished material is shown in figure 10. At

magnifications up to 5000X, this surface is es- sentially smooth and shows no signs of rippling.

Figure 9. SEM of underpolished sapphire filament

surface.

253
Figure 10. SEM of perfectly flame polished sapphire

filament surface.

c. Residual Stress Petrographic examinations revealed no evidence of radial stress patterns under polarized

light. It is appreciated that the Youngs' modulus exhibited by sapphire makes the detection of

elastic strain difficult.

Samples were examined for evidence of bending after polishing in a manner expected to produce a nonsymmetric stress distribution if a

radial stress pattern were present.

The curvature to be expected for a given

surface stress may be calculated as outlined

below [12].

Curvature in a uniform

member

M

E •! (1)

where M is the bending moment, E is Youngs'

modulus and I is the moment of intertia of the

member.

Also:

M

O-f =

I (2)

where or is the longitudinal stress, y is the distance of the neutral plane from the plane of

maximum stress and M and I are as above.

Equations (1) and (2) may be combined and

lead to:

Filament Curvature = - E V(3)

Consider a stress of 7 x 10« N/M^ (10^ psi)

and assume that the neutral plane is midway in

the polished section (i.e., parallel to and 6.25

X lO-'^ m from the polished surface). Thus,

Filament

Curvature =

7 X m

4.65 X 10" X

24.1 m-i

6.25 X 10-

The radius of curvature, p, is given by the re- ciprocal of the filament curvature. Therefore,

1 = - 4.2 X 10-2 m

24.1

This represents a considerable degree of

bending. Although the calculation is not rigorous, it may be considered to given an approximate measure of the curvatures expected and

has been verified by measurements on highly

stressed ribbons [13].

Examination of filament in the as-grown and

polished conditions prepared as outlined above,

revealed no noticeable curvature and suggest,

on the basis of the above calculation, that the

components of any tempered stress distribution are much less than 7 x 10^" N/m^ (10^ psi).

4. Discussion

Proctor, et al., [6, 7] obtained strength

values in pure bending in the range 4.8 to

6.9 x 10» N/m- (7.0 - 10.0 x 10= psi), which

were the highest reported for sapphire in a form other than small (<10-'^ m diameter)

whiskers. Similarly, Noone, et al., [8] obtained

values in four-point bending close to 5.5 x 10^

N/m" (8 X 10= psi). The main disadvantage in

comparing these values with those of the present work is that bend tests on cylindrical samples only place a small fraction of the sample

surface (a line in the case of four point or pure

bending) under maximum stress, unlike tensile

testing where the entire gauge volume is

stressed uniformly. However, Noone, et al., [8]

repeated their measurements in tension testing

using similarly flame polished 5.0 X 10-* m diameter AI2O3 rods with an average gauge length

of 1.53 X 10-2 They obtained an average tensile strength value (14 samples) of 2.6 x 10^

N/m^ (3.77 X 10= psi). The ratio of strength

in bending to strength in tension (2.1) for the

data reported by Noone, et al., [8] is very close

to the value (2.12) for similar measurements

in cast iron [14] where the extent of plastic

deformation prior to fracture was very small.

This factor when applied to the results of Proctor, et al., [6, 7] produces strength in tension

values of 2.3 X lO** N/m^ (3.25 X 10= psi) for

samples with a test length of 6 X 10-=* m. Careful selection of more perfect smaller gauge

length samples (3 X 10-^ m long) was required

to produce bend strength values in the region

of 6.9 X 10" N/m^ (10*= psi). The correction

factor (2.1) reduces this value to near 3.3 x

10" N/m- (4.75 X 10= psi) for tension testing.

In comparison, the maximum tensile fracture

stress values obtained during the present investigation, in excess of 3.5 x 10" N/m^ (5 x 10=

psi) using 2.5 X 10-^ m gauge lengths, can be

considered encouraging. It should also be noted

254
that the material used in the present work was

not subjected to any selection process at any

point during the investigation.

The optimum flame-polishing conditions produced surfaces which contained none of the

helical rippling reported by the two previous

investigations [6, 7, 8]. This is the result not

only of the nonrotation of the filament during

treatment, but also indicative of less severe surface melting taking place. Noone, et al., re- port that the 5 x 10"^ m long molten collar produced at the rod surface was about 2.5 to 5.0 x

10"^ m deep, suggesting a quite radical change

in state of a considerable volume fraction of a

5 X 10-* m diameter rod (approximately 20%).

By contrast, the optical microscopy evidence

presented here indicates that the molten zone

is less than 10"^ m, thick which means that

about 10 vol percent of the filament is melted

during optimum flame polishing. Thus, viscosity

and surface tension enable the geometric integrity of the filament to be maintained and improved, as evidenced by the SEM photomicrographs.

Flame polishing and any other thermal treatment may be expected to change the surface

and/or internal characteristics of sapphire filament. As suggested earlier, some consideration

has been given to the thesis that a major role

in strength enhancement arises through either

strain relieving or in the creation of a stressed

state akin to that present in tempered glass.

The results of the petrography examinations

and curvature experiments suggest that the role

played by strain is small compared with the

changes in strength measured. Therefore, it may be reasonable to consider the strength

versus flame polishing characteristics of the

filament, as summarized in figure 2, in terms of

fracture initiating microcracks present in the

material.

Griffith analysis gives the fracture stress in

the presence of a crack as: _ h

'

where o-t is the fracture stress, y is the surface

energy, 2c is the internal crack length (c is the

surface crack depth), p is the crack tip radius,

and a is the interatomic spacing. Evidently the

stress concentrating effect of the crack is greatly dependent on the value of c/p. Thus, the

sharper the crack profile the greater its stress

concentration effect and the lower the fracture

stress observed.

As the filament passes through cool flames,

range A in figure 2, these microcracks may

widen, due to thermal expansion, causing the

depth or length to increase in an irreversible

fashion. A small decrease in tensile fracture

stress is measured as the result of the increase

in crack length.

During optimum and near optimum flame

polishing conditions, a substantial increase in

strength (1.1 x 10^^ N/m^ (1.6 x 10^' psi)) over that exhibited by the as-grown material is

observed. The structural evidence presented

here is conclusive that this maximum enhancement in strength is coincident with the attainment of a more perfect surface. Indeed, the

profilimeter data in figure 7 suggests a simple

correlation between decreasing surface roughness and increasing tensile fracture stress.

However, a considerable strength enhancement

of about 5.5 x 10« N/m^ (8.0 x 10* psi) is

observed for treatment in flames which are too

cool for surface melting to occur (range B).

This may be attributed to thermally activated

diffusion leading to crack blunting, since an improvement in the surface quality has been

noted. Any such process will result in a reduction in surface energy and it is not improbable

that they will occur at elevated temperatures.

Heuer and Roberts [51 and Kirchner et al., [151

reported that strength increases were measured

on annealing single crystal and polycrystal

aluminia, respectively. Lange and Radford [16]

showed that the strength of polycrystalline

AlsOr,, weakened by the introduction of cracks

through thermal quenching, was rapidly re- covered on annealing at 1700 °C. While in these

relatively weak materials the surface of the

samples was of primary importance, it is not

unreasonable to suspect that similar mechanisms will be operating within the bulk. Thus,

it is highly possible that the thermal environment during flame polishing contributes to the

strength enhancement by means other than the

removal of surface flaws, through remelting.

These considerations suggest that the creation of the molten high temperature surface

zone, which does not destroy the geometric integrity of the filament, may also be the creation

of an efficient thermal environment for diffusion controlled blunting of internal microcracks.

This hypothesis could be confirmed by concrete

evidence that fracture origin is internal, before

and after flame polishing. Some evidence is

available to support this theory. Filament tensile samples fracturing at high loads [i.e.,

greater than 2.1 x 10" N/m^ (3 x 10^ psi)l

generally shatter, at least in the immediate

break vicinity. LaBelle and Hurley [17] reported that preliminary examination of the fracture

surfaces pointed to internal fracture origin at

high fracture loads in tension. This work has

been extended, using as-grown samples embedded in epoxy to reduce shattering and loss of

the primary fracture surfaces. SEM examination of six fracture surfaces points to internal

nucleation in all cases. Details of the investigation will be reported at a later date [18] . Thus,

in the sapphire filament used in the present

255
work fas-grown strength > 2.1 X 10^ N/m^

(3 X 10^ psi)], there is evidence to support the

supposition that fracture origin is internal,

perhaps associated with the microscopic spherical voids. While there is conclusive proof that

optimum flame polishing results in the creation

of a more perfect surface, there is also reason

to believe that thermal healing of internal microcracks may be the primary cause for the

enhancement in tensile fracture stress observed.

Over-polishing results in a fall in strength,

due at first to the reduced cross-sectional area,

and then, as the flame temperature increases,

to the formation of resolidification voids. These

voids are indicative of solidification proceeding,

both from the central core outwards and from

the resolidified shell inwards. Unlike the lO""

m spherical voids present in the as-grown material, these voids have non-symmetric shapes

and higher stress concentrating factors.

5. Summary

Continuous flame polishing of nominal 2.5

X 10"^ m diameter sapphire filament enables

strength enhancement of about 1.1 x 10® N/m^

(1.6 X 10^ psi) over values exhibited by as- grown material. Maximum tensile strengths obtained, 3.8 X 10® N/m^ (5.5 x 10^ psi), are

the highest reported for sapphire in a form

other than small diameter whiskers. These

strengths are obtained after polishing in oxygen-hydrogen flames which produce a stable

molten zone 8 x 10"^ m long by — 10"^ m deep

at the filament surface. The geometric integrity

of the material is completely maintained and

surface examination by SEM shows that a

highly smooth surface is created. Nevertheless,

evidence concerning the internal nature of fracture in the as-grown material suggests that

thermal healing of microcracks may be playing

a more important role in the strength enhancement process than the creation of a more perfect surface.

and F. H. Cocks for valuable discussions of

many aspects of the work.

This work was supported by the Air Force

Materials Laboratory, Nonmetallic Materials

Division, Wright Patterson Air Force Base,

Ohio, under Contract No. F33615-69-C-1369.

The author wishes to thank Mr. John Bailey

who carried out much of the flame polishing

and mechanical testing, and Drs. G. F. Hurley

[1

[2

[3

[4

[5

[6

[7

[8

[9

[10

[11

[12

[13

[14

[15

[16

[17

[18

6. References

Proctor, B. A., Fracture of glass, Appl. Mat. Research 3, 28-34 (1964).

Morley, J. G., Andrews, P. A., and Whitney, I., Strength of fused silica, Phys. Chem. Glasses

5, 1-10 (1964).

Brenner, S. S., Mechanical behavior of sapphire

whiskers at high temperatures, J. Appl. Phys.

33, 33-39 (1962).

Wachtman, J. B., and Maxwell, L. H., Plastic

deformation of ceramic oxide single crystals,

J. Amer. Cer Soc. 37, 291-299 (1954).

Heuer, A. H., and Roberts, J. P., The influence of

amealing on the strength of corundum crystals, Proc. Brit. Ceram. Soc, 6, 17-27 (1966).

Morley, J. G., and Proctor, B. A., Strength of

saoDhire crystals, Nature 196, 1082 (1962).

Mallinder, F. C. and Proctor, B. A., Strengths of

flame polished sapphire crystals, Phil. Mag. 13,

197-208 (1966).

Noone, M. J., and Sutton, W. H., Investigation of

bonding in oxide-fiber (whisker) reinforced

metals, National Technical Information Service

(NTIS) AD 664-176.

Popov, D., The growth and uses of gem grade

corundum crystals, Book, Growth of Crystals,

Ed. A. V. Shubnikov and N. N. Sheftal, Volume

II, (Chapman and Hall, London, 1962).

LaBelle, H. E., and Mlavsky, A. I., Growth of

sapphire filaments from the melt, Nature 216,

574-575 (1967).

Kingery, W. D., Introduction to Ceramics, 639-

641 (John Wiley & Sons, Inc., New York, 1960).

Crandall, S. H., and Dahl, N. C., An introduction

to the mechanics of solids, 288-294 (McGrawHill Book Co., Inc., New York 1959).

Hurley, G. F., to be published.

Rozner, A. G., Comparison of tensile strength

measured in tension and bending. Trans.

A. I. M. E. 233, 1646-1647 (1965).

Kirchner, H. P., Gruver, R. M., Platts, D. R., Richel, P. A., and Walker, R. E., Chemical

Strengthening of Ceramic Materials, Summary

Report, Contract NOOO19-68-C-0142, Naval

Air Systems (1969).

Lange, F. F., and Radford, K. C, Healing of sur- face cracks in polycrystaline AI2O3, J. Amer.

Cer. Soc. 53, 420-421 (1970).

LaBelle, H., and Hurley, G. F., Strength and

fracture of long single crystal sapphire fila- ments, Sampe Journal, 6 17-23 (1970).

Pollock, J. T. A., and Hurley, G. F., to be published.

Discussion

Matsko: Did the unidirectional heating of

the filament present any problem?

Pollock: Yes, in the paper there is some

evidence presented which indicates that there

is some asymmetry of the actual temperature

distribution.

Noone: Didn't you find that some of the

strain that you measured was due to the fact

that you were melting more AI2O3 on one side

of your filament therefore producing this complex stress state ? You are adding another variable by melting more on one side than the other.

Pollock: As I said, there is some asym256
metry I don't think it's particularly important. GlELlSSE: Did you mention the cross section

No evidence of strain in the crystals could be size of the filament? .r,-, • ^:^^^fp^

detected, before or after treatment. Pollock: I didn't. It was 10 mils in diameter.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

257

Preparation of Smooth Crystalline, Damage Free,

Sapphire Surfaces by Gaseous Etching

W. A. Schmidt and John E. Davey

Naval Research Laboratory, Washington, D.C. 20390

Sapphire has a combination of properties such as, relative chemical inertness,

exceedingly high resistivity, optical transparency and crystalline structure that suggest

its use as a substrate material for various semiconductor devices. Its successful use

for silicon vapor phase epitaxy substrates indicates that similar results might be

possible with other semiconductors. Previous epitaxial studies have shown that surface

crystalline disorder and topographic imperfections inhibit epitaxial growth; therefore

highly polished well ordered sapphire surfaces were needed. Mechanical polishing

alone is insufficient as it results in cold flow and work damage.

Initial experiments in removing surface damage were performed with hot

phosphoric acid and with hydroxide etches. Undesirable preferential etching was observed for both processes and the processes did not seem amenable to the routine

production of high quality surfaces. Two other surface treatments for which successful

results have recently been reported were investigated.

The first was a hydrogen firing technique in which the substrates were heated in

hydrogen in a Mo wound resistance furnace at 1500 °C. The second was a simple

laboratory process in which the sample is heated by means of RF heating of a carbon

susceptor to 1350 °C in an atmosphere of helium and fluorotrichloromethane (Freon11). Commercial reagents were used throughout and a fused quartz tube was used for

the reaction chamber for the second technique.

Reflection electron diflfraction (RED) was used to determine the surface crystalline

order, and electron microscopy (EM), using high resolution replication techniques, was

used to examine the topography structure of the surfaces. A nimiber of different

substrates from various industrial sources, with different surface topographies and

different orientations were used.

Hydrogen firing at 1500 °C results in an etch rate of 0.1 10" m/min. Firing for

times up to 45 min did not produce consistent surfaces on the 0°, 60° or 90° orientations. While hydrogen firing did produce high quality well ordered crystallograph

surfaces by RED, their topographic condition was poor. Resolvable surface structures

could be detected on some with standard optical microscopy (200x, dark field) and for

all high resolution EM measurements. The final surface finish quality was related to the quality of the prefired surface, indicating that complete damage removal was not

accomplished.

Freon firing at 1350 °C resulted in an etch rate of 1.5.10"' m/min. This technique

consistently produced well ordered, high quality, surfaces for the 60° and 90° oriented

surfaces. EM viewed 60° and 90° freon fired surfaces had no resolvable surface

structure after a 5 min etch and were well ordered crystallographically as measured

by RED. The same etch produced strongly etched surfaces on 0° oriented material.

Key words: Crystalline sapphire surfaces; fluorinated-hydrocarbon etch; fluorotrichloromethane; freon etching; hydrogen annealing; sapphire; sapphire substrates; sapphire surface preparation.

1. Introduction

Sapphire has a combination of properties

which make it a desirable material for use as a

semiconductor substrate. [1-2]^ It is resistant

to attack by most of the materials used in

present semiconductor processing and etching

technologies. It is crystalline, transparent,

insulating and it has a relatively high thermal

conductivity. Table 1 is the table presented by

Mueller detailing the particular advantages for

its use v^^ith silicon technology. [1] Since the

transparent M^indow of a sapphire substrate

(typically 0.5 mm thick) extends from the near

UV to about 62,000 A, optical methods can be

used for examining the band gap, and many

other electronic transitions v^^ithin the lattice

' Figures in brackets indicate the literature references at the end of this paper.

Table 1. Advantages of the silicon-on-sapphire

technique

Electrical advantages

1. Excellent isolation (p > 10" ohm-cm).

2. Good high frequency performance (loss tan <10"'

at 800 MHz).

3. Extremely low capacitance devices and circuits.

4. No possibility of "4-layered latch-up".

5. No allowance necessary for space-charge spread.

6. Considerably greater resistance to transient radiation. Mechanical advantages

7. Ability to withstand regular Si processing, temperature of diffusion, chemical etches, etc.

8. High strength, especially at high temperatures.

9. Good heat conductivity (50x glass).

10. Smoothness, no voids.

11. Very high packings density (10° complementary

devices per sq. in).

12. For many circuits, fewer processing steps are re- quired.

259
of the semiconductor film, by irradiating

through the substrate.

The successful use of sapphire as a substrate

for Si, Ge and GaAs [1-7] vapor phase epitaxy

by various authors indicates that it may be a

useful substrate for vacuum deposition of these

and other semiconducting materials.

In previous work by the authors on the va- cuum deposition of GaAs on GaAs and GaAs

on quartz, sapphire, and GaP [8-9] it has been

shown that surface crystalline disorder and

topographic imperfections inhibit epitaxial

growth; highly polished well ordered sapphire

surfaces were needed.

Mechanically polished sapphire is an improvement over a completely amorphous material such as fused quartz, however, it is not adequate for use as an epitaxial substrate. The

surface of a mechanically polished sapphire

sample is essentially amorphous due to work

damage and cold flow.

In an attempt to restore the mechanically

polished surface to the same crystalline order

as the bulk material, various chemical etchants

were employed. The initial experiments were

performed with hot phosphoric acid and with

hydroxide etches. Undesirable preferential etching was observed with both processes, and their

complexity was such that the incorporation of

these etches as routine steps in normal wafer

processing was difficult.

Two articles by Manasevit [10-11] on the

gaseous etching of sapphire with sulfur

fluorides and fluorinated hydrocarbons and Allison's [2] account of hydrogen etching indicated that the desired surface quality was obtainable. To evaluate these processes the

Manasevit procedure was used with fluorotrichloromethane CFCI3 as the freon etchant and

samples were hydrogen fired at 1500 °C as described in the Allison article.

2. Experimental Methods

The two forms of gaseous etching investigated, required similar apparatus. The sample

was held on an inert support in a high temperature gas flow of the etchant gas.

The high temperature hydrogen etching was

performed in a commercially available resistance heated furnace. The hydrogen furnace was

tubular in design with a refractory tube which

contained the hydrogen atmosphere and the

molybdenum resistance heating elements. The

sample carrier was coated with alumina to prevent reaction with the samples. Fifteen minute

preheat and cool down periods in the cooler

end sections were required to prevent thermal

shock to the furnace walls and to the sample

holder.

Samples were hydrogen etched at two temperatures in this furnace at 1500 °C the temperature used by Allison, [3] and at the maximum temperature of the furnace, 1700 °C.

The He freon apparatus is shown in figure 1. A quartz reaction tube was used to contain the

etchant and carrier gasses around the heated

samples. Samples were heated to 1350° with

R.F. using a spectographic grade ultra pure

carbon rod as the sample holder and R.F.

susceptor. The hot zone of the carbon rod was

held separate from the quartz tube walls by

quartz spacer rings and rods, which were attached to the susceptor ends outside of the hot

zone. A prefired alumina plate was used to support the samples to eliminate contamination

from the carbon. The cold trap on the down

stream side of the tube was used to trap reaction products which clogged the vent jet in

early runs. Temperature measurements were

made with an optical pyrometer looking indirectly through a front surface mirror onto

either the edge of the sample or the alumina

plate under it. No temperature correction was

made for the presence of either the mirror or the quartz tube in the optical path. The design

of the gas feed apparatus was dictated by the

choice of the type of freon. Freon 11, fluorotrichloromethane, is a liquid at room temperature

with a vapor pressure approaching 1 atmosphere. Bubbling a low flow of He gas through

a three inch column of freon 11 was sufficient

to transport freon vapors into the low pressure

system. Initial experiments indicated that passing the entire He flow through the bubbler produced too rich a mixture for reasonable control

of etch rate. The total flow rate of the He is

determined by the need to maintain a low pres- sure of He in the main tube with a vent jet size

not likely to clog with reaction products, A low

He flow was bubbled through the freon 11 and

the mixture was returned to the feed line at the

outlet of the mainline flowmeter. The mass flow

of the freon was measured by weighing the

freon container before and after a few test runs

to establish the mass flow rate of freon associated with a specific flow of He in the bubbler

line.

The flow conditions used were:

Helium 410 ml/min

Helium in freon bubbler 70 ml/min

Freon 0.135 gm/min

All of the samples were examined for their

surface quality before and after each etching

treatment. The determination of surface quality was made with three types of equipment.

Optical microscopes were used to cover the

range from Ix to 500x magnification with both

brightfield and darkfield illumination. An optical interference microscope was used to measure surface planarity. A JEM-6 electron microscope was used to examine surfaces under high

magnification (10*x). Standard carbon film

260
i«.UMMA PLATE ^ ^ COIL sapphire

FRONT SURFACE — .-\ MIRROR

CAR30N SUSCEPTOR OOOOOOOOODOO

COLD TRAP

HEL lUM

FREON II

Figure 1. He-freon etching apparatus.

techniques were used to make the surface replicas. Reflection electron diffraction and x-ray

diffraction analysis were used to determine the

crystalline order of the surfaces.

The sapphire used was obtained from two

sources. Sapphire windows and Czracholski

pulled boules of single crystal sapphire were

obtained from the Union Carbide Co. The

boules were cut and polished in the NRL optical

shop. "Substrate" grade blanks similar to those

used commercially for silicon vapor phase epitaxy were obtained from the Adolph Meller Co.

The surfaces of the samples were quite different, but consistent within a group, (fig. 2) Note

distance between marks is 1 /^m. Sapphire

which was cut and polished at NRL (fig. 2A)

using a 1/4 diamond final polish had extensive scratches detectable with even a 200x optical microscope. Substrate grade material (fig.

2B) was also heavily scratched. Generally the

window grade material had a lower scratch

density than the mechanically polished or the

substrate grade material. Only a few scratches

were seen with the electron microscope on the

60° and 90° (fig. 2D) "window" grade surfaces,

while the 0° surfaces were not as high quality

(fig. 2C). The crystalline order of the "substrate" grade samples, however was far

superior to the other grades of samples (fig. 3).

The "substrate" grade samples were ordered as observed by the RED pattern. The other surfaces were completely amorphous. All of the

samples used, from all sources, displayed single

crystal patterns under x-ray analysis.

3. Experimental Results

3.1. Hydrogen Etching

Hydrogen etching of 60° and 90° samples

produced inconsistent surfaces. All of the

samples etched in hydrogen had excellent

crystallographic quality (fig. 4). The surface

topography of these samples (fig. 5) was not

consistent. Some samples had scratch free sur- faces as examined under the electron microscope, while other retained residual scratches

(fig. 5B) and still others had extensive striations (fig. 5A). Entire surfaces were densely

covered with striations over 1 /nm long.

After 45 minutes of hydrogen etching at

1500 °C residual scratches and striations were

found in some of the samples. With the etch

rate of 10"^ meters/minute, the same as Allison, [4] this corresponds to a removal of

4.5.10"® meters, and while the samples had

numerous scratches prior to etch none were

this deep. It was assumed that the scratches

remaining after hydrogen etching were the

result of work damage. The origin of the striations is at this point unclear. They were found

on the etched "substrate" grade material more

frequently than on the other grades. With hy439-921 O - 72 - 18 261
(A) 1/4 MICRON DIAMOND FINAL POLISH (B) SUBSTRATE GRADE SURFACE FINISH

(C) WINDOW GRADE SURFACE FINISH 0° CUT (D) WINDOW GRADE SURFACE FINISH 60° CUT

Figure 2. Sample surfaces before gaseous etching by electron microscope examination.

(A) WINDOW GRADE

(B) SUBSTRATE GRADE

Figure 3. Reflection electron diffraction of

sample surfaces before gaseous etcthing.

drogen etching as with other methods of sur- face preparation the basal plane or 0° surfaces

were preferentially etched.

3.2. Freon Etching

With the exception of the 0° surfaces all of

the Helium-freon etched surfaces were scratch

free and single crystal, (fig. 6). In this figure

surfaces etched the standard 5 min as well as a surface etched for an extended time (15

min) are shown. No residual scratches or discrete structures were visible at this magnification other than artifacts of the replication

process. The etch rate was typically 7.5.10'®

m/min. In a typical 5 min etch run 37.5 10'^

meters of material was removed from a surface, effectively removing all work damaged

material. This etch rate was varied by changing

the level of the freon content in the gas. Higher

etch rates resulted in heavy etching of the edge

of the samples, however, and were not used for

that reason. The 0° surface for both He and

He-freon were preferentially etched (fig. 7)

with the resulting etch structure from freon

(A) resembling those found by Manasevit[2]

who used chlorotrifluoromethane as the etchant

gas. All of our electron microscope data was

taken at magnifications about one order of magnitude greater than those of Manasevit, however, there was no indication of the extensive

262
(A) WINDOW GRADE (B) SUBSTRATE GRADE

Figure 4. Reflection electron diffraction of hydrogen etched surfaces.

(A) STRIATIONS (B) RESIDUAL SCRATCHES

Figure 5. Typical surfaces after hydrogen etching.

etch pit formation seen by him except on the

basal plane where our etch structures were still

about one order of magnitude smaller than his.

It is assumed that this difference is a result of

etchant chemistry, which is one factor obviously

different in the two experiments. The hydrogen

annealed 0° surface is typical of other workers

results. A common complaint with the use of

fluorinated hydro carbon etchants is that the

surfaces that result are non planar or have

dimpled surfaces. This effect was not observed

by the authors.

4. Conclusions

The use of freons as vapor high temperature

etchants for sapphire has been established

previously, but here the use of one, fluorotrichloromethane, has been shown to be useful in a simple apparatus for producing high quality

scratch free ordered, single crystal surfaces.

The process has been shown to be more reliable

than the hydrogen anneal used by industry,

which produces questionable surfaces, containing dense striations and residual scratches.

Table 2 summarizes the results found.

Table 2. Summary of results

Hydrogen etching

1500°— 1700° C 15-45 min.

Each Rate 10"' meter/min.

Results in ordered crystalline surfaces

extensive striations

residual scratches

0° surface, preferentially

etched

He-Freon etching

1350° C 5 min.

Etch Rate 7.5.10-" meter/min.

Results Ordered-crystalline high Quality

Scratch Free Surfaces

0° Surface, Preferentially

Etched

Further work is being pursued to incorporate

this process into a standard vacuum evaporation

process. The corrosive reaction products prevent direct inclusion, and a further insitu cleaning process (thermal bake. Ha Fire low temp)

must be used to compensate for the contamination during transfer through atmospheric

ambient.

263
(A) 5 MINUTE STANDARD ETCH

60° AND 90° SURFACES

(B) 15 MINUTE EXTENDED TIME ETCH

60° SURFACE

(C) REFLECTION ELECTRON DIFFRACTION ANALYSIS OF TYPICAL SURFACE

Figure 6. Helium-freon etched surfaces.

(A) BY HE FREON (B) BY HYDROGEN

Figure 7. Preferential etch of 0° cut material.

5. References

[1] Mueller, C. W., Thin-film devices on dielectric

substrates J. of Vacuum Science and Technology 7, 147 (1970).

[2] Allison, J. W., Dumin, D. J., Heiman, F. P., Mueller, C. W., and Robinson, P. H., Thin

film silicon: Preparation, properties and device app. Proc. IEEE, 57, 1490 (Sept. 1969).

[3] Manasevit, H. M., Thorsen, D.C., Heteroepitaxial

GaAs on aluminum oxide I. early growth

studies. Metallurgical Transactions, 1, 623 (Mar

70).

[4] Manasevit, H. M., Single crystal gallium arsenide on insulating substrates, Appl. Phys. Ltrs. 12,

156 (15 Feb. 1968).

[5] Manasevit, H. M., Simpson, W. I., Single-crystal

silicon on a sapphire substrate, JAP, 35, #4

1349 (1964).

[6] Manasevit, H. M., Nolder, R., Cadaff, J., et al.,

264
Heteroepitaxial silicon-aluminum oxide interface, Pare I Trans. Met. Soc. of AIME 233,

540 (Mar 65).

Part II, Trans. Met. Soc. AIME 233, 549

(Mar 65).

Part III, Trans. Met. Soc. AIME 242, 465

(Mar 68).

[7] Zanowich, R. J., Morritz, F. L., Heteroepitaxial

deposits of Ge and GaAs/Ge on sapphire and

beryllium, J. Electrochem. Soc. 114, #6 146C

(Jun 67).

[8] Schmidt, W. A., Davey, J. E., Zn Doping of

GaAs Films Using Coevaporation of the Elements, Naval Research Laboratory Report

#7100, (April 1970).

[9] Davey, J. E., and Pankey, T., Epitaxial GaAs

films Deposited by vacuum evaporation, J. Appl. Phys. 39, #4, 1941, (March 1968).

[10] Manasevit, H. M., Morritz, F. L., Gas phase

etching of sapphire with sulfar fluorides, J. Electrochem. Soc. 114, #2 (Feb. 1967).

[11] Manasevit, H. M., Gas phase etching of sapphire

III fluorinated hydrocarbons, J. Electrochem.

Soc. 115, #4 (April 1968).

Discussion

Heuer: Do you know how those substrate

grade samples were prepared ? Schmidt: No. The only answer I got in a query of that sort was that it was proprietary

information and they weren't willing to give it out. My feeling is that it's probably mechanically polished with quick hydrogen anneal.

Clearly the scratches would indicate most of

the work damaged material was removed. There

is high crystalline order. I would expect a quick

hydrogen anneal was performed after mechanical polishing.

Heuer: What were the hexagonal features

shown on as-received material ? Schmidt: The hexagonal patterns are found

on 0° material and are etch pit structures. Both

the as-delivered and etched surfaces had these

patterns.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).

265

The Strength of Gas Polished Sapphire and Rutile

R. W. Rice, P. F. Becher, and W. A. Schmidt

U.S. Naval Research Laboratory, Washington, D.C. 20390

Surface polishing of a-AljOs using a helium-freon gas mixture at elevated tenn- peratures is shown to yield bend strengths comparable to flame polished sapphire.

However, the gas polishing process proved to be more versatile as evidenced by the

success with flat bar surfaces. Further, the process can be readily applied to some

other materials as demonstrated by the substantial improvements in the strength of

Ti02 after gas polishing. Limited attempts at polishing other ceramic materials are

also discussed, as well as observations on strength variations and fracture behavior.

Key words: Chemical polishing; rutile; sapphire; spinel; strength; twinning.

1 . Introduction

Further development by Schmidt and Davey

[1] ^ of a chemical polishing technique for sapphire utilizing hot, flowing, helium-freon gas

mixtures showed that very good surfaces were

obtained relatively easily. The high quality of

the surfaces produced suggested that high

strengths could be achieved by such polishing. Further, it was felt that this polishing

technique might also be applicable to some other

ceramic materials. This paper presents results

of the initial investigation of the strengths of

gas polishing ceramics, especially oxides (e.g.,

AI2O3, and TiOa crystals). Some observations

on fracture are also made.

2. Experimental Technique

The materials used have been previously described [2-4], as has the data on orientations

and test techniques. During the gas polishing

process, the test bars were placed in graphite

fixtures which supported the ends of the

samples and prevented contact between sample

surfaces. After polishing, sample handling was

minimized and any handling and storage was

undertaken in a manner to prevent damage to

the surface. During the course of polishing and

mechanical testing (3 point flexure), it was

found that a strip of thin transparent pressuresensitive tape placed on the compressive sur- faces aided in handling and helped retain fracture sections of high strength specimens for

later examination.

3 Results and Discussion

3.1. a—AI2O3

Sapphire bars up to about 4 cm in length

could be polished in the present apparatus.

These longer bars were found to polish best on

surfaces nearest the gas exit end of the suceptor, indicating that some initial preheating of

'Figures in brackets indicate the literature references at the end of this paper.

the gas would be beneficial. In order to minimize any finishing gradients, most bars were

polished for 10 minutes in one orientation, then

polished again for the same period after being

turned 180°. Results show a trend for a more

uniform, but generally a less flat, surface re- sulted if bars were gas polished in the as- ground rather than a mechanically polished

condition. During the polishing schedule it was

observed that greater amounts of removal gave

somewhat smoother surfaces. The resultant

surfaces were very glossy and free of any

sharp surface discontinuities. Except for some

slight rounding of the edges, removal was very

uniform on all surfaces, indicating little efi'ect

of crystallographic orientation. Chemical etching failed to reveal any inhomogenity of sur- face texture and pit densities were generally

in the order of lOVcm-. An exception occurred

when only a slight amount of material was re- moved from the surface of the bars ; then etch

pit densities were generally much higher in re- sidual grinding or polishing stria.

Although not studied in detail, observations

on bars that were not extensively polished, at

least over part of their length, indicated that

removal of less than 0.1 mm. of the as-ground

or mechanically polished surface generally re- sulted in low strengths (of the order of those

achieved by conventional finishes). A more detailed characterization of similar sapphire

surfaces polished in the same apparatus is

given elsewhere [1] and is thus not further

discussed here.

Bend test results, listed in table 1, indicate

that quite high strengths can be obtained,

generally being more reliable than those obtained by flame polishing similar flat bars [4]

In subsequent examination, fracture origins

were located on about half of these specimens,

being detected somewhat more frequently on

lower strength bars. It was found that about

50-60 per cent of the fractures originated from

the edges, but no clear trend in strength or other parameters was noted. Fracture mirror

267
dimensions were observed to be generally consistent with other fracture data obtained in a broader study of ceramics (to be published

later). Initial optical examination indicated

that most samples do not contain twins. The

few samples observed to have twins were of

average or lower strengths.

Trial gas polishing of dense, high purity,

sintered AloO^ bodies having either fine or large

grain sizes produced very dull surface finishes.

The gas polishing was observed to preferentially remove material at and adjacent to the

grain boundaries ; however, no sharp local radii

of curvature at boundaries were detected.

Strengths of such polycrystalline specimens

were unaltered from those obtained by either

grinding (parallel to the tensile axis) or mechanical polishing.

3.2. TiOz

Both as-grown (blue) and air annealed

(yellow) bars were gas polished after which

each type exhibited a yellow color indicating

that reduction was not occurring as found in

flame polishing [4]. In addition, material was

removed from rutile single crystals at a much

higher rate than from ALO., subjected to the

same polishing conditions. Some of those TiO^

specimens developed extremely glossy surfaces,

while others developed a frosted appearance

over part or all of their surfaces, indicating

that polishing was very sensitive to gas flows

and temperature conditions (or possible con- tamination from other materials in the same

run). Again, as with the sapphire, removal appeared to be quite uniform from all surfaces.

As shown in Table 1, those bars having very

glossy surfaces had strengths from a few to

several-fold times those obtained with conventional mechanical finishes. On the other hand,

those with even small frosted patches on the

side of the bars had strengths comparable to

those obtained by conventional machining techniques. Fracture origins were located on all eight specimens, with four being from the tensile surface and four from the edges, yielding

no obvious trend with strength or surface condition. Fracture mirrors were usually observed

and were found to be consistent with data from

other tests. Optical observation of the highest

strength specimen showed several twins with

intersecting cracks and/or other twins along

one edge of the specimen near the fracture.

Subsequent chemical etching in fuming sulfuric acid ( 300-330 °C) revealed a dense mosaic

of intersecting twins and associated cracks adjacent to the fracture surface around the fracture origin. Since these appeared or became

visible very near the fracture origin, interaction of mechanical twins may be limiting

strengths. The twins originally observed along

the edge of this sample lay in a region of the

tensile surface subjected to much reduced stress

(70-100 X 10^ psi, 49-69 X 10^ N/m^) suggesting that twinning might be important in

lower strength rutile samples. This would be

enhanced by regions of stress concentration

such as the edges of the bar (as noted above)

and surface or internal defects.

3.3. Other Materials

A trial polish of a MgO crystal resulted in

a badly decomposed crystal which could be

crumbled into a white powder. Further work

with this material was therefore not attempted.

Unstabilized and CaO-stabilized ZrOo crystals which were in positions adjacent to the

MgO crystal during the same polishing run

were an opaque white through a substantial

portion of the bars. The resultant samples were

very weak and one was found to be deliquescent. Any success with gas polishing of ZrOa

was felt to require extensive change in the

technique which was not attempted.

ALOa-rich spinel crystals from a separate

gas polishing run had a thin white surface

coating. Sample strengths of these were lower

than those normally obtained by grinding or mechanical polishing. However, a spinel bar

obtained in another run (which included the

MgO and Zr02 samples) developed a fairly

glossy finish over part of the surface. The

Table 1. Room temperature bend strengths of sapphire and rutile single crystals

Material

Average strength

as machined

Average strength

as gas polished

Finish 10' psi 10'N/m= Range 10' psi lO'N/m Tests

Sapphire Ground 48±9 33 111-495 231±36 159 11

Mechanically

Polished 53±15 37 140-362 207±25 143 9

Rutile Ground 13±1 9 43-156 '78 ±30 54 4

Mechanically

Polished

22±7 15 15-20 '18±1 12 4

'Rutile crystals having glossy finishes after gas polishing.

Rutile crystals with frosted finishes after gas polishing.

268
strength obtained in this region of the bar was

comparable to those of mechanically polished

specimens. However, the strength achieved may

have been limited by precipitation of ALO^

during the cooling from the polishing temperature (—'1350 °C). Therefore, gas polishing of

spinel crystals, especially stoichiometric ones,

is considered to be promising with further

development.

Specimens of polycrystalline spinel were also

included in both of the previous gas polishing

runs. These developed mat or frosted surface

finishes, and strengths were in the same range

as those obtained by grinding, possibly being

on the low side.

Finally, gas polishing of polycrystalline B^C

and SiC was also attempted. These samples

were coated over a greater portion of their sur- faces with a fairly thick, black deposit; and,

in general, there appeared to be little, if any,

polishing action. The strengths of both the

B4C and SiC were of the order of one-third to

one-half that of as-ground samples. Thus, polishing of these materials does not appear to be

promising.

4. Summary and Conclusion

Very excellent quality finishes can be obtained on sapphire by gas polishing which is

not detectably influenced by crystallographic

orientation. These sapphire bars exhibit much

improved strengths, having bend strengths

averaging over 2 x lO'^ psi (1.4 x lO^N/m^),

with 5 X 10^ psi (3.4 X lO^N/m^) being

achieved in these initial attempts. Also the

strengthening effected by gas polishing can be

achieved much more reliably than can be obtained by flame polishing. In addition, some

rutile crystals developed very smooth surface

finishes and exhibit fairly high strengths. The

present results further suggest that development of this technique may also be fruitful

with spinel and possibly for other ceramics.

Gas polishing of a variety of polycrystalline

ceramic specimens yielded mat or frosted finishes and gave no clear strength increases or even decreases in strength. Thus, although

apparently limited to single crystals, the polishing technique has proven to be a very versatile

and useful procedure for obtaining good sur- faces and substantial increases in strength in

sapphire and rutile. Twinning may be limiting

further increases in strength of rutile.

Note added in proof: Recently, using only

czochralski grown sapphire, average flexure

strengths of 8 ± 0.3 X 10^ psi (—5.5 X

lO^N/mO with maximum strengths cf 1.2 x

10« psi (—8.3 X lO^N/m^) have been obtained.

Also, average bend strengths of 2.1 ± 0.25 x

10^ psi (-'1. 5 X 10^ N/m^) have been obtained

for rutile while strengths up to 1.5 x 10^ psi

(—'1.1 X lO^N/m^ for stoichiometric spinel

crystals were achieved. Fine etch pits also can

occur during polishing, especially with rutile,

which can either limit strength directly or may

act as twin sources.

5. References

[1] Schmidt, W. A., and Davey, J. E., Preparation of

smooth, crystalline, damage-free, sapphire sur- faces by gaseous etching. This volume.

[2] Rice, R. W., The effect of sputtering on surface

topography and strength of ceramics. This

volume.

[3] Rice, R. W., Effect of grinding direction on the

strength of ceramics. This volume.

[4] Becher, P. F., and Rice, R. W., Flame polishing

of flat oxide bars. This volume.

Discussion

Rice : One other piece of data. We tried this suggestion, I think if you have stoichiometric

experiment on alumina rich spinel and there is spinel, gas polishing may give you high strength

some possible improvement in strengths. As a also.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic MAcmNiNc and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

269

IV. SURFACE AND SUBSURFACE ANALYSIS

AND CHARACTERIZATION

Session Chairman

W. H. Rhodes, AVCO Corporation

Analysis and Characterization of Ceramic Surfaces

for Electronic Applications

R. C. Sundahl and L. Berrin

Bell Telephone Laboratories, Inc., Allentown, Pennsylvania 18103

The criteria which are used to evaluate the surfaces of ceramic substrates for

use in the electronics industry must be related to the specific application of the

surface. This article emphasizes one such application—the use of a ceramic surface

as a support for complex thin film conductor patterns which serve to interconnect

silicon integrated circuit chips. Those ceramic surface parameters which are found

to be critical to this application are (1) topographical properties, (2) chemical properties and (3) crystallographic properties.

These properties are characterized using such tools and techniques as optical,

transmission electron and scanning electron microscopy, profilometry, electron microprobe analysis, ion microprobe analysis. Auger electron spectroscopy and electron

diffraction. Special emphasis is placed on the relationship between the results of

such analyses and the performance of the surface as part of the SIC interconnect

system.

Key words: Aluminum oxide; Auger electron spectroscopy; crystallographic texture;

ion microprobe; scanning electron microscopy; substrates; surface defects; surface

segregation; thin film adhesion; thin film circuit imperfections.

1. Introduction

The analysis and characterization of sur- faces resulting from ceramic forming, machining, and finishing methods are necessary to

obtain a fundamental understanding of the

effects of such operations on surface properties.

It is through this understanding that ceramic

surfaces can be improved so as to meet the

rigid requirements of the new technologies. In

the electronics industry the use of a ceramic

surface as a substrate for a thin film circuit

imposes certain demands on the properties of

that surface [1-5] \ Thin film properties are

particularly sensitive to a v^^ide variety of

physical and chemical properties of the substrate surface, hence a wide variety of techniques have been developed to evaluate these substrate properties. In the discussion of more

familiar techniques, such as the optical microscope and stylus instrument, the article emphasizes problems associated with the interpretation of observations on ceramic substrates.

Several techniques somewhat less familiar in

their application to ceramic surfaces are also

discussed.

' Figures in brackets indicate the literature references at the end

of the paper.

Before describing techniques for the analysis

of ceramic substrate surface properties it would

be instructive to first discuss in brief the cer- amic surface as part of a thin film circuit.

Figure 1 shows a typical silicon integrated circuit package, a 1,024 bit IGFET (insulated

gate field effect transistor) memory used in the

Electronic Switching System [6] . A high density

alumina substrate, about 80 mm x 30 mm

X 1 mm thick, provides support for the Ti-PdAu thin film conductor pattern. This thin film

circuit, fabricated from a vapor deposited

metallization by photolithographic methods,

provides the electrical and mechanical interconnections between the forty beam lead integrated

circuits. The circuit also provides access to and

from the outside world through the leads which

are thermocompression bonded at the edge of

the substrate.

The adhesion of the metal interconnect film

and the integrity of the finely detailed interconnect pattern are critically influenced by the

ceramic surface properties. Other factors important in determining compatibility of the

ceramic with the integrated circuit assembly

are resistance to chemical and thermal degradation during processing, high electrical resistance, appropriate dielectric constant and

271
Figure 1. An alumina substrate, approximately 3cm X 8cm, supports

the thin film interconnect circuit for a 102Ji. hit memory. The 32

chips in the top center of the circuit are 32 bit memory chips, each

with 26 leads, 50 ixm wide and on 100 ixm centers [6].

sufficient mechanical strength to withstand

bonding operations and handling.

The techniques for analyzing and characterizing the ceramic surface can be divided into

three general categories of surface properties:

(1) Surface topography—the techniques used

to ascertain the surface roughness and its

effect upon the delineation of fine metal

lines on the surface. Also included are the

number, size distribution and identity of

discrete surface defects. These defects

have been found to have a critical effect

on thin film circuit pattern integrity.

(2) Surface chemistry—techniques used for

a direct analysis of the chemical composition of the ceramic surface, which, in

general, differs from the bulk ceramic.

The surface chemistry of a substrate has

been found to have an important influence

on circuit characteristics such as film adhesion, aging stability of thin film resistors and leakage currents across the

substrate surface.

(3) Surface crystallography—techniques discussed here include methods of characterizing crystallographic texture; that is,

the preferred orientation of the individual crystallites. Also covered are methods for studying the crystal structure on the surface, which in general

differs from the bulk structure, and the

mechanical damage introduced into the

surface by machining operations. There is

not a great deal of information available

on the surface crystallography of ceramics, probably because it has been felt in

the past that these properties are not important to substrate performance.

The techniques described here are discussed

in relation to their use for a particular problem,

the characterization of substrate surfaces for

thin film interconnect circuits. However, many

of these techniques have potential usefulness

in a wide variety of problems associated with

ceramic surfaces.

2. Surface Topography

One of the most common techniques used to

characterize surface topography is the stylus

trace. The technique provides a rapid method of

evaluating surface roughness for routine work

and the results can be characterized by a single

parameter, the AA or arithmetic average. However, one is often frustrated in his attempt to

relate the results of a stylus trace to the performance of the ceramic surface in a thin film

circuit package. It has been found that a more

detailed assessment of surface topography is

necessary to enable one to predict substrate

surface performance. This has been recognized

by White et al., [7] in an earlier paper, in

which they describe the use of computer processed scanning electron microscope images and

rastered stylus traces. These techniques are

still limited to relatively small areas, as measurements are time consuming.

We have combined rather conventional optical examination techniques, under a low 'power

microscope, with a statistical analysis of the

observations. This analysis provides information about the number and size distribution of

discrete defects in the substrate surface, as

well as the effect of these defects on thin film

interconnect patterns. More detailed information about the defects and other features of the

ceramic surface have been provided by scan- ning electron microscope (SEM) and transmission electron microscope (TEM) observations.

2.1. The Stylus Trace

The operation of this instrument is simple

in concept. A diamond stylus with a tip which

has a radius of curvature typically 2.5 ixm, is

272
brought into contact with the sample surface

under a load of less than 500 Kg/cm-. As the

stylus is moved across the surface, the vertical motion of the stylus is converted into an

electrical signal by a displacement transducer,

the signal amplified, and then displayed on a

strip chart recorder. Amplification, the ratio

of recorder pen motion to stylus motion, varies

from 1,000 to 50,000. The highest gain corre- sponds to a sensitivity of about 50A vertical

displacement of the stylus. Horizontal displacement of the stylus is directly related to

motion of the chart paper and amplification of

the horizontal scale is typically 100, or 10 to

500 times less than the verical scale. This ac- centuation of the vertical motion of the stylus

should be kept in mind v^^hen examining stylus

profiles. Deviations in the surface of only a few degrees slope appear as very sharp peaks

when viewed on the stylus trace.

Typical recorder traces of an alumina substrate surface are shown in figure 2. This

sample was prepared by an isostatic pressing

process [8] and sintered to 99 percent of its

theoretical density. It is instructive to study

these traces and determine their usefulness.

Figure 2A, an as-fired surface, shows the typical small modulations in the profile and peaks

and valleys of approximately equal size. After

grinding (fig. 2B) the peaks and the undulations are obviously removed. The valleys in

the profile probably represent regions of grain

pullout. Careful polishing with diamond paste

essentially eliminates pullouts (fig. 2C). This

last profile, with a ten-fold increase in vertical

magnification, shows only an occasional pit,

about 0.3 nxn diameter, on the surface.

The stylus trace can be characterized by

several parameters; namely, (1) an average

peak-to-valley distance, (2) the maximum peakto-valley distance, (3) the average half-height

widths of the peaks, and (4) the average number of peaks per unit length. Although a single

parameter provides only a very limited description of surface topography, it is quite common

to characterize surface roughness by the arithmetic average (AA). The AA is determined

electronically from the transducer output by

summing peak areas above and below a zero reference point, the center line, and dividing by

the stylus trace distance.

The stylus trace does provide a convenient

method for monitoring ceramic surface preparation processes, for changes in ceramic prop1.2 7pn

-50l

:t:r.

-HI - 1

-

1270 A

^

1 —^1 3 —

Figure 2. Surface stylus traces of alumina substrate surfaces in three

stages of finishing; (A) As-fired, AA = 0.27 /jltk (10.5 fiin) (B)

Ground with a ^00 grit diamond button wheel, AA = 0.11 fi.m (U.5 /j-in), (C) lapped with a diamond polishing compound, AA < 0.005 fxm (0.2

liin).

273
erties or process steps will generally be

reflected by corresponding changes in the

appearance of the stylus trace and the AA

value. However, it is difficult to characterize

surface topography quantitatively with a stylus

trace. First, the common use of the AA value

alone has some serious drawbacks. Two sur- faces with widely differing stylus traces may

have identical AA values [5, 9]. Also, low AA

or other average values may underweight the

effects of a few large surface defects (see fig.

2C) which may have a critical effect on thin

film conductor or capacitor pattern yield. Finally, the stylus instrument provides a sample of

only a very small area of the surface and without a great amount of effort a few large defects

on a single substrate surface would go undetected

2.2. The Optical Microscope

Optical microscopy techniques are well documented in the literature. An example of one

such technique is described by Lo [10] in

another paper presented at this Symposium.

Our concern in this discussion is with the

effect of surface defects on the integrity of

thin film interconnect circuits [11]. A circuit,

such as seen in figure 1, may contain several

meters of 50 /xm lines on 100 /xm centers. Surface defects affect the photolithographic fabrication process so as to cause opens in conductor lines or shorts between adjacent lines.

Observations on the effect of various defect

types were made by use of fine line test patterns

which simulate the worst case conditions en- countered in an interconnect circuit.

The test pattern used consists of 15 m of 50

/xm conductor lines spaced on 100 /xm centers.

This pattern is generated on a 20 cm^ area of

p.-

substrate using standard thin film (0.1 /xm Ti

and 1.5 /xm Au) and photolithographic techniques. These patterns are examined in detail

under a binocular microscope at 60X, using

both reflected and transmitted lighting. Examples of surface defects and their effect on the

test pattern are seen in figure 3. A surface pit

in figure 3A has resulted in a complete open in a conductor line. A burr or protrusion, as seen under low angle reflected lighting conditions in

figure 3B, has led to a short between several

adjacent lines. This latter defect can easily

be detected with the unaided eye, using an intense light source with a low angle of incidence.

Observations recorded during such an examination include (1) the number of various defect

types and (2) the number, size and types of

imperfections produced in the conductor lines

by these defects.

There are two implications of this approach

which are important to this discussion. First,

the ceramic surface can now be characterized

by a parameter directly related to its performance, rather than some parameter only vaguely

related to use, such as surface roughness. The

parameter chosen here is fine line imperfection

frequency f(A,d). Here, f is a function both of

the width d of the conductor lines of the test

pattern (50 /xm in this example) and of the criterion used to define a conductor line imperfection. The criterion used here is any short

between adjacent lines or any line which is

more than 1/3 open (A. = 1/3).

The second implication is that this quantitative parameter can be used to predict performance of the ceramic surface in a given

thin interconnect circuit. If one assumes, rea- sonably, that substrate surface defects are randomly distributed over the substrate, then

Figure 3. Optical micrographs of alumina surfaces with a conductor

line test pattern present. (A) A surface void, resulting in a complete

open in a conductor line, (B) a burr protruding from the surface,

resulting in shorts between adjacent lines. Low angle incident lighting.

274
the probability that a len^h of line L will have

no line imperfection is [11]

Y = e-^f

This is, of course, identical to the production

yield of a thin film interconnect circuit with

conductor line length L, fabricated on a substrate with a characteristic / value. This statistical treatment can be extended to permit

estimates of the effects of changes in A and d

on yields [11].

One positive consequence of this approach

is the presence of a reference grid on the surface. It has been our experience that visual

observations are much more accurate if the

test pattern, which acts as a reference grid,

is present. The grid also provides a convenient

reference to estimate surface defect sizes. The

method unfortunately is very tedious, especially in cases where there are only a few defects

in a 100 cm^ area. However, attempts to automate these measurements with the aid of an

image analyzer thus far have not met with success, mainly because of the complexity of the

analysis and the required contrast and resolution. Therefore, despite the large amount of

effort required to obtain statistically significant data and the systematic human error

associated with such a technique, we have found

this statistical method of characterizing ce- ramic surfaces to be a most useful approach.

2.3. TTie Scanning Electron

Microscope (SEM)

Although the SEM was introduced into general use only recently, it has rapidly become a standard tool for surface characterization. The

SEM is particularly useful for providing more

detailed information about the nature of the

ceramic surface because of the unique characteristics of the instrument. These characteristics include (a) high resolution—about 500

is nominal, (b) unusually high depth of field

about 300 times that of the optical microscope

and (c) little or no sample preparation. The

details of the operation of the SEM has been

covered earlier in this Symposium [7].

Figures 4, 5 and 6 illustrate the application

of the SEM to the characterization of surface

defects [11]. These observations allow one to

determine the source of substrate defects and

to modify process variables so as to minimize

their occurrence. A common surface defect is seen in figure 4. At the lower magnification we

see a ground ceramic surface with a surface pit

present. In this case the pit had no effect on the

conductor pattern. At the higher magnification

the microstructure inside the pit is revealed in

greater detail. The smooth, equiaxed grain

structure is typical of an as-fired surface and

suggests that the pit was formed during or, more probably, before the sintering step. Electron microprobe analysis of this defect reFlGURE 4. SEM pictures of surface void observed on a ground and annealed isostatic pressed substrate with a

fine line test pattern present. The smooth equiaxed grain structure within the void indicates formation prior

to or durmg sintering. (A) lOOOX (B) 5000X.

275
Figure 5. SEM pictures of surface void of a polished (unannealed) isostatic pressed substrate with test pattern

present. The platelet structure within the void indicates the presence of foreign contamination. (A) lOOOX.

(B) 5000X.

Figure 6. SEM pictures of a large white spot observed on a polished and annealed isostatic pressed substrate.

The fine grain structure indicates that only partial sintering occurred in this region. (A) 500X. (B) 2000X.

vealed no significant contaminants in this type

defect.

A second and less common type of surface

defect, seen in figure 5, exhibits an anomalous

morphology inside the pit. This growth of acicular platelets is indicative of a local contaminant. This conclusion was substantiated when

an electron microprobe examination revealed

a high concentration of calcium within the void.

It would appear then that this type of defect is

somehow associated with contaminants in the

raw material.

A very common defect observed in high density alumina ceramics is an isolated low density

276
region. Figure 6 shows the catastrophic effect

of this defect upon a 50 /xm wide test line. It

is seen that this region is actually made up of

very small individual grains only partially

sintered together. It is suggested that this phenomenon is the result of hard agglomerates

existing in the raw material which retain their

integrity during the ceramic forming and sintering operations. The defect type exhibited no

measurable contamination when examined in

the electron microprobe.

There are some problems associated with the

use of the SEM for the examination of ceramic

surfaces and these warrant a brief review. On

insulating samples the electron beam charges

the surface negatively. If this charge is not

dissipated the picture is distorted and local

bright spots appear. This problem is easily

overcome through the use of a thin conductive

coating, a few hundred angstroms of carbon

will do, on the surface of the insulator. One

must be careful, however, to completely cover

the area to be examined. An examination of

figure 4 reveals that an area inside the pit was

shadowed during vacuum evaporation of carbon

onto the sample, resulting in local charging and

bright spots.

A second area of difficulty is the interpretation of SEM images. The contrast effects are quite complex. In the standard secondary

emission mode of operation contrast results

from variation in secondary emission yield

over the sample surface. As discussed in detail

by White [7], the yield increases as the angle

of incidence of the beam with the surface decreases, giving a topographical contribution to

the image. However, secondary yield also varies

in a complex manner with surface composition,

which in turn may vary from sample to sample

or even over the surface of a presumably homogeneous sample. These effects are not completely understood and can lead to confusion.

One attractive feature of the SEM is its

versatility. There are several other modes of

operation [12] available on most commercial

instruments, in addition to the secondary emission mode. These include the "backscatter"

mode, in which only the elastically scattered

electrons are detected. A combination of information from the backscatter and secondary

modes can resolve some ambiguities in interpretation of contrast effects. An inclusion of

information from the sample current image

may lead to further understanding of the nature

of the surface. The SEM can also be used to

study cathodoluminescence and some instruments are equipped with dispersive or nondispersive x-ray detectors, converting the

instrument into an electron microprobe.

2.4. The Transmission Electron Microscope

(TEM) [13-15]

Again, the principles and uses of the TEM

are well-documented in the literature. The purpose here is to briefly describe techniques we

have found useful in studies of ceramic surface

topography.

When a high degree of resolution is sought

and the surface to be observed is not overly

rough, replica TEM techniques provide micrographs superior to those obtained from the

SEM. Thus the method is useful for polished

and as-fired surfaces, but is difficult to employ

for the examination of ground surfaces or sur- face defects.

The sample preparation for the TEM is cer- tainly more complex than for the SEM. Two

methods that are useful for replicating a

surface are the one-step plastic technique and

the two-step plastic-carbon replica technique

(see fig. 7). The one-step method is relCERAMIC SURFACE TO BE REPLICATED

SURFACE WITH

PLASTIC IN PLACE

STRIPPED REPLICA

SHADOWING

DIRECTION

r

CARBON

DIRECTION

PLASTIC '-METAL STEP I PLASTIC REPLICA

^ETAL

'='^^H^f=ls*? CARBON

STEP 2 FINAL REPLICA

Figure 7. Two common techniques for

fabrication of transmission electron

microscope replicas of ceramic surfaces. (A) One-step plastic replica,

(B) Two-step plastic-carbon replica

with heavy metal shadowing.

atively simple. The replica is formed by

placing on the ceramic surface a dilute solution of collodian or Formvar and allowing the

solvent to evaporate. The film is removed from

the surface by stripping it off with adhesive

tape or floating it off under water. A replica

prepared in this manner is a negative reproduction of the original surface. This technique

is not applicable to very rough surfaces as the replica is quite fragile. Also, the plastic

tends to decompose during long exposure to the

electron beam. The two-step process utilizes a

thicker first replica which then acts as a substrate for a thin carbon replica. A positive

439-921 O - 72 - 19 277
replica results from this procedure. The twostep process is useful for rougher surfaces and

the carbon is quite stable in the electron beam.

Figure 8 shows the result of a one-step replication of the surface of a commercially available tape cast substrate. This method appears

entirely satisfactory for such application as

fine detail is revealed. Some artifacts are evident in the micrograph, some of which are due

to tearouts during the film stripping.

Figure 8. Transmission electron micrograph of an

as-fired aluminum oxide surface replicated with

plastic (8600X).

In the TEM electron absorption is not very

different for the small differences in specimen

thickness and very little difference in contrast

results from thickness effects. Contrast in the

TEM arises primarily by differential scattering from different parts of the sample. Since

the relative degree of scattering is governed

primarily by differences in the mass density of

a specimen, the most useful method for enhancing contrast is the shadowing technique. This

method consists of increasing the density on

certain sides of the specimen by oblique vac- uum coating with a heavy metal, so that the

beam is scattered readily as it passes through

those portions of the specimen. It is essential to know the shadow angle so as to obtain

quantitative information; i.e., height and depth

measurements, of the surface topography.

3. Surface Chemistry

In the past it has been the general practice

to determine the effect of the chemical properties of the ceramic substrate surface through

some operational measurement of substrate

performance, such as adhesion measurements,

aging experiments or environmental tests. The

role of chemistry has then been inferred indirectly from the results of the experiment and,

in some cases, knowledge of the bulk chemical

composition of the substrate. Recently, however, there has been a proliferation of new analytical techniques which provide direct information about the chemistry of the surface.

These include the electron beam and ion beam

methods, such as electron microprobe analysis,

Auger electron spectroscopy, photoelectron

spectroscopy and ion microprobe armlysis. These

tools can be combined with the more indirect

methods of analysis, such as infrared spectroscopy, ellipsometry, contact angle measurements, gas adsorption measurements and the

above mentioned performance tests to provide

a more complete picture of ceramic surface

chemistry.

3.1. The Electron Microprobe [16-18]

The electron microprobe is, of course, not a new method of chemical analysis. In fact, in

a strict sense, it is not a method of surface

analysis at all. However, a description of the

principles behind electron microprobe analysis

provide a basis for discussing other beam techniques.

A schematic of the electron microprobe is seen in figure 9. The optics are similar to that

of the SEM; the electron beam is rastered across the sample and a display is formed co- incidently on a cathode ray tube. In this case,

though, the information is the intensity of

emission of a particular x-ray line. This instrument may also be used as a crude SEM, just as the SEM can be used as a crude electron

microscope. The basic difference in the two

instruments is the lower resolution in the electron microprobe. The electron microprobe is

limited to about 1 pxa beam spot rather than a

nominal lOOA for the SEM. This is compatible

with the inherently lower resolution of x-ray

emission mode of operation.

Figure 10 illustrates the complex processes

which occur when a high energy electron beam

interacts with a sample surface. Application of

the microprobe to studies of ceramic surfaces

requires that the investigator keep in mind several considerations illustrated by figure 10:

(1) The incident beam penetrates several

microns into the surface, losing energy

through the generation of heat and sec- ondary electrons. This heat may cause decomposition of the sample if it is a

278
Figure 9. Schematic diagram of the electron microprobe.

PRIMARY BEAM -E(,

Figure 10. Schematic diagram of electron loss and

x-ray emission processes in the electron microprobe.

Primary excitation occurs in a region in which the

diffusing electrons maintain an energy E greater

than the threshold energy E'th (A£">0).

poor thermal conductor. Radiation damage due to the incident beam may also

modify the sample. The generation of

secondary electrons results in the sur- face charging of a ceramic insulator.

This again can be eliminated by the use

of a thin conductive coating. Finally,

penetration of the beam means that x- rays are generated from a depth of

several thousand angstroms rather than

from the surface alone.

(2) Diffusion of primary electrons in the

sample and the generation of secondary

x-rays results in loss of lateral resolution, v^hich is on the order of 2 /xm.

(3) The complexity of the various processes

makes it difficult to obtain a quantitative interpretation of the data. This

problem is exaggerated for the lighter

elements as absorption corrections introduce the most error. However, there are available elaborate iterative computer programs which do provide quantitative data from the microprobe measurements [19].

(4) The measurements of the light elements

are also limited by low signal yield, absorption by the conductive coating and

low sensitivity of the detectors to long

wavelength x-rays.

Applications of the electron microprobe to

the study of ceramic surfaces has thus far been

limited to the semiquantitative analysis of inclusions and other macroscopic forms of con- tamination [18], to interdiffusion at metalceramic interfaces [16] and to the segregation

of impurities at grain boundaries [17]. The re- sults of grain boundary segregation analyses

have some implications relative to the ceramic

surface, as one would expect that impurities

that segregate at internal surfaces; e.g., grain

boundaries, might also segregate on as-fired or annealed surfaces.

3.2. Auger Electron Spectroscopy

(AES) [20-26]

The technique of Auger electron spectroscopy

(AES) is a new addition to the catalog of tools

available for the analysis of surfaces. The instrument is in many ways similar to the electron microprobe. In both instruments an incident beam of electrons is used to excite the

atoms and the surface chemistry is characterized in terms of the decay products of these

excited atoms. The main difference is that the

decay products detected by AES are electrons,

not photons. In other words, AES is an electrons in, electrons out technique where the electron microprobe is an electrons in, photons out

instrument.

The characteristics of the AES of primary

interest for the study of ceramic surfaces are:

279
(1) The instrument has a high selectivity

for the surface atoms. For electrons

with an energy of 500 eV, a typical en- ergy of an Auger electron, one expects

a mean escape depth of a few atom

layers.

(2) AES is able to detect less than 0.01

monolayer of contaminant on the surface.

(3) Below 1000 eV virtually all relaxations are associated with Auger emission,

thus the relatively high sensitivity of

AES to the presence of light elements.

Applications of AES have been reported in

metallurgy (interdiffusion of thin films, grain

boundary segregation, surface diffusion), solid

state technology (photoemission from semiconductors, contamination on silicon wafers, segregation of impurities at surfaces), catalysis

(chemisorption on platinum), and solid state

and molecular physics (surface band structure,

molecular spectra). Unfortunately, very little

work has been done on the characterization of

ceramic surfaces with AES. This is undoubtedly due, in part, to the newness of the technique. However, there are some fundamental

limitations which may have impeded its application. These problem areas will be discussed

after a brief description of the Auger process

itself and some of the instrumentation used for

its detection.

The Auger process of electron ejection is a

relaxation process complementary to x-ray

emission. Figure 11 A shows an excitation process, the same as for x-ray emission, in which

an incident electron ejects an electron from a lower energy shell. The electron beam characteristics are typically 3 keV and 100 ixSl. X-ray

emission (fig. IIB) is seen to result from the

relaxation of an excited atom, in this case magnesium, through an electron transition and the

simultaneous emission of a photon. In figure

lie the same excitation and decay process is

seen, but this time accompanied by the ejection

of an Auger electron. To conserve energy, the

ejected electron in this case has a kinetic en- ergy of K-Li — L'z, where L'., is the L2 level perturbed by the double ionization of the Mg atom.

This perturbation results in some uncertainty

as to the theoretical position of the lines, and

in most cases these positions have been determined experimentally [21, 22].

The case of a solid and, in particular, an insulator is complicated by the broadening of en- ergy levels into bands. In MgO, for example,

Mg^+ has no M electrons and consequently no

LMM transition would be expected for a purely

ionic compound. However, transitions are observed, and it is uncertain as to the origin of

the electrons which decay into the lower energy

level and the origin of the Auger electrons.

These transitions have been explained alterna-

(A) EXCITATION OF Mg ATOM (8) DECAY WITH PHOTON EMISSION

L|-V-V' fLi-Vo-Vi

» « K

1 • K •—• K

(C) DECAY WITH KLzLi AUOER (0) LMM TRANSITIONS IN MgO

TRANSITION

Figure 11. A comparison of Auger and x-ray emission processes for Mg, MgO. (A) The excitation of

the atom with an incident electron is common to both techniques. (B) Relaxation of excited atom

via x-ray emission. (C) Relaxation accompanied by

Auger ejection of an electron—a KL^Lj' transition. (D) Auger transitions in MgO.

tively in terms of a deviation from ionicity or crossover transitions from the 0= band, as shown in figure IID.

The major problem associated with the detection of Auger electrons is the low signal-tobackground ratio. The incident electron beam

produces a large number of secondary and

inelastic electrons over all energies below the

incident beam energy. The Auger electrons

make up a very small fraction of this total signal and herein lies the problem, the detection

of the weak Auger signals in a high background. A standard analog technique for solving such a problem is to examine the differential of the secondary electron distribution. The

slowly varying background signal has a first

derivative near zero, whereas the small, sharp

Auger peaks have a significant first derivative.

In this manner an increase in signal-to-background ratio of several magnitudes is gained

[20].

Figure 12 shows an Auger electron spectrum

of MgO. The first derivative of the energy distribution is plotted as a function of electron

energy. Note that the appearance of the major

Mg LVV transition is quite similar to the differential of a gaussian peak. The MgO spectrum

consists of the low energy Mg LVV transitions,

depicted in figure IID, as well as the high en280
t

200 300

BIAS VOLTAGE

Figure 12. An Auger spectrum of MgO. The low energy Mg LVV transitions are by far the most intense. A series of weak transitions are observed between 100 V and 300 V which have not been identified.

The oxygen and magnesium KLL transitions differ in intensity by

an order of magnitude as Auger yield drops off above 1000 V.

ergy oxygen and magnesium KLL transitions.

It is to be noted that, despite the similarity of

the KLL transitions for the two components,

the Mg transition is significantly lower in peakto-peak amplitude. This reduction in intensity

for higher energy transitions occurs because

(1) it is more difficult to excite the K level of

Mg than 0 with the 3 kV beam and (2) above

1000 eV some of the excited states decay via

photon emission.

Figure 13 is a schematic of a retarding field

analyzer used to perform AES measurements

[25, 26]. The sample, situated at the center of

the spectrometer, is exposed to an incident

beam of electrons, in this case supplied by a

modified cathode ray tube electron gun. The

electron optics of the spectrometer are conCOLLECTOR

ISOLATION TRANSFORMER

Figure 13. Schematic diagram of the retarding field analyzer.

structed from a series of nested spherical grids,

situated in front of the electron collector. A

bias voltage Vb applied to the grid structure

allows only secondary electrons from the sample with an energy greater than eVg to be collected. Thus this instrument provides a collector current i, which is the integral of the sec- ondary electron energy distribution above

cVb. In order to obtain a plot such as shown in

figure 12 a second derivative, dHc/dYs^, is needed. This double differentiation of ic is performed

electronically using ac modulation and detection techniques [25]. Briefly, the bias voltage

Vb is modulated with a small ac signal, resulting in a corresponding ac component in v. The

first harmonic of ic (at frequency wo) is proportional to dic/dVB and the second harmonic

(2(oo) to dHc/dVB^ The detector system is then

tuned to the 2wo component of i and the output

plotted vs. the sweep voltage Vb on the x-y re- corder. Figure 12 is a direct result of such a measurement. It is then the combination of the

retarding field analyzer, with its wide acceptance angle, the concept of measuring the derivative of the secondary electron energy distribution, and the use of modern ac signal processing techniques which have made AES a

practical tool.

As with other electron beam techniques there

are some problems associated with the application of AES to the study of ceramic surfaces.

These problems are categorized as follows:

(1) Surface charging of insulators. This

problem is analogous to that described

281
previously for the SEM and electron

microprobe. The problem is more critical here since the charge cannot be dissipated with a thin conductive coating.

For, if one used a 500A carbon coating

here, the only component that would be

detected is carbon. In general, for energies greater than 200 eV, the effect of a positive surface charge is to displace

the spectrum uniformly toward a lower

voltage. The extent of this voltage shift

can usually be estimated if one has a

strong peak of carbon or oxygen, for

example, to act as fiducial marks on

the spectrum. The spectrum at low en- ergies (< 100 eV) is more strongly

affected by sample charging. In general

this lower energy region of the spectrum should be avoided since complications, such as illustrated in figure IID,

make interpretation of low voltage

Auger spectra ambiguous.

(2) Surface contamination. In certain applications the impurities picked up

during machining, firing steps or cleaning operations may play an important

role in determining the performance

at a ceramic surface. However, contamination from the ambient is usually

only a nuisance. In this case it is best

to prepare the surface in situ in an

ultrahigh vacuum system. The sample

can be cleaned in the vacuum station

by ion beam etching or by the electron

beam itself (electron desorption) but

great care is needed so that the "interesting" contaminants are not re- moved with the "nuisance" contaminants.

(3) Quantitative analysis. The processes

involved in an AES study of ceramic

surfaces are not as well understood as for the electron microprobe. One is generally limited to empirical calibration

methods through the use of standards

to obtain quantitative information [21,

22]. Experiments of this nature have

demonstrated a linear relationship between the peak-to-peak amplitude of

the Auger signal and surface concentration for less than one monolayer coverage. Most work in this area to this

date has been in the form of semiquantitative correlations between the

measured signal amplitude and some

macroscopic property, such as ceramic

performance.

The next series of figures illustrate the application of AES to one problem associated

with ceramic surfaces. Alumina substrates

were prepared by grinding the surface with a 400 grit diamond wheel followed by an anneal

at 1500 °C for 30 minutes in air. The substrates

then went through a cleaning treatment commonly used prior to metallization and then were

immediately placed in the spectrometer. Figure

14 shows evidence of the contaminants left behind by the organic cleaning solvents and the

ambient. The surface is obviously masked by

a carbonaceous residue. A low oxygen peak is

detected and there is some evidence of calcium

present.

ISOSTATIC PRESSED SUBSTRATE

J \

/ CI LMM

/ S LMM

-ANNEAL 30MIN ISOCC

[/ N KLL

1 Co LMM

0 KLL >i

°

X. 75

-.5

C KLL

1 1 1 1 1 100 200 300 400 500 600

Figure 14. Spectrum of isostatic pressed substrate sur- face after a 30 min anneal in air at 1500 °C. Sample area not previously exposed to electron beam.

A weak calcium peak is observed on this surface,

besides contamination peaks associated Avith S, CI, C

and N.

After the substrate surface is cleaned with

the electron beam the Auger spectra shown in

figure 15 was revealed. The increase in the

oxygen and calcium signal is evident, as well as the corresponding decrease in the carbon

signal. There are also charging effects at low

voltages. A careful analysis of the low voltage

spectrum reveals a modified Al spectrum with

a weak Si peak superimposed on it. Although

a bulk chemical analysis of the substrate shows

that about 650 ppm Mg is present, there is no

indication of any Mg from Auger emission. By

measuring the Mg signal from MgO (fig. 12)

a rough estimate of the sensitivity of the instrument to Mg segregation is provided. It is

estimated then that there is less than 10 percent Mg present on the alumina substrate surface. This is not too surprising as MgO may

volatilize from the substrate surface during the

annealing procedure.

The presence of a Ca signal on annealed

substrates led to some interesting results

regarding impurity segregation at the ceramic

surface. An Auger spectrum of an as-ground

282
ISOSTATIC PRESSED SUBSTRATE - ANNEAL 30 MIN, 1500'

-

- EXPOSED TO E -BEAM 1 HR

/I' f/M+Si LVV

// 1 Al KLL

HAI LVV ' J

'

- X.2 j" CO LMM

1 XlOO J X.75 1 I

1200 1300 1400

1

1 1 1 1 1 1

0 KLL

1

1 1 40 60 80 100 200 300 400 500 600 700

BIAS VOLTAGE

Figure 15. Spectrum of same sample as Figure H after a one hour

exposure to the electron beam.

sample (not annealed) showed no evidence of

calcium. Annealing the substrate for longer

times at 1500 °C brought out the calcium peak.

Figure 16 clearly shows the evidence of this segregation of Ca on the alumina surface.

0.3

Figure 16. Normalized AES signal

amplitude for the Ca LsMijM^s

peak on the isostatic substrates

vs. annealing time in air at 1500 °C.

3.3. Photoelectron Spectroscopy [27—29]

The technique of photoelectron spectroscopy

is similar in many respects to Auger spectroscopy. It differs in that it is a photons in, electrons out method, as shown in figure 17. The

excitation source is an intense beam of x-rays

(usually an aluminum or magnesium target)

and the emitted photoelectrons (or Auger electrons) are analyzed with an elaborate, high

resolution electron spectrometer. The attractive feature is that the energy of a photoelectron is directly related to one energy level. The

instrumentation has been developed to exploit

A

)

hv

lU /

^ -/—

/ * • */

T—• * /

}(-• K

Figure 17. Emission of a photoelectron from

an Mg atom.

the inherently high resolution of this technique.

Finally, since there is no incident electron

beam, there is little sample charging and a considerable reduction in secondary electron emission and background noise.

The sensitivity of this technique is generally

not as high as Auger spectroscopy since the

cross section for x-ray excitation is low. The

spatial resolution in photoelectron spectroscopy

does not compare favorably with either Auger

spectroscopy or the electron microprobe. The

sample area exposed to the x-ray beam is about

1 cm-. Finally, the surface selectivity is not as good as AES. Photoelectron energies normally

are on the order of thousands, rather than hundreds, of electron volts and correspondingly,

average escape depths are about lOOA.

This high resolution spectrometer can also be

used to study Auger transitions using both

x-ray and e-beam excitation modes. An exam283
pie of a photoelectron Auger spectrum for the

Mg KLL transition [27] is seen in figure 18.

A higher resolution is obtained (five transitions

resolved) than by corresponding AES methods.

AMPLIFIER

MAGNESIUM KLL AUGER

1060.0 1080.0 I 100.0 IliO.O 11400

KINETIC ENE^OY

Figure 18. The Mg KLL Auger spectrum measured

on a high resolution electron spectrometer [27]. Five

of the predicted six peaks are detected. The sixth

peak, at 1183 V, is too weak to be detected.

The high resolution of photoelectron spectroscopy has resulted in its application to a large

number of problems associated with chemical

bonding, especially with organic molecules. Experiments have also provided information about

the band structure of metals, alloys and compounds. Although no literature is reported on

its application to ceramic surfaces, the technique does provide a unique method by which

detailed information about the chemical structure near the surface can be obtained. There

is, therefore, every reason to believe that this

technique will find application to problems re- lated to ceramic surfaces.

3.4. The Ion Microprobe [30, 31]

A potentially very powerful tool for surface

characterization is the ion microprobe. This

instrument uses an incident beam of ions to

sputter material from the surface. The sputtered ions are analyzed with a mass spectrometer. One version of the instrument is

shown in figure 19. An intense 2 fjxa diameter

beam of ions can be rastered across the surface

supplying a two dimensional display of the

output, similar to the electron microprobe and

SEM. Other versions [30] use a stationary ion

beam about 200 jxm in diameter. Etching rates

on the order of 5A/s or higher are available.

The sensitivity of the ion microprobe is

unparalleled. It is determined by the beam

current and mass spectrometer design. Sensitivities of parts per billion are claimed for a monolayer of sample. The surface selectivity is,

by definition, confined to the first layer of atoms.

MULTIPLIER L

MAGNET ION SOURCE \

CONDENSER LENS ISOTOPE RATIO SWITCH

BEAM SWEEP

OBJECTIVE LENS

SAMPLE

Figure 19. Schematic diagram of the ion microprobe.

In addition to the high selectivity, it is a

unique method for determining a composition

profile as the ion beam etching proceeds. A

problem, however, is that ions are contributed

by the sides, as well as the mottom of the etch

pit and this tends to smear out the composition

profile. The apertured detector can be adjusted

to only accept ions from the bottom of the pit

but this has an associated problem of beam

stability. Another method is to etch by rastering the beam across the surface. This eliminates the problem but results in a very slow

etch rate. An additional feature of the ion microprobe is its ability to perform an isotopic

analysis, enabling one to quickly and accurately

determine diffusion profiles.

Under carefully controlled conditions, the ion

microprobe can provide accurate quantitative

information from ppb up to the mole fraction

range of concentrations. Figure 20 shows the

results of an investigation [32] in which the

bulk composition of a sample of moon rock was

accurately characterized over a wide range of

concentrations of its constituents. The ion microprobe with its obvious potential for the

study of ceramic systems has found limited use

to date because of the complexity and associated high cost of the instrumentation.

The techniques described above provide direct information about the chemical composition of the surface. There are also a series of

methods that will not be discussed, but the

reader should nevertheless be made aware.

These methods provide indirect information

about the surface chemistry and include infrared spectroscopy, ellipsometry, contact angle

measurements, and gas adsorption and desorption measurements. Although the data obtained

cannot be directly associated with specific sur- face components, these techniques are often

convenient to perform and in many instances are more sensitive to subtle changes in the sur- face chemistry than are the beam techniques.

284
I

I Ill I I I I I 1 11 iiiil I I 1 1 1 ml I I M I ill 10 100 1000 I 10 100

^ V V

' ppm %

ION MICROPROBE ANALYSES OFGLASSES

Figure 20. Ion microprobe analyses of lunar material

vs. standard chemical analyses of the same components, as performed in several laboratories using

the optimum technique for each component 132}.

4. Crystallography of Ceramic Surfaces

In this section three techniques will be con- sidered that are used to characterize the crystallography of ceramic surfaces. These techniques are low energy electron diffractiori

(LEED), reflection high energy electron diffraction (RHEED or RED), and x-ray texture

analysis.

In many cases the crystallographic structure

of the surface is identical to the bulk. Exceptions, however, can be of some importance, particularly in the field of electronic ceramics.

Studies on single crystal surfaces have shown

the surface crystallography to be reconstructed

within the first few monolayers. This surface

structure has been shown to be sensitive to the

presence of impurities and surface treatment.

Machining operations have been shown to introduce microscopic changes in the crystallography at the surface, increasing the dislocation density and associated residual strains.

On a macroscopic scale, polycrystalline ce- ramics in some cases exhibit crystallographic

texture, or preferred orientation of the individual grains. Examples of texture in bulk ceramics include hot pressed materials, ferrites

sludge pressed in a magnetic field and the

pressing of acicular particles. There is recent

evidence that such effects are also associated

with the surface of tape cost substrates. Such

texture effects may be of importance for any

application of an anisotropic material, such as

alumina, when a second rank tensor property;

e.g., dielectric constant or thermal expansion

coefficient, is critical.

4.1. Low Energy Electron Diffraction

(LEED) [33-35]

The first low energy electron diffraction ex- periments were performed by Davisson and

Germer [36] in 1927, demonstrating the wave

nature of electrons. Only recently, however,

has the technique become a useful tool for the

study of crystal surface phenomena. The instrumentation is similar to the AES spectrometer (fig. 13). In fact, the AES spectrometer

is a derivative of the LEED apparatus. In

LEED a normal incident beam of slow electrons (5-500 eV) interacts strongly with the

surface and yields a reflection diffraction pattern, which is displayed on the phosphor-coated

electron collector.

There are several consequences of the low

energies involved.

(a) The resultant low penetration depth of

electrons means that only the first few atom

layers contribute to the diffraction pattern. The

pattern then is truly representative of the surface.

(b) By satisfying the Laue conditions for diffraction one always obtains a two dimensional

diffraction pattern for any orientation of a

single crystal. Thus, no pattern is obtained for

a polycrystal, limiting LEED to the study of

single crystal surfaces.

(c) Quantitative diffraction theory is inadequate to predict the atom arrangements from

the diffraction pattern. Usually LEED is used

to detect changes in the symmetry of the sur- face crystal structure.

(d) This technique is very sensitive to surface

contamination and, as such, experiments re- quire an ultrahigh vacuum system.

LEED has been successfully used to study

structures of insulating crystal surfaces.

Chang [37] and Charig and Skinner [38] studied the surface structure of the basal and

prismatic planes of sapphire and the epitaxial

growth of silicon on these surfaces.

4.2. Reflection High Energy Electron

Diffraction (RHEED) [35,39-42]

This technique is applicable to a polycrystalline as well as single crystal surfaces. There is

little literature reported for ceramic materials

but RHEED has potential utility despite some

experimental difficulties. The experimental ar- rangement is shown in figure 21. A high en- ergy e-beam (10 keV to 100 keV) impinges on

the sample surface at a low angle of incidence.

Alignment of the e-beam is critical and the familiar problem of surface charging exists. To

neutralize the surface charge, the sample can

285
S'Giimpl'e

imuB'r'iiiip'ul'Qi't'or

magnetic

lens

electron

gun

def I ec tioipi plates

deflection plates phosphor

screen

r ^ ^ 1

Figure 21. Schematic diagram, of reflection high energy electron diffraction system.

be flooded with a low energy electron beam. As

with LEED, the RHEED method is sensitive

to the first few atom layers, is aifected by sur- face contamination and it is best to conduct

RHEED experiments in an ultrahigh vacuum

system (P < 10"^ torr).

A consequence of the low angle of incidence

is the high sensitivity of the diffraction pattern

to surface topography. Two extreme examples

are illustrated in figure 22. The diffraction pattern of an atomically flat surface (fig. 22)

consists of a series of vertical streaks, a conse- quence of the low penetration depth of the beam

into the surface (about lOA) and refraction.

Diffraction from a rough single crystal surface

(fig. 22B) resembles a transmission diffraction

geometry more than a reflection geometry.

This change in microscopic geometry leads to

the familiar spot diffraction pattern. In most

cases the geometry is a combination of the two

situations. In the case of a rough polycrystalline surface the spot pattern degenerates into

Debye rings (fig. 22C). The sharp rings with

uniform intensities about their circumference

are indicative of a structure made up of ran- domly oriented, highly perfect crystallites.

A major advantage of RHEED then is the

sensitivity to surface structure, topography

and orientation, combined with the applicability to polycrystalline surfaces. These positive

aspects of the technique are balanced by some

negative aspects. Sample preparation is difficult. The diffraction pattern tends to overweight

those components which protrude from the surface. Finally, the experiment geometry and dynamical diffraction effects complicate any

attempt at quantitative interpretation of intensity data.

There are available in the literature ex- amples of the application of RHEED to a wide

variety of problems. The most common use of

this technique is the identification of phases

present on a surface or in a thin film. Using

A

B

Figure 22. Diffraction of high energy electrons from

a surface. (A) Diffraction from an atomically smooth

single crystal surface produces a rod diffraction

pattern. (B) Diffraction from a rough single crystal

surface produces a spot pattern. (C) A ring pattern from a rough polycrystalline surface.

techniques analogous to the classical x-ray

methods of phase identification, the Debye pattern can be indexed according to the d-spacings

and line intensities and the results compared to

tabulated data. The technique is limited by lack

of resolution, accuracy of the d-spacing meas286
urements (usually '—1%) and difficulties in interpreting line intensities. Line intensity measurements have been applied quantitatively to

studies of phase transformations and chemical

reactions occurring at a surface. The literature

is replete with examples of the use of RHEED

in studies of corrosion reactions. A detailed

study of the profile of the diffraction lines provides semiquantitative information about such

phenomena as mechanical damage to the sur- face [43], residual stresses in thin films and

surface crystallographic textures. Mechanical

damage and residual stress cause the profiles

to broaden and preferred orientation causes

the Debye rings to break up into arcs. The extreme limit of mechanical damage

would be an amorphous layer. Such a layer is

found to be present on metallic surfaces after

polishing and a considerable effort has been

devoted to the characterization of this so-called

Beilby layer [44]. As one might expect, the

diffraction pattern of an amorphous film represents the extreme limit of line broadening

only a diffuse electron diffraction pattern is

observed. The diffuse diffraction pattern provides a unique opportunity for quantitative

interpretation of intensity data, for dynamical

effects are no longer important. One example

of a quantitative treatment is the work of

Drobek [45] in which he analyzed the structure

of amorphous thin films of SiOa deposited by

various techniques on silicon wafers. Two of

the diffraction patterns obtained by Drobek are

seen in figure 23. A diffuse pattern, characteristic of an amorphous material, was observed

for the surface of the 2000A film (fig. 23A).

However, upon etching down closer to the sili- con surface, evidence of some crystallization

was detected (fig. 23B). It was found that this

crystallization propagated upward toward the

surface during long periods of aging at room

temperature. Quantitative treatment of the intensity data revealed correlations between film

structure, deposition techniques, aging behavior

and electrical properties.

4.3. X-ray Texture Analysis [46-48]

Measurements of crystallographic texture, or preferred orientation, of polycrystalline metallic materials has evolved to the state of a standardized technique, in many cases automated, and the techniques are described in

detail in many texts on x-ray diffraction. As

was stated in the last section, the RHEED

pattern from a polycrystalline surface is sensitive to the presence of preferred orientation

and, indeed, if the phenomenon in question is

significant only very close to the surface,

RHEED is the best technique for the study of

texture. However, the results from such a study

are difficult to interpret quantitatively and, in

Figure 23. RHEED diffraction patterns of 2000A

thermal SiO: film grown on Si. (A) Surface of film,

(B) all but 50A of the film stripped by etching [45].

most cases, the standard x-ray methods are

superior.

To understand the principles of x-ray texture

analysis it is important to understand the

standard form for presentation of texture data,

the pole figure (fig. 24A). The "pole" referred

to in the term "pole figure" is the normal to a

particular crystallographic plane of the structure under study. The orientation of a particular crystallite can be defined in terms of the

angles that a particular pole P (say the (00.1)

pole of a hexagonal crystal) makes with the

reference axes of the ceramic surface (fig.

24B). The polar plot in figure 24A is then a

plot of the orientation of pole P. A plot of the

orientations of the (00.1) poles for each crystallite would provide statistical information

about the extent of preferred orientation in the

polycrystalline ceramic. Such a probability plot

is a pole figure. If the pole figure exhibits uniaxial symmetry about its center, the texture is

termed a "fiber texture."

287
P-0

90° 270°

180°

(A)

SOURCE f OSCILLATION DETECTOR

Figure 25. The geometry of experimental ar- rangement for an x-ray texture study is identical to that for the standard x-ray

diffractometer arrangement, with the additional feature of a sample goniometer which

allows extra degrees of rotational freedom.

Oscillation of the sample permits an in- crease in the number of crystallites ex- amined.

(B)

Figure 24. The orientation distribution of the individual crystallites which make up a polycrystalline sample is plotted on a

polar coordinate diagram (A),

the pole figure. The reference axes of the pole figure are the

normal N and the axes perpendicular to N at /3 = 0. In this case N is chosen as the normal

to the sample surface and the

/3 = 0 axis is an arbitrary di- rection in the plane of the sample. The vector P is a normal to a chosen set of planes for a particular crystallite in the sample.

Detailed information about the orientation

distribution is obtained from x-ray reflection

intensity measurements, for a given reflection,

performed on an x-ray diffractometer equipped

with a pole figure goniometer. The goniometer

allows orientation of the flat sample into positions corresponding to a, /3 ^ 0 (fig. 25). In

the case of a fiber texture, the most important

information is at « = ;8 = 0 and so a quantitative investigation of the fiber component of

the texture can be performed with a simple

diffractometer arrangement. Both types of ex- periments have been performed recently on

different types of tape cast substrates [49,50].

The results of the pole figure measurements

[50] indicated that many tape cast alumina

substrates have a very strong basal fiber texture (i.e., the (00.6) x-ray reflection was much

stronger than found for a randomly oriented

sample) . Measurements [49] of intensity ratios

1(00.6/1(11.3) substantiated these results and

also demonstrated that the texture is a sur- face-associated phenomenon. Measurements for

ground samples showed significant decreases in

basal texture. This texture behavior has also

been correlated with optical and dielectric constant observations [50].

A careful consideration of these observations

leads to several conclusions about the crystallography of ceramic surfaces. First, the apparent

association of preferred orientation with the

surface, the lack of orientation in the "green"

tapes [49] and the rotational symmetry of the

texture about the surface normal suggest that

it is the presence of the surface itself which

has caused the texture in sintered tape cast

alumina substrates. One might suspect that

small crystallites which happen to be oriented

so as to expose the low surface energy closepacked (00.1) surface will grow during the

grain growth stage of sintering at the expense

of other grains with higher surface energies.

This hypothesis leads to several other conclusions. Since surface energy is generally quite

sensitive to surface impurity level, one might

expect that the wide variation in observed textures [49, 50] from supplies to supplier is as- sociated with surface impurity variations. If

this is indeed the case, the problems associated

with correlating some performance measure

(i.e., thin film adhesion, mechanical strength)

with some material parameter (i.e., grain size,

bulk chemistry, texture) are even more difficult, for it is certain that the latter parameters

are interrelated in a very complex manner. As

an example, one might suspect that texture

would have an influence on mechanical strength,

for the internal stresses present in a sintered

body which are due to the anisotropic thermal

expansion coefficient would be relaxed if the

grains were oriented. However, mechanical

properties of ceramics are most certainly also

related to grain size, surface chemistry and

surface topography and it is difficult to imagine

how one might conduct a controlled experiment

in this situation.

5. Conclusion

This paper has described a wide variety of

experimental techniques which are useful for

288
characterizing ceramic surface properties. The

techniques were discussed in tei-rns of their

application to substrates used in thin film integrated circuits. However, these techniques do

have a more general application whenever

ceramic surface performance depends in a critical manner on the topographical, chemical or

crystallographic features of the surface.

The techniques used for characterizing sur- face topography are well known. However, it is

sometimes difficult to predict surface performance from the results of such experiments.

One such approach to the problem, in which the

surface topography is characterized directly

in terms of a performance test, was described.

The problem associated with the characterization of surface chemistry is the converse

problem. The usual approach has been to characterize surface chemistry in terms of some indirect measurement. However, there are now

available several newer analytical techniques

which provide direct information about the

chemical composition of the surface. The most

promising of these techniques are Auger electron spectroscopy and ion microprobe analysis.

There is little information available on

either the crystallography of ceramic surfaces

or subsequent effects on surface performance.

This is probably because, in most cases, sur- face crystallographic effects are masked by

more important parameters. There are many

special situations where crystallographic features do play an important role and this area

of surface characterization should be con- sidered in more detail in the future.

The authors wish to thank E. J. Sedora, who

carried out a considerable amount of the experimental work described here; J. Colby, for his

comments on the use of the electron microprobe,

SEM, and ion microprobe; Y. Nakada, who

provided us with the unpublished results of

measurements of substrate textures; and J. Drobek for his RHEED results and comments

on the RHEED technique.

6. References

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25 in Electronic Ceramics, Amer. Ceram. Soc. Spec. Publ. No. 3, Columbus, Ohio, (1969).

[2] Rigterink, M. D., Present quality of ceramic

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[3] Thomas, D. G., Ceramics and glasses—some uses

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[10] Lo, W. C, Ceramic surface texture by reflective

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[11] Sundahl, R. C, Berrin, L., and Sedora, E. D.,

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[12] Thornton, P. R., Scanning Electron Microscopy,

(Chapman and Hall, London, 1968).

[13] Hall. C. E., Introduction to Electron Mieroscony

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[16] The Electron Microprobe, Ed. T. D. McKinley,

K. F. H. Heinrich, and D. B. Wittry (Wiley,

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[17] Gray, T. J., and Zope, J. K., Electron microbeam

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[19] Colby. J. W., Quantitative microprobe analysis

of thin insulating films, p. 287 in Advances in

X-ray Analysis, Vol. 11 (Plenum Press, New

York, N. Y., 1968).

[20] Harris, L. A., Analysis of materials by electronexcited Auger electrons, J. Appl. Phys. 59,

1419 (1968).

[21] Taylor, N. J., The technique of Auger electron

spectroscopy in surface analysis, in Techniques

in Metals Research, Vol. 7, Ed. R. F. Bunshah

(Interscience Publ., New York, N. Y., 1971).

[22] Chang, C. C, Auger electron spectroscopy, in

Modern Methods of Surface Analysis, Eds. P. Mark and J. D. Levine (to be published by

North-Holland Publ. Co., Amsterdam Neth.,

1971).

[23] Palmberg, P. W., and Rhodin, T. N., Auger elec- tron spectroscopy of fee metal surfaces, J. Appl.

Phys. 39, 2425 (1968).

[24] Weber, R. E., and Johnson, A. L., Determination

of surface structures using LEED and energy

anaylsis of scattered electrons, J. Appl. Phys.

40, 314 (1968).

[25] Weber, R. E., and Peria, W. T., Use of LEED

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surface contaminants, J. Appl. Phys. 38, 4355

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[26] Palmbere. P. W., Auger electron spectroscopy

in LEED systems, pp. 29.01-11 in The Structure and Chemistry of Solid Surfaces, Ed.

G. A. Somerjai (J. Wiley and Sons, Inc., New

York, N. Y., 1969).

[27] Siegbahn, K., Nordling C, et al., ESCA (Almquist and Wiksalls AB, Upsalla, Sweden,

1967).

[28] Siegbahn, K., Nordling, C, et al., ESCA Applied

to Free Molecules (North-Holland, Amsterdam, Neth., 1970).

[29] Hercules, D. M., Anal. Chem. 42, 20 (1970).

[30] Soucha, A. J., Analysis of surfaces utilizing

sputter ion source instruments, in Modern

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[31] Anderson, C. A., Analytical methods for the ion

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[32] Anderson, C. A., Hinthrone, J. R., and Fredriksson, K., Ion microprobe analysis of lunar material from Apollo 11, pp. 159-67 in Proc. Apollo 11 Lunar Science Conf., Vol. 1, Ed.

A. A. Levinson (Pergamon Press, New York,

N. Y., 1970).

[33] Bauer, E., Low energy electron diifraction

(LEED), pp. 559-639 in Techniques in Metals

Research, Vol. 2, part 2, Ed. R. F. Bunshah

(Interscience Publ., New York, N. Y., 1969).

[34] Lander, J. J., Low energy electron diffraction

and surface structural chemistry, pp. 26-116

in Progress in Solid State Chemistry, Vol. 2, Ed. H. Reiss (Pergamon Press, New York,

N. Y., 1965).

[35] Estrup, P., Surface studies by electron diffraction, in Modern Methods of Surface Analysis,

Ed. P. Mark and J. D. Levine, to be publ. by

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30, 705 (1927).

[37] Chang, C. C, LEED studies of thin film silicon

overgrowths on ^-alumina, pp. 77.1-14 in The

Structure and Chemistry of Solid Surfaces,

Ed. G. Somerjai (J. Wiley and Sons, Inc., New York, N. Y., (1969).

[38] Charig, J. M., and Skinner. D. K., The (0001)

surface of a-alumina-LEED observations, pp. 34.1-20 in The Structure and Chemistry of

Solid Surfaces, Ed. G. Somerjai (J. Wiley and

Sons, Inc., New York, N. Y., 1969).

[39] Pinsker, Z. G., Electron Diffraction (Buttersworth, London. England, 1953).

[40] Vainshtein, B. K., Structure Analysis by Electron Diffraction (Pergamon Press, New York,

N. Y., 1964).

[41] Marcus, R. B., The characterization of solid thin

films by electron microscopy and diffraction.

Measurement Techniques for Thin Films, Eds.

B. Schwartz and N. Schwartz (The Electrochemical Society, Inc., New York, N. Y., 1967).

[42] Bauer, E., Reflection electron diffraction (RED),

pp. 501-557 in Techniques of Metals Research,

Vol. 2, Pt. 2 (J. Wiley and Sons, Inc., New

York, N. Y., 1969).

[43] Schmidt, W. A., and Davey, J. E., Preparation

of smooth, crystalline, damage-free, sapnhire

surfaces by gaseous etching, this proceedings.

[44] Bowden, F. F., and Tabor, T., p. 145 in Symposium on Properties of Metallic Surfaces (Inst,

of Metals, London, 1952).

[45] Drobek, J., Electron diffraction study of phase

transformations in grown and deposited SiO;

films, private communication, to be published.

[46] Cullity, B. D., Elements of X-ray Diffraction

(Addison-Wesley, Reading, Mass., 1959).

[47] Taylor, A., X-ray Metallography (J. Wiley and

Sons, Inc., New York, N. Y., 1961).

[48] Barrett, C. S., and Massalski, T. B., Structure of

Metals, 3rd Ed. (McGraw-Hill Book Co., New

York, N. Y., 1966).

[49] Y. Nakada, private communication, to be published.

[50] DiMarcello, F. V., Key, P. L., and Williams, J. C, Preferred orientation in AI2O3 substrates

for thin film circuits. Presented at the Fall

Meeting of the Electronics Division, Amer.

Ceramic Soc, San Francisco, Oct., (1970).

Discussion

H. J. Green : Are there any special requirements for the operation of the ion microprobe?

SUNDAHL: Usually, it is operated under a vacuum of 10 '' to 10-" mm Hg. Initially carbon

contamination occurs until the first monolayer

is sputtered away.

WiEDERHORN: What is the resolution in the

ion microprobe?

SUNDAHL: It depends on the instrument

used. Some instruments have a beam diameter

of about 300 microns. The instrument discussed

here has a beam diameter of about 2 microns.

Between 2 microns and 300 microns resolution

can be obtained depending upon how it is

focused. The resolution normal to the surface

depends on the beam current and instrument

sensitivity. Resolutions approaching one monolayer are claimed by some designers.

WiEDERHORN: How much do they cost?

Sundahl: Well, one instrument I think is

selling for one quarter of a million dollars. It's not one you go out and buy everyday. I would

suspect that there aren't too many organizations that are going to buy it. I would also

presume that it will be possible to perform

such experiments on instruments for a $1,000

a day.

Rhodes: I'd like to ask one question with

regards to the electron spectroscopy. You mentioned something about studies of segregation

of impurities at grain boundaries. Could you

elaborate on that at all and say a little bit about

the geometric resolution limits of the instrument?

SuNDAHL : The lateral resolution is poor with

a beam diameter of about a millimeter. The

way you perform a segregation study on the

grain boundary is to look at a fractured surface. That surface represents the segregation

at the grain boundary since you only are looking at the first few monolayers. A great deal

of study is being done on steels, correlating

segregation, embrittlement, and mechanical

properties. However, there has, up to this point

at least, been little done on ceramics. We hope

to do some measurements in the near future.

290
Heuer: I was interested in your application

of Auger spectroscopy to detect Ca segregation

during annealing of alumina substrates. Wasn't

there Ca segregation during firing? Was the

Ca an accidental impurity in the starting

powder?

SUNDAHL: The segregation does occur during the original firing. As the sample was an- nealed the calcium apparently diffused along

the grain boundaries up to the surface and then

diffused across the surface. It tends to segregate on the surface in this manner. The calcium

impurity level in the fired substrate was 640

ppm CaO by weight. The SiO. level was at 900

ppm and MgO at 1000 ppm. The other impurities were almost nil. The CaO impurity level

is greater than the solubility limit and is the

major source of the surface layer of Ca, although measurements on annealed sapphire

lead us to suspect that a small fraction of the

contamination could originate from the furnace.

Rice : I've seen the same sort of thing in so- called reagent grade magnesium oxide that has

been hot pressed. If you anneal it, you'll find

the small quantities of silica and other impurities coming out. In certain materials you

get, for example, anisotropic crystal structures

formed on the surface. There definitely is preferential migration to surfaces of certain

materials.

SUNDAHL : That is right. This is also true for

many semiconductors. In the study of highly

pure metals experiments have been performed

which quantitatively characterized this behavior.

Matsko: Are these Auger spectrometers

commercially available?

Sundahl: They are commercially available.

There are two companies who make Auger

spectrometers. One is PEI, Physical Electronics

Industries in Edina, Minnesota, and one is

Varian Associates in Palo Alto, California. The

particular instrument I showed is not the ideal

instrument to use. There is a better instrument

now available which has a cylindrical analyzer.

In fact, there was some work done on the anlyzer right here at NBS. The cylindrical analyzer

is a much more sensitive instrument and much

easier to use. We built this one ourselves. There

are some people, at Sandia in Albuquerque,

New Mexico, who have done quite a bit of

work in this area. R. L. Park is one such

person.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 1970, Gaithersburg, Md. (Issued May 1972).

291

Surface Characteristics of Ceramic Substrates for

Hybrid and Microwave Electronic Circuits

Jackson K. Emery ^

Accumet Engineering Corporation, Hudson, Massachusetts 07149

Several methods currently used for finishing ceramic substrates in ready-fordeposition state are discussed, including firing, glazing, burnishing, grinding, lapping,

and polishing. Each method tends to exhibit a diff"erent and typical combination of

surface characteristics, such as variations in thickness, parallelism, camber, waviness,

roughness, and localized defects; each of these in turn may influence the deposition

process, performance, and reliability. Analysis of some typical surface details, particularly roughness and waviness, as significant surface factors capable of affecting

process or function, is presented with comments on instrumentation and measurement

procedures.

Significant differences exist in the use of terms relating to substrate surfaces and

their characteristics; a need is indicated for more particularized use of such terms,

and for recognition of uniform definitions. Clarification of terminology may facilitate

improved communications and measurements, which in turn may correlate with cur- rent and future requirements in circuit production. Several existing applicable standards are referenced, and definitions of terms are suggested. A number of choices

of surface characteristics are presented which may be related with process controls

and function, but no recommendations for particular values of surface details are

intended.

Key words: Camber; ceramic; fired, flatness; ground; lapped; microwave; polished;

roughness; substrate; surface; thick film; thin film; waviness.

1. A Case for Better Substrates

Microwave and hybrid circuit utilization

promises rapid growth up to a forty-fold in- crease in ten years [1]- due to predicted ex- pansion of applications and technological

breakthroughs. To the extent that the common

ceramic substrate will remain a factor in this

growth, a close reexamination of its functional

surface characteristics may be warranted. In

fact, it is increasingly recognized that the sur- face quality of substrates is critically related

to circuit performance and reliability [2].

If we examine this rapid growth closely we

may see the beginnings of a subtechnology in

substrate surface characteristics and control.

Such an emergence could become important to

the designer, processor, and user of hybrid and

microwave circuits, and can be expected to involve disciplines usually outside the normal

bounds of electronics, to include materials finishing, microsurface analysis, and dimensional

metrology.

Many circuit designs, including vast numbers

of both hybrid and microwave circuits, have

been reliably applied on ceramic substrates

having the normal surface characteristics of

as-fired or glazed material. Until recently no

economically acceptable alternatives were available. Now, however, we can observe increasing

instances of circuits having unusually stringent

requirements, as in fineness and uniformity of

^ President, Accumet Engineering Corp.

^ Figures in brackets indicate the literature references at the end of this paper.

line or using low allowable percent spread in

circuit component performance resulting in

narrower tolerances being sought. At the same

time, low-cost production capabilities are becoming available for improving substrate sur- face quality, and better measuring equipment

and procedures can be readily acquired. When

the several surface parameters are regarded

critically in terms of new low tolerances and

corresponding functional results, the challenge

of exact surface quality in terms of its individual components, their selected emphases and

cost trade-offs becomes increasingly complex.

In this paper distinctions are proposed between the various characteristics of substrate

surface quality, and their terminology, control

and measurement are discussed.

2. Standards and Terminology

In reading the literature on substrates one

may conclude that some inconsistencies and

omissions exist in describing concepts and in

understanding and in definitions of terms. To

encourage uniformity, attention may be directed to applicable existing standards, which

include:

MII^I-lOA, 1962

MIL-S-55620 (EL), 1969

ACMA, Standards for High Alumina Ceramics, 1969

ANSI-B46.1, Texture, 1962

ANSI-B87.1, Decimal Inch

Various commercial standards by ceramic

makers, processors, and users.

293

439-921 O - 72 - 20
Some evidence may be noted regarding a need

for improved and more detailed standards and

terminology for better understanding. A number of definitions of terms are listed at the end

of this paper including several less familiar

terms for possible consideration.

One important exception to standards adherence appears to be justified in the choice of

the most effective symbol denoting microinches.

Although ANSI Standard B87, Decimal Inch,

1967, recommends discontinuation of the

familiar term ^in in favor of the term mk, the use of mk in the publications and communications of the electronic industry is virtually

nonexistent at this time. In this paper the

term /xin v^^ill be used to avoid misunderstanding.

3. Units of Measurement

Although SI (metric) units are now officially

specified for the National Bureau of Standards (NBS) publications, "customary" units

(inches) are currently used almost exclusively

for designation of lengths, widths, diameters,

thicknesses, feature locations, and other details

of substrates in most current catalogs, brochures, and drawings issued by substrate

manufacturers, processors, or users in the

United States of America. Even in brochures

of importers substrate dimensions are published

in inches for size, thickness, camber, glazed

thickness, and finish; various strengths and

thermal conductivities are usually shown in

"customary" units, not in SI units. One major

producer states, "All dimensions are in inches."

The ACMA standard gives "customary" units

in their table 1, but also gives "metric" units

in table 2,

In view of the almost nonexistent usage of

SI units in this context, they will not be used

in this presentation to avoid confusion and

difficulty.

4. Dimensions

The dimensions affecting surface features of

substrates are principally thickness, parallelism, and localized defects of various kinds.

Various standard dimensions are offered by

substrate manufacturer, chiefly for length,

width, diameter, and thickness in a variety of

decimal and fractional inch sizes. The variety

of sizes and the frequent use of special sizes

having odd dimensions reduces the occurrences

of like dimensions to the point that very few of

the many sizes are available for immediate delivery. The beginning of a less expensive and

immediately available substrate market may

depend on standardization of sizes in suitable

steps, but in a limited number of varieties ; possibility such sizes could be selected from existing lists, but bear an interrelationship based

on a preferred number series, or other appropriate guideline. By standardization through re- duced variety of sizes gains in ecomomy and

availability may justify increased attention to

evaluation and improvement in substrate sur- face characteristics.

5. Tolerances

Substrate surface quality relates to tolerances of thickness, parallelism, flatness, camber, and texture (roughness and waviness).

Some of these are included in the definitions at

the end of this paper. Tolerances are not uniformly standard, but for principle dimensions

of length and width (not precisely surfacerelated) the major manufacturers specify plus or minus 1 percent, although some users may

specify plus or minus 0.5 percent, or plus 0

minus 0.5 percent. Four Thickness tolerances are fairly common : plus or minus 0.001 in, plus or minus 0.002 in, plus or minus 0.003 in and

plus or minus 0.004 in for thicknesses 0.010 in,

0.020 in, 0.025 in and 0.050 in. Lapped or polished substrates may have thickness tolerances

of 0.0005 in to 0.0002 in or less. Camber tolerances appear to vary widely according to tradition or application for many thousands of an

inch per inch to extreme cases of 0.0001 in per

inch. Flatness tolerances may be specified separately from camber for critical applications.

Parallelism tolerance, when required, is usually

less than thickness tolerance. Waviness tolerance is almost never defined or expressed; although waviness may be critical functionally in

some applications, it is usually ignored. Roughness tolerance is most frequently expressed as a

single value representing the allowable high

limit, particularly for very fine fired, glazed, or

polished substrates in the order of 10 /xin or less.

However, a growing utilization of a roughness

tolerance concept having a high and low limit

appears in hybrid circuit substrates where the

variation in printed components such as resistors must be extremely small; typically, such

roughness tolerance may be 24 /xin AA plus or minus 1.5 fxin AA, particularly for precision

lapped substrates.

6. Surface Structure and Porosity

Surface quality is dependent on both density

and uniformity of the substrate. Nonuniformity in ceramics, such as shown in figure 1, may

penetrate deeply and may often run the entire

substrate thickness. Some porosity, however

fine, random or discontinuous, is inherent in

most ceramic structures and can be a cause of

holes, pits, pocks, blisters, and blebs (see MILS-55620-EL). Isolated pores may occur just

under the surface or be intersected by the

294
Figure 1. Non-homogenous structure in 1 in x 2 in

substrate of 99.5 percent alumina; lighter areas show

about 1 fiin roughness, while darker areas are 6 to 12 fiin, pattern is full thickness of 0.0B5 substrate.

finished plane of a fine polished substrate.

Exposing such pores results in inherent structural voids, often mistakingly thought to be

"pullouts" resulting from improper grinding or other surface treatment; "pullouts" can be

correctly identified only when they exhibit evidence of "tears" or "comet-tails." Surface discontinuities can more or less unfavorably

influence the superimposed circuit, depending

on the circuit location or vulnerability and the

pore dimension. When porosity is poor it not

only may effect function and reliability, but can

cause deceptively high roughness measurements [3]. Even in fine quality 99.5 percent

alumina individual pores commonly measure

from 0.0001 in to 0.0005 in width and can be

occasionally much larger and closely clustered.

According to MIL-1-lOA, porosity shall be

determined with dye penetration in accordance

with ASTM publication D-116-44; this test

may be most useful for inter-connecting porosity. However, since porosity is frequently discontinuous, the importance of isolated pores

may not be easily determined by the ASTM

test. In such cases substrate surfaces should be

visually examined in accordance with MIL-S55620 (EL).

Remarkable details of fine grade as-fired

Alumina (99.5%) may be observed at high

magnification by use of a scanning electron

microscope; figures 2 and 3 show large differ- ences in grain size, and spaces between grains

Figure 2. As-fired 99.5 percent alumina, scanning

electron microscope, (SEM), lOOOx.

Figure 3. As-fired 99.5 percent alumina, scanning

electron microscope. 10,000x.

which may suggest significant subsurface porosity ; also of possible interest is the improvement shown at similar magnifications after

precision polishing to 2 fxin AA, as shown in

figures 4 and 5.

295
7. Flatness

It is widely recognized that controlled flat- ness of a substrate, along with controlled sur- face finish is a significant factor in circuit performance and reliability [2]. Flatness error may be considered to include waviness, but can

refer to only one side of a substrate at a time.

Flatness is not strictly equivalent to camber,

since the flatness errors of both sides become

components of camber along with variations of

thickness, parallelism, and local defects.

Flatness correction, especially through precision lapping and polishing, offers a direct and

economical means for eliminating waviness and

minimizing camber in precision circuit work.

Substrates finished by a lapping operation make

possible a rapid and inexpensive flatness well

within a most rigorous requirement, without

the necessity of either substantial stock re- moval or prior grinding.

8. Waviness

The familiar succession random ripples,

irregularities, and undulations across a surface

is called waviness ; the usual spacing from crest

to crest is considerably greater than that of

the spacing of the roughness crests, but less

than the surface width. Waviness in one component of texture, according to ANSI B 46.1,

and is generally characteristic and attributable

to the surface forming process. Inherent in both

Figure 4. Precision polished substrate from same lot of 99.5 percent alumina as in fig. 2 and 3. lOOOx,

showing open pocks (SEM)

Figure 5. Precision polished substrate from same lot of

99.5 percent alumina as in fig. 2 and 3, 10,000x,

showing fine line scratches from diamond polishing,

most principal lines less than 0.00001 in wide, open

pocks from inherent pores show intact wall detail, no "comet-tails" (SEM).

glazed and as-fired ceramic, waviness may be

regular as may occur in surfaces produced with

a rigid grinding wheel, or may be extremely

random in both vertical heights and horizontal spacings characteristic of firing processes.

Actual height variations due to waviness can

be measured by sliding a substrate across a

small flat anvil with a spherically tipped

spindle lightly touching the upper surface.

Waviness can also be determined on a conventional stylus type instrument (see figure 6)

with a strip chart added. Another measurement

Figure 6. Standard stylus type roughness measuring

instrument (Courtesy Industrial Metrology Div.,

Bendix, Ann Arbor, Mich.)

296
method for both waviness and flatness utilizes a Moire pattern which can yield a direct visual

profile of the entire substrate area through a

range of variations from 0.010 in to 0.0001 in or less (see figures 7 and 8).

9. Camber

As commonly defined, camber is a threedimensional concept and is a composite dimension encompassing the combined profiles of

both surfaces of a substrate, including variations of flatness, vi^aviness, thickness, parallelism, local raised zones (burrs, blebs, ridges)

and other defects. Generally accepted by a go/no go "slot gage" inspection, a substrate

vi^hich is just under the high limit of the thick- ness tolerance will be rejected for negligible

camber, while substantial camber near the low

thickness limit will be accepted. An alternative

method for measuring camber uses a simple

dial indicator gage and measures camber as the

maximum indicator difference between measurement of substrate height with convex side

up and a measurement with convex side down.

Camber of substrate 0.040 in or less in thickness

can also be measured as the indicator gage

movement when the substrate is placed on the

anvil with convex side up and then sprung down

into contact with the anvil. Camber components

may be assessed visually and numerically even

on dull finished substrates, produced by any

manufacturing method including firing, grinding, matte-lapping, or polishing, by observing

the Moire band pattern formed on a fine-line

grid (see figures 7 and 8). Polished substrate

flatness can also be measured by observing

light interference fringes on an optical flat

Figure 7. Drawing of typical

Moire line pattern on a 1000-line

grid formed by a 2 in X 2 in

as-fired substrate showing combined flatness and waviness of

about O.OOJ4. in. (Each line cor- responds to a contour interval of

about 0.0005 in.)

Figure 8. Drawing of typical

Moire line pattern on a 1000-line

grid formed by a 2 in X 2 in

precision polished substrate

showing absence of waviness and

overall flatness of about 0.0002

in.

where the manner of reading is identical to

that of the Moire band, except the fringe unit

of curvature is about 0.00001 in (one-half the

wave length of visible light).

10. Thickness and Parallelism

Thickness variations from substrate to substrate are undesirable particularly when screen- ing because of the effect on screen clearance

or contact and consequent effect on the pattern.

Some screening departments find advantageous

a narrow control of thickness to as little as

0.001 in total variation, provided camber and

waviness are also minimized.

Parallelism is usually regarded as simply variation of thickness across a particular substrate. When parallelism error is in the form of a wedge it may cause screening problems. Both

parallelism and thickness variations between

substrates can cause problems in automated

processing. For measurement of thickness and

parallelism the instrument contacts should have

only a small bearing on the substrate, for ex- ample from 0.08 in to 0.125 in to avoid measuring local defects or camber.

11. Roughness

Roughness measurement has been extensively

analyzed for several decades and has given rise

to authoritative standards. Roughness is described in several ceramic references as microinches CLA. However, the term in jxin AA con- forms to ANSI Standard B 46.1. Actually CLA

(centerline average) and AA (arithmetic aver- age) are essentially equal. Too frequently we

still see the term RMS (root mean square aver- age) which was discontinued as a roughness

297
standard in 1955 in favor of AA; RMS values

are about 10 percent greater than AA values.

The standard method of B 46.1 for evaluation of roughness is the diamond stylus tracing

technique, as illustrated in figure 6. In this

method a polished diamond stylus tip, preferably having a 0.0001 in radius, is made to tra- verse the specimen in a straight line under a

light load (1 gram), thus generating a signal

which is proportional to the roughness or texture of the specimen. The accuracy of the re- sult is dependent upon reasonable care and

calibration of the electronics and the stylus of

the instrument; under reasonably controlled

conditions agreement between instruments and

operators should be within 20 percent variation,

and with care and accurate calibration can be

within 10 percent or less. The standard roughness value is read as an amplified signal from

either a meter or a strip chart in units of /xin

AA, according to ANSI B 46.1. It may be of

interest, especially to those who apply very

fine circuit conductors and components that the

true peak-to-valley height on the substrate is

from 4 to 5 times greater than the AA value.

It has been observed [2] that surfaces having

very different real profiles may produce similar AA roughness readings; this problem can

be safely neglected only when surface textures

produced by similar finishing processes are compared with correct instrumentation. Roughness measurements by air probes usually include a waviness component [4] which may be

misleading. Light-section microscopes can be

useful by measuring individual scratches or

flaws, but are less useful for characteristic

roughness especially when AA values are low.

An interference microscope will show visually as a topographic fringe pattern the true peakto-valley height of roughness detail, across a

small area, preferably for finishes of 20 /xin or

less; usually the field is very small, typically

from about 0.08 in (60x) to 0.01 in (480x).

Ordinary band patterns from light interference

or from fineline grids will disclose flatness but

not roughness detail.

Standard roughness readings are usually

traced at selected locations, as within the four

quadrants of the ceramic, to establish the variation of the roughness. It should be remembered that stylus techniques merely permit

sampling of the surface as very narrow traces,

which is usually very rapid, consistent, and

adequate.

12. As-Fired Substrate Surfaces

Surfaces of as-fired substrates may be slightly more dense than the interior by as much as

0.5 percent [2]. These surfaces will exhibit at

low oblique viewing a medium gloss and rippled appearance, causing reflected images of

straight lines to be wavy. Ordinarily the roughness of 99.5 percent ceramics will be relatively

fine, perhaps 10 mk AA or less, while that of

96 percent or 94 percent Alumina will be

coarser perhaps 20 mk AA or more. In as-fired

substrates, camber is usually present, often to a degree discernible visually and over a relatively broad range, usually being greater on

thin or large sized boards. Surface defects inherent with the production process, such as

burrs, pits, pocks, and fins or ridges, either occur in profusion or be negligible, according

to the quality of the process. Tape-process substrates frequently show low corners or raised zones which can result in thickness variations across the substrate up to 0.002 in or more.

13. Glazed Substrate Surfaces

Glazed substrates have surface characteristics very similar to those of as-fired substrates,

except for the very thin smooth glassy surface

produced by refiring, often in the order of 1

^in AA roughness. Glazed surfaces show a tendency for a raised zone near the edges. Glazing

is usually on one side only. Defects of waviness,

flatness, camber, and variations of thickness are similar to those of as-fired ceramics.

14. Burnished (Tumbled) Substrate

Surfaces

Substrates having sufficient as-fired defects

to render them unsuited for coating are sometimes burnished by various tumbling or vibratory processes, utilizing balls, stones, sand, or abrasive grains. Following an extended cycle

of many hours, or days, the substrates exhibit

slightly rounded corners or edges, smoothed

chips, and reduced burrs. Little change of waviness occurs, although a narrow spread in roughness may be possible [5]. Generally camber (or

flatness) and variations in thickness are not

altered l3y the process. Burnishing tends to con- centrate on sharp aspects while it changes the

broad surface very little. Removal of burrs by

burnishing may improve handling and metallization quality.

15. Ground Substrate Surfaces

Grinding a substrate by use of a rigid abrasive wheel removes the irregular surface layers

and the characteristic ripple of the as-fired condition. Camber may be improved by precision

grinding, or it may be made worse by altering

stress balance by incorrect techniques. Ground

finishes may contain relatively deep and ob- scure scratches from a raised grain on the

grinding wheel or they may show surface voids

as grain pullouts due to improper grinding conditions. Very fine diamond wheels correctly ap298
plied can produce finishes of fair uniformity to

8 fiin or less.

16. Lapped Substrate Surfaces

Lapping as a finishing process for ceramic

substrates consists of moving a tool or lap

against the substrate in the presence of fine

abrasive grains either as a rigid layer impregnated into the lap as a thin slurry. Typically

the lap material is softer than the work material, and the grains are harder. The pressure

between the lap and the substrate is controlled

from about 0.5 to 10 psi to suit the conditions.

As-fired substrates should be lapped directly

without grinding. A proper choice of lapping

conditions can rapidly and economically convert

irregular ridges, waves, peaks and valleys into

a wavefree uniform surface with minimized

camber, little stock removal, and reduced variation of thickness. A number of options exist

for control of thickness and parallelism. Surface roughness can be narrowly limited at a

desired level from 50 to less than 1 /xin AA

across one surface, or from substrate to substrate. Preferences vary from finishing both

sides 2 /xin AA maximum, to 30 plus or minus

2 ^in AA on one side only.

Four different typical standard grades for

lapped hybrid substrates are as follows:

Grade 1. Front only: no waves, reduced

camber, controlled finish

Grade 2. Both sides: no waves, reduced

camber, controlled finish

Grade 3. Both sides; reduced camber, controlled finish;

Front only: no waves; back only:

waves 50 percent removed: Narrow variation for thickness

Grade 4. Both sides: no waves, reduced

camber, controlled finish: Narrow

variation for thickness and parallelism.

Lapping can be effective for correction or

control of surface characteristics on almost

any substrate material, purity, shape, size, or thickness from 0.075 in or more to 0.005 in or

less. Most substrates can be successfully production lapped in ordinary or unusual configurations with holes, recesses or scored grids.

Flatness of lapped ceramics may be typically

0.0002 in ; the flatness departure from a perfect

plane can be determined by noting band curvature in the Moire pattern formed against a finelined grid, as in figure 8. A visual indication

of the change in surface character from as-fired

to precision lapped or polished may be noted

by comparing details observed in a scanning

electron microscope at medium to high magnification, as shown in figures 2, 3, 4, and 5. Note

the complete absense of the tears or "comettails" characteristic of "pullout."

Possible advantages for use of lapped substrates for hybrid circuits include improved

handling and screening, extended screen life,

reduction of rejects, minimized resistor trimming, increased capacitor yield and easier as- sembly.

17. Polished Substrate Surfaces

Precision polishing is an exactly controlled

process and may be defined simply as an ex- tremely refined lapping technique, utilizing the

finest lap materials and abrasives, which often are especially compounded diamond powders.

Polished substrates which exhibit a bright reflective appearance but with poor flatness, camber, and varying smoothness, have limited

value. Precision polished substrates are characterized by superior optical grade flatness, no

"pullout", negligible camber, and low roughness

at a selected level from 10 to 1 /xin, held to a

narrow spread across the entire surface and

from piece to piece and from lot to lot. The

finest polishes are made in 99.5 percent alumina, but lower purity and lower density materials can also be polished when required. Surface flatness of most precision polished ceramics

can be controlled to less than 0.0001 in; such

flatness levels are easily evaluated by observing

light interference fringes on an optical flat

where a regular pattern will be seen similar

to that of figure 9.

18. Thick Film Substrate Surfaces

Fabricators have experienced some disappointments in achieving narrow tolerances for

printed resistors [6] . It seems clear that a maFlGURE 9. Unretouched Polaroid photo of polished 2 in

X 2 in alumina substrate showing light interference

fringes for an overall flatness error less than 0.00005

in (Note each line here corresponds to a contour

interval of half the wavelength of light, or about

0.000012 in.)

299
jor improvement in the spread of such resistor

values depends not only throug-h control of process variables but through minimizing of substrate variables, especially those of camber, flatness, and roughness [6]. Flatness [1] can be

critical in automated thick film printing, especially with fine mesh screen, very narrow lines

(0.002 in or less) and up to five or more se- quentially printed layers [7]. Some processors

have already reported up to 30 percent increase

in yield by printing resistors on uniformly

lapped substrates.

19. Thin Film Substrate Surfaces

Thin film circuits, which may employ 94

percent to 99.5 percent purity substrates, are

subject to failure [8] or poor performance in

the event of local overheating due to excessively

weak spots in the film, or open circuits and inconsistent resistivity due to discontinuities in

the film; such defects can arise from burrs,

pits, or local undulations typical of asfired ce- ramics. Since precision lapped or polished substrates are virtually free from these defects,

many substrates for thin film circuits are now

lapped.

20. Microwave Substrate Surfaces

Most users of lapped alumina substrates for

microwave applications request precision polishing in the very low roughness range from

about 10 to 1 fiin, chiefly below 5 /xin, although

the practical limit may currently be about 0.3

/xin AA. At this time the above range of polished surface quality is manufactured routinely

in production lots primarily in 99.5 percent

purity only. Most frequently both sides are

specified for the identical finish; alternatively,

one side (sometimes called "front") will require

typically a 2 /xin finish while the opposite side

("back") will require in one case a 10 /xin finish,

while in another case a 30 /xin finish, and in

still another case the as-fired finish is retained.

Microwave substrates are often specified with

low flatness (0.0001 in to 0.0005 in), low camber (0.0002 in to 0.001 in), and zero waviness.

Comparable surface quality, with certain specific limitations, may be applicable to materials

other than alumina, including beryllia, ferrite,

and others. For example, most beryllia can be

polished to 3 /xin or less only under very favorable conditions and with relative difficulty,

while high quality ferrite can be polished to 1

/xin and to optical flatness of 0.000001 in to

-0.000010 in.

2 1 . Conclusions

Adequate recognition of the complexity and

importance of substrate surface quality requires that distinctions be made between the

several separate but interrelated surface characteristics. Proper specification of significant

substrate surface characteristics, their dimensions and tolerances, may be enhanced by cor- relation of functional effects of each surface

detail. Attractive economics can be effected

through cost tradeoffs based on functionally

determined specifications from the most stringent to the most relaxed levels. Additional eco- nomics may be possible indirectly through increased implementation of and familiarity with

improved measuring methods, new finishing

techniques, updated standards, and exact terminology.

22. Definitions

The following list of terms excludes those covered

elsewhere in this paper or in aoplicable references

such as the Standard MIL-S-55620 (EL).

Centerline: a line parallel to the general direction of

the surface profile such that the sums of those areas between it and those parts of the profile on either

side are equal.

Lay: direction of a predominant surface pattern, usually designated not numerically but by orientation, as

longitudinal, transverse, diagonal, curved, or random.

Matte dullness of a surface; absence of reflectivity.

(Note: a matte surface may have extremely low

roughness.)

Polish: reflectivity of a surface.

Profile: an intersection, disclosing detailed texture of

a surface, between the surface and a perpendicular

plane oriented in a partcular direction.

Roughness: a rating of the profile deviations as an

arithmetic average (AA), measured in microinches

normal to the centerline.

Texture: the repetive or random deviations from the

nominal surface which form the pattern of the surface, consisting of roughness, waviness, lay, and flaws.

Waviness: the more widely-spaced irregularities or undulations of surface texture, the spacing of which

is greater than that of individual roughness peaks but

less than the surface width.

(Note: roughness may be considered to be superimposed

on waviness.)

Also see ANSI Standard B 46.1, Surface Texture.

23. References

[1] Hybrid Microelectronics Review, July, 1970, Matthew F. Romano, Editor.

[2] Szekely, George S., Alumina substrates reliability

criteria, a user viewpoint, Nepcon West, Feb.

11, 1969.

[3] Lenz, Robert G., GMC research laboratories, private correspondence, May 11, (1970).

[4] Pifer, B., and Collis, W. P., Special Report No.

476, American Machinist, August 10, 1959 (Surface Finish Measurement).

[.5] Indars, G., I. G. Engineering Co. private communication, July, (1970).

[6] Van Hise, Jon A., Process variables in thick film

resistor fabrication, Solid State Technology,

July, 1970.

[7] Sweet, Alvan F., Insulation forum. Insulation/

Circuits, March, (1970).

[8] Groves, H. T., and Keister, F. Z., Know your thin

film materials. Electronic Design 19, Sept. 12,

(1968).

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2^, 1970, Gaithersburg, Md. (Issued May 1972).
Ceramic Surface Texture by Reflective Replica Technique

Wing C. Lo

Bell Telephone Laboratories, Inc., Allentown, Pennsylvania 18103

Surface texture is a parameter of general concern in the study of thin film metallizing adherence on ceramic substrates. It refers to the geometrical character of the

surface irregularities recurring many times across the ceramic surface tending to form

a pattern. In this context, surface texture is determined by the size, shape, arrangement and distribution of the surface constituents.

This note presents a convenient method by which the surface texture of ceramic

substrates can be visually assessed and recorded photographically using a simple

metallographic microscope.

Examples in applying this technique for distinguishing and differentiating surface

textures of different types of ceramic substrates are given. Use of the technique for

identifying surface defects is also shown.

The simplicity of the method should warrant its use for monitoring the physical

surface quality of ceramic substrates.

Key words: Alumina substrates; surface replica; surface texture.

1 . Introduction

Surface texture refers to the geometrical

character of the surface irregularities recur- ring many times across the ceramic surface

tending to form a pattern.^ It is determined by

the size, shape, arrangement and distribution

of the surface constituents. Surface texture

here does not refer to the crystallographic orientation of a material.

The surface texture of ceramic substrates

used in thin film applications is a parameter of

general interest. Theoretically, the breaking

load of a metal lead bonded to a thin film metallized substrate or the electrical properties of

the thin film on the substrate can be influenced

by the surface texture.^ Thus, the uniformity

and homogeneity of the surface texture of substrates is one of the requisites for maintaining

the quality in a finished product. The need for

assessing the surface texture is apparent.

One aspect of the surface texture that warrants consideration in thin film applications is

the area of the surface accessible to and actually covered by the thin film material. This

physical quality of the surface cannot be

assessed numerically in terms of a root-meansquare value, a peak-to-valley value or the commonly used center-line-average (c.l.a.) value

obtained with a mechanical profile tracer. The

fact that c.l.a. values are not functionally dependent on or related to the real surface area

has already been pointed out by J. E. Marian.^

A qualitative, and possibly quantitative evaluation, of this physical quality can be easily

achieved by visual observations with the aid of

an optical microscope or a scanning electron

microscope. (SEM).

^ Reason, R. E., The Measurement of Surface Texture, (Cleaver- Hume Press Ltd., Great Britain, 1960). ^-A paper on the subject is in preparation by the author.

' Symposium on Properties of Surfaces, ASTM Special Publication No. 340.

However, with an ordinary metallographic

optical microscope, one cannot view the surface

of the ceramic directly. This is because of the

small amount of light reflected from the textured surface. It is even more difficult to record

the observed image photographically. The object of this note is to present a convenient technique by which these difficulties may be over- come. This is the reflective metallographic

replica technique which has been in use by

metallographers for some time.

The advantages of this method are its simplicity and general availability of the equipment required. One does not have to have

tedious preparations of the sample before it

can be studied. In the case of the SEM, it is

necessary to coat the ceramic surface with a conductive film before the surface can be

viewed. The limitation in the depth of focus or penetrating power attainable with optical

microscope objectives is not always a disadvantage in this application. With the optical microscope one can estimate the depth of surface

defects by focusing the top and bottom planes of

the defect and measuring the difference in these

levels. This estimation of depth is more difficult

with the SEM. To assess the homogeneity of the

surface texture one would like to view a large

enough area of the sample at a glance. This can

be accomplished at a magnification of about

200X. At this magnification the field of view is

several orders of magnitude greater than the

basic pattern of the surface texture. It is high

enough to reveal the geometrical features that are of interest. With the SEM, useful magnifications starts at about 500X. Of course, when

the asperities on the surface are in the neighborhood of 1 /xm or less, then one has to resort

to the use of the SEM which has substantially

better resolution ('-'lOOA).

301
2. Description of Method

The material used in this technique comes

from a kit marketed by Struers Scientific Instruments, Copenhagen, Denmark, under the

trade name "Transcopy Reflective Metallographic Replica". It can be purchased through

W. J. Hacker & Co., Inc., West Caldwell, New

Jersey 07006.

The kit consists of a bottle of Transcopy

liquid and a box of Transcopy-replicas. The

Transcopy-replica is essentially a bright aluminum sheet coated with a green colored plastic

film on one side.

The procedure for preparing a Transcopyreplica specimen is shown in figure 1. The first

step is to place a small drop of the Transcopy

liquid on the surface of the cleaned ceramic

sample (fig. la). A small piece of the Transcopy replica is immediately placed over the

liquid with the green plastic side facing the

sample surface (fig. lb). Hold the replica

against the surface with uniform pressure ' for

approximately two minutes. Place a double

sided adhesive tape on the back side of the

replica (fig. Ic). Attach a glass slide onto the

adhesive tape (fig. Id). Turn the assembly over

and remove the ceramic sample (fig. le). The

mounted replica (fig. If) is ready for viewing

under the metallographic microscope. The green

color of the film serves as a filter for viewing.

From experience better replicas have been

obtained with ceramic samples by substituting

the Transcopy liquid with a similar replicating

solution marketed by Ladd Research Industries,

Inc., Burlington, Vermont.

It should be noted that the replica is not useful for low magnifications (below lOOX) as the

influence from the reflective aluminum layer

will become too great with the increased depth

of focus.

3. Results

The inadequacy of cla values for describing

the surface texture of ceramics is illustrated

in figure 2. The c.l.a. values cannot relay the

difference between a ground and as-fired surface. Figure 3 demonstrates the vastly different

surface textures of 99.5 percent alumina ce- ramics from different suppliers that meet the

same 25 /xin (0.64/xm) min c.l.a. requirement.

At low magnifications (150X to 200X) the

limitations of the usefulness of the scanning

electron microscope is illustrated in figure 4. The texture of the two substrates can be easily

differentiated visually from the photographs

obtained by the replica technique. Their difference is less apparent when one compares the

images recorded by the scanning electron microscope.

* A pressure of 80 g/mm ^ is adequate.

-TRANSCOPY UOUID ^ ^ ^ X y X CERAMIC

GREEN SIDE OF REPLICA

(a) APPLYING TRANSCOPY LIQUID

.-TRANSCOPY - REPLICA

CERAMIC .......... ^ ^TRA

(b) ATTACHING TRANSCOPY REPLICA (ALLOW

REPLICA TO SET FOR - 2 MINUTES) DOUBLE SIOEO__—* V ' ' J ' ^ '- ' -REPLICA

AOHESIVETAPE /

^ ,,,,,,

(c) APPLYING DOUBLE SIDED ADHESIVE TAPE

GLASS

SLIDE

(d) ATTACHING GLASS SLIDE

CERAMIC

r/ //' /// ///' ^'M//

(e) REMOVING CERAMIC

r //' /y. //' y/y^

(f) MOUNTED REPLICA

Figure 1. Procedure for preparing a transcopy

replica specimen.

Use of the replica technique in revealing vast

differences in surface textures of ceramic which

one may get from suppliers is shown in figures

5 through 7. The difference in surface textures

of ceramics may be due to differences in body

composition and/or in manufacturing methods.

Figure 5 compares the surface textures of 99.5

percent alumina ceramics reported to have the

same body composition by the supplier but

made by different processes, viz., the drypressed process and the tape-cast process. Fig- ure 6 illustrates the difference in surface texture of 99.5 percent alumina substrates of the

same body designation, made by the tape-cast

process but from different lots shipped by the

supplier. Figure 7 shows the textures of a drypressed 94 percent alumina ceramics in different shipments from the same supplier. According to the supplier, the difference here is due

to a change in the finishing operations of the

ceramics from grinding to lapping.

Changes in surface texture can also be

brought about by post-manufacturing processes.

Figure 8 shows the change in surface texture

of a 99.5 percent alumina ceramic by refiring

it to cone 31. The change in a 94 percent alumina ceramic surface texture effected by an acid

etchant (HF and HNO.^ solutions) is shown in

figure 9.

Applications of the replica technique in identifying surface defects and damages are shown

in figures 10 and 11. Figure 10a is a replication of the surface defect that results in a

302
32 40

41 40 35

332/2425

35

29

CENTER LINE AVG. (MICRO-INCHES)

EACH VERTICAL DIV. = lOO^^in EACH HORIZONTAL DIV. - .002 in CUT-OFF WAVELENGTH = .01 NOTE: l>iin=0.0254>im

TALYSURF MEASUREMENT

Figure 2. Surface texture of ceramics by replica technique vs center line average values.

white area on a substrate when inspected with

illumination from its back side. This defect is

a pit on the back side of the substrate directly

below the location of the white area. When il- luminated from the back side, the location of

this defect on the viewer side of the substrate

will look white due to greater amount of light

that can get through to the viewer at this location. The presence of relatively large grains in

the pit suggests the presence of impurities

which caused exaggerated grain growth or the

formation of new phases. This could establish

the basis for further investigation for remedial

actions by the manufacturer. Figure 10b is the

replica of a fissure on the substrate surface

which shows up as an area that retains fluorescent penetrants. The fissure is possibly caused

by impurities imbedded in the surface. Figure

11 shows pull-outs and streaks of distorted

material from grinding operations. The same

grinding process produces more pits or pullouts

on the 94 percent than on the 99.5 percent

alumina ceramics.

4. Conclusions

Although the surface texture of ceramic substrates used in thin film applications may not

be the only parameter that governs the desired properties of the finished product, it is a

variable that should be controlled.

The replica technique offers a simple and

practical aid for ceramic suppliers to monitor

their manufacturing processes and for users to

check on the uniformity in the quality of the

parts they receive.

The author wishes to thank Messrs. J. E.

Clark, W. B. Grupen, F. L. Rowland and P. I. Slick for their encouragement and comments.

303
m

tar- • ; •

. - *» ^ *

» -I »

a CERAMIC BODY A ; SURFACE AS-FIRED

b CERAMIC BODY A ; SURFACE TUMBL,ED

c CERAMIC BODY A ; SURFACE LAPPED

d CERAMIC BODY B ; SURFACE GROUND

Figure 3. Surface textures of 99.5 percent alumina ceramics meeting 25

fjLin (0.64. mm) CLA finish.

SUBSTRATE 1 SUBSTRATE 2

a,b OPTICAL MICROSCOPE USING REPLICA TECHNIQUE

c,d ELECTRON SCANNING MICROSCOPE

Figure 4. Comparison of low magnification micrographs of surface

textures taken with optical microscope using replica technique and

with electron scanning microscope.

304
^ih>^»lfil'^i^^lif^^ ' .o^

Figure 7. Surface textures of

ground dry-pressed 9i. percent

alumina ceramics from different

lots shipped by same supplier.

Figure 6. Surface textures of tape-cast 99.5 percent alumina substrates

from different lots shipped by same supplier.

305
tOOMm

b

a AS-FIRED

b REFIRED

100 um

Figure 8. Change in surface texture of

99.5 percent alumina ceramics by

rejiring.

100 urn

a AS-RECEIVED

b ACID ETCHED

Figure 9. Change in surface texture of

94. percent alumina ceramics by acid

etching.

306
o PIT b CRACK

Figure 10. Defects on ceramic

substrate surfaces.

99.5% ALUMINA 94% ALUMINA

4

AS-FIRED

Figure 11. Grinding damages on 99.5 percent and 94. percent alumina

ceramic surfaces.

307
NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machimng and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).
Quantitative Surface Finish Characterization by CESEMI

E. W. White and H. A. McKinstry

Materials Research Laboratory, Pennsylvania State University,

University Park, Pennsylvania 16802

and

A. Diness

Office of Naval Research, Metallurgy Program

Arlington, Virginia 22217

A feasibility for quantitative characterization of surface finishes based on the

computer evaluation of scanning electron microscope images (CESEMI) has been

established. Measurement of the number, size, shape and orientation of isointensity

(or selected intensity interval) regions constitute the basic analysis. Brightness in SEM images varies as a function of the steepness of slope of the surface at a given

point. In their analysis, the characteristics of grooves, pits, ridges, etc., can be studied

separately. Detailed interpretation of the results are not attempted in this preliminary

study.

A feasibility for recording and analyzing profilometer rasters has also been

demonstrated. Such recordings are, in essence, true three dimensional images of

surface topography insofar as the profilometer stylus faithfully follows the surface

morphology. Analysis of these profilometer recordings by the basic CESEMI computer

programs yield direct descriptions of the topography. The SEM image and profilometer

"image" analyses appear to be complementary techniques. One obvious advantage of

the SEM image analysis is that it is a no-contact technique hence, there are no

problems introduced by stylus damage.

Key words: Ceramic Materials; CESEMI (Computer evaluation of scanning electron

microscope images); digital magnetic tape recordings; profilometer; profilometer

rasters; quantitative characterization; scanning electron microscope; surface finish

analysis; surface morphology; surface roughness; surface topography; three dimensional images.

1. Introduction

Quantitative characterization of surface finish parameters should be useful in evaluation

of the mechanisms of material removal by various finishing processes as well as in assessing

the relation of final surface finish characteristics to property.

The analysis of ceramic surface finish is the

problem of describing quantitatively the important morphological characteristics of the

surface topography which result from a particular finishing process such as grinding, burnishing, sputtering, polishing, etc. Parameters

of interest include: surface roughness; microfracture characteristics; size, shape and depth

of pits, etc. Also of importance are the orientation of elongate features such as scratches,

grooves, ridges and fractures.

It is proposed that much of this information

might be gleaned from computer analysis of

scanning electron microscope images. As a complementary approach, the analysis of rastered

profilometer traces (successive line traces over

a given area) may be used to get a reading of

the complete three-dimensional surface topography.

Instrumentation and computer programs developed for our research on particulate characterization appear to be directly applicable

to studies of ceramic surface finish. This paper

is intended as a feasibility study to set forth

possible approaches and how they might be implemented. To achieve this goal the basis for

image formation in the SEM will be reviewed.

Results from the analysis of five representative

finishes will be discussed. The results from a

trial profilometer reading are also presented.

2. Description of SEM Image

Forming Process

The SEM employs a finely focused (100-

200A diameter) beam of high energy electrons

(5 keV —V 30 keV) to dislodge low energy (1-

50 eV) electrons from the specimen surface. A

highly sensitive electron scintillation detector

monitors the currents (10"" —> 10"^ amps) of

the emitted secondary electrons. The detector

signal is used to moderate the brightness of a cathode ray tube (C.R.T.) as both the electron

beam on the specimen inside the SEM and the

C.R.T. spot are rastered synchronously. This

results in the generation of an image of the

specimen surface viewed on the C.R.T.

The SEM gives quite high resolution (200A)

images of the specimen surface at a depth of

field some 300 times greater than for normal

309

439-921 0 - 72 - 21
light microscopy at a given value of magnification. Upon hitting the specimen surface the

electron beam penetrates to a considerable

depth (up to several microns depending on the

keV and specimen composition). Throughout

the entire volume of beam penetration, lov^ en- ergy secondary electrons are generated but only

those generated vi^ithin 10-50A of a surface can

escape and become part of the image signal.

Figure 1 shows schematically the sort of

intensity of secondary electron emission to be

expected from a surface consisting of familiar

geometric configurations. From examination

of figure 1 it follov7S that:

(a) deep holes and cracks yield essentially

no secondary electron signal.

(b) the steeper the slope of a surface, the

higher will be the secondary electron

emission.

(c) sharp ridges resulting from the intersection of two steep slopes will yield

the highest signal. Any obtuse angle

of intersection (peak or ridge) will

yield brighter emission than either of

the intersecting surfaces.

(d) notches or grooves resulting from the

acute angle of intersection of two sur- faces will yield intensities lower than

from a surface which is normal to the

beam.

Not all the emitted secondaries may be collected by the detector. The fraction detected

depends upon (1) the degree of collimation or angular selectivity of the detector and (2)

reabsorption of the secondaries by prominent

features on the specimen surface. These losses

introduce shadow effects in the image as

illustrated in figure 2 for an extreme case. Most

observed images exhibit shadow effects somewhere between these two extremes.

Figure 2. Idealized secondary electron images with

exaggerated shadow effect from selected geometric

features.

3. SEM Instrumentation

The SEM and related recording instrumentation used in our study are shown schematically in figure 3. The x-ray channels are not

used in this study but could have been. All

SEM images were recorded at 300X magnification. Intensities were measured on a gray scale

of 2000. There were 250 raster lines with 250

sample points per line giving sample grid of

62,500 data points. The detector was optirnized

to suppress shadow effects, thus, image brightSIGNAL M

U ANALOG

PREPROCESSOR

LT

TO

DIGITAL 1 P

LE CONVERT

X.

,,

BUFFER

CONTROLLER

Figure 1. Idealized secondary electron image from

selected geometric features.

Figure 3. Schematic block diagram for the

SEM and, instrumentation for digital magnetic tape recording.

310
ness for this study closely follows the model

shown in figure 1. Additional details of the instrumentation and

the basic computer programs have been published previously [1, 2, 3, 4]^

4t Profilometer Recordings of Surface

Topography

The SEM images give the impression of a

three dimensional picture but since the intensities are a complex function of several parameters, the actual surface topography is not

measured. A novel technique was developed to

obtain the surface topology directly. Within the

range of its intended use (resolution etc.,) the

signal output from a profilometer is a direct

function of the surface height. Thus, by ar- ranging the profilometer to raster the sample

surface, a three dimensional representation of

the surface is obtained. The recording system

and computer programs developed for the

measurement of SEM images are suitable for

handling this profilometer input. Multiple con- tours thus made give a topographical map of

the material's surface. Our test runs are rather

crude in that a makeshift "raster" was utilized,

however, they serve to establish the general

feasibility.

5. Computer Analysis

The basic computer programs calculate the

characteristic features contained above (or

below) a selected intensity level in the image.

Each continguous field is recognized regardless

of its shape complexity. Basic parameters calculated for each such field are its area and

perimeter. An ellipse is fitted to the perimeter

of each field and its characteristics are calculated including the lengths of major and minor

axes and the orientation of the major axis. (Although the axial orientation is used in the program it is not a part of the output generated

in this study). From these basic parameters

other characteristics such as aspect ratio, ratio

of the measured perimeter to ellipse perimeter

and etc., are calculated as desired.

6. Results

Five specimens from 48 samples provided

by R. J. Stokes of Honeywell and Roy Rice of

NRL were chosen to represent a variety of

surface finishes, materials and finishing techniques. Secondary electron images were recorded photographically and simultaneously on

digital magnetic tape. In all cases the secondary

electron detector was set to minimize shadow

effects.

• Figures in brackets indicate the literature references at the end of this paper.

In addition to the conventional SEM image,

a y-axis modulation photograph was obtained.

These displays tend to emphasize topographic

features. Unfortunately there is an electronic

distortion in these particular images so that

corresponding features on the two photos do

not register exactly in the verticle direction.

Computer generated binary area and perimeter maps are illustrated in figures 4 and 5. These are photographic reductions of several

pages of computer print. Due to the higher

density of printed characters per line than the

carriage spacing in the print, these representations have a rectangular distortion. Although

they are not actually used in the computer analysis, such printouts serve to illustrate certain

functions of the programs.

Figures 4 through 8 show the images on

which computations have been performed. All

results are in the basic program units and not

true dimensions. One data point corresponds

to 1.33/xm or an area of 1.77 V^^.

The frequency distribution of the contiguous

area sizes at two different levels are plotted

in figures 9 and 10 suggesting the kind of additional information that can be obtained from

the computer analysis. The SEM pictures which

show small features have a similar distribution

for both peaks and pits. The materials with a

smoother appearance have a broad size distribution as one might expect.

The profilometer trace analysis shown in fig- ure 11, indicates the kind of result that one

would expect. The contour was drawn at one

level and shows the nearly random pattern that

should result from a surface of novaculite

(crypto-crystalline quartz) ground flat on a

grinding wheel using a 600 mesh silicon car- bide abrasive.

This work is part of a program being sponsored by the Office of Naval Research, Metallurgy Division under Contract N00014-67-A0385-0007.

Most of the samples used in the study were

supplied by Roy Rice of the Naval Research

Laboratory and R. J. Stokes of Honeywell. Profilometer instrumentation was made available

by the department of Industrial Engineering

and Professor A. B. Draper and Mr. J. M.

Samuels, Jr., of that department have assisted

by way of several useful discussions.

Other personnel collaborating in this program include: Dr. G. G. Johnson, Jr., Dr. H.

Gorz, Mr. and Mrs. J. Lebiedzik, Miss K. Mayberry and Mr. R. Hoover.

7. References

[1] McMillan, R. E., Johnson, G. G., Jr., and White,

E. W., Computer Processing of Binary Maps of

SEM Images. Proceedings of the 2nd Annual

311
Scanning Electron Microscope Symposium,

IITRI, pp 439-444 (1969).

[2] White, E. W., Gorz, H., Johnson, G. G. Jr., and

McMillan, R. E., Particle Size Distributions of

Particulate Aluminas from Computer Processed

SEM Images. Scanning Electron Microscopy/

1970. Proceedings of the 3rd Annual Scanning

Electron Microscope Symposium, IITRI, pp

57-64, Chicago, Illinois (1970).

[3] Matson, W. H., McKinstry, H. A., Johnson, G. G.,

Jr., White, E. W., and McMillan, R. E., Computer Processing of SEM Images by Contour

Analysis, (in press—Pattern Recognition).

[4] White, E. W., Mayberry, K., and Johnson, G. G.,

Jr., Computer Analysis of Multi-channel SEM

and X-ray Images from Fine Particles. Technical

Report No. 3, Office of Naval Research, Metallurgy Program, Contract No. N00014-67-A0385-0007, December 1970.

Si

Figure 4. Lucalox gas polished to AA finish of 3.2fim. Sample supplied

by R. Rice.

A. Secondary electron images with detector conditions set to minimize

shadowing effects. Regions of steepest slope have greatest brightness. B. Y-axis modulation of same area as A. An instrumental distortion

has elongated this image toward the top so that the registration is not accurate with respect to A.

C. Computer generated binary map at the fifty per cent level. All

data points whose intensity is greater than the average value are

indicated by a black dot. D. Computer generated perimeter at the 95% level. The computer

printer introduces an elongation in the top-bottom direction. The

bright areas in A can be matched to their corresponding configurations. E. Computer generated perimeter map at the 50% level. F. Computer generated perimeter map at the 10% level (inverted).

312
A lOOU B

Figure 5. Silicon face (100) ground by an alumina

wheel to an AA finish of 0.25fim. Sample supplied by

R. J. Stokes.

A. Secondary electron image showing oriented sur- face features.

B. Y-axis modulation of same area as A.

C. Computer generated binary map at 10% level (inverted).

D. Computer generated binary map at 95% level.

Figure 6. Alumina, diamond polished to an AA finish of 0.32ij,m. The

sample was obtained from R, Rice.

313
A 100 u B

Figure 7. Sapphire surface after R.F. sputtering.

The AA finish is 2.0^m.

Figure 8. Magnesia surface ground by a diamond wheel to an AA sur

face finish of 1.1 fim. Sample obtained from R. J. Stokes.

Figure 9. Frequency Distribution of Contiguous Areas at the 95% level.

314
300-

FiGURE 10. Frequency Distribution of Contiguous Areas at the 10% level.

315
Figure 11. Profilometer image of Novaculite surface.

The rastered profilometer image contoured at a

50 percent level by computer processing. Multiple

contouring would give a topographic map of the

surface.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

316
An Assessment of Surface and Subsurface Damage Introduced in Ceramics by

Semifinish Grinding Operations

B. G. Koepke

Honeywell Corporate Research Center, Hopldns, Minnesota 55343

The nature and extent of grinding damage introduced by surface grinding a number of ceramics having widely varying mechanical properties has been studied.

Surface damage has been characterized by optical and electron microscopy and subsurface damage has been observed using etch pit techniques. Grinding damage was

found to depend strongly on wheel type and rate of material removal as well as on

the mechanical behavior and microstructure of the workplace. The results indicate that

grinding damage is mainly composed of mixtures of three types depending on the

mechanism of material removal.

When material is removed efficiently from low impact resistance ceramics (e.g.,

magnesium oxide and ferrite) the surfaces are generated by brittle fracture and are composed of regions of transgranular and intergranular fracture. When grinding is inefficient; i.e., when the grinding wheel loads up, material is removed by plastic flow.

The resultant surfaces on deformable ceramics (e.g., magnesium oxide, ferrite, and

silicon) are smooth and burnished, but may contain thermal cracks due to the heat

generated. In this instance subsurface damage consists of a discrete, highly deformed

layer containing cracks in most cases. When material is removed efficiently from high impact resistance, nondeformable

ceramics (e.g., alumina and boron carbide) material is removed by plastic flow and by

transgranular and intergranular fracture. The presence of extensive plastically flowed

regions on the ground surfaces of extremely hard ceramics is surprising and points

out the extremes of stress and temperature existing under the wheel-workpiece interface during a grinding pass. Key words: Ceramics; etch pits; fracture; grinding damage; plastic deformation;

surface condition ; surface grinding.

1 . Introduction

Ceramics are finding- increased use in modern

technology because of their wide range of physical properties. Optimization of the physical

properties is generally attained by close control

of primary fabrication variables. For many applications extremely close dimensional tolerances are called for. Precise ceramic machining

is therefore usually performed as a secondary

fabrication operation on material in the as-fired

(fully hardened) condition.

To date the most common ceramic machining techniques are those that use abrasives of

one form or another for grinding, lapping and

polishing operations. Although abrasive machining techniques are highly developed they

are quite operator dependent and surprisingly

little is know^n about either the nature of as- machined surfaces or the material removal

phenomena leading to the finished surface

Even less is known about the condition of the

material immediately beneath the finished surface. These points are of concern because it is

known that many physical properties of ce- ramics, particularly mechanical properties, are

highly structure sensitive and the successful

operation of a precision ceramic device might

ultimately depend on its surface condition [2].

Abrasive machining generally involves a se- quence of steps; namely, rough grinding (hog-

' Figures in brackets indicate the literature references at the end

of this paper.

ging) in which the workpiece is cut to shape,

semifinish grinding in which the final shape

is obtained, and finally finish grinding, lapping

and polishing in which the desired surface finish and tolerances are obtained. From the

standpoint of surface integrity, a successful

machining operation is one in which the damage introduced during a given step is completely removed by the following step. The most

likely place for the introduction of persistent

damage is during the rough and semifinish

grinding operations. In many instances this

damage is not satisfactorily removed by the

later machining steps. For this reason it is felt

that an understanding of the nature and extent

of the damage introduced in ceramics by semifinish grinding is essential if optimum workpiece quality is to be achieved.

This paper presents an assessment of the

state of ceramic materials subjected to semifinish grinding. It will compare the damage re- sulting from semifinish surface grinding operations on diff'erent ceramics having widely

varying mechanical response to the high

stresses imposed by the machining operation.

The paper will distinguish between the surface

condition, referred to as surface damage, and

the subsurface damage. The various types of

grinding damage encountered will be discussed

in terms of the properties of the workpiece, the

characteristics of the grinding system and the

tool-workpiece interaction.

317
2. Choice of Materials

The materials used for this study were chosen

primarily on the basis of their mechanical behavior. Ceramics may be divided according to

three categories of mechanical behavior;

namely, completely brittle, semibrittle and ductile [3, 4]. In completely brittle ceramics dislocations cannot move and plastic shear by slip

is not possible even at stresses approaching

theoretical strength. In semibrittle ceramics

dislocations are mobile but slip is restricted to a limited number of slip planes. In ductile

ceramics dislocations are extremely flexible and

cross slip over many planes results in complete

plasticity.

The categories of mechanical behavior of the

ceramics used in this investigation are shov^n

in table 1. As table 1 indicates, the mechanical

behavior of these materials depends on temperature with a trend towards increased deformability at higher temperature. Differences

in mechanical behavior can be expressed by the

number of independent slip systems operating

at different temperatures. According to the von

Mises criterion a crystal needs five independent

slip systems to undergo a general change in

shape. Note in table 1 that this condition only

prevails in magnesium oxide, spinel and silicon

at high temperatures. Alumina does not achieve

the necessary number of independent slip systems. To the author's knowledge the slip systems in boron carbide have not been determined. However, since the material has a hexagonal crystal structure, its mechanical behavior is presumably similar to alumina.

3. Experimental Procedure

3.1. Materials

Samples of the ceramic materials indicated

in table 1 were obtained in both single crystal

and polycrystalline form from various sources

as indicated in table 2. These materials are

listed according to their hardness and it should

be noted that a hardness differential is observed

even among different samples of the same materials; e.g., polycrystalline alumina.

Prior to machining studies, samples were

prepared having approximate dimensions 1/2

in X 1/4 in X 1/8 in. In the case of single crystals of magnesium oxide these could be prepared

by cleavage. Silicon single crystals and polycrystalline samples were all prepared by meTable 1. Mechanical behavior of ceramics used for this study

Material Temperature Number of

independent

slip systems

Mechanical

behavior

Reference

Magnesium oxide (single crystal and polycry- R T 2 Semibrittle 3, 6

stalline) >156o'°C 5 ductile 3, 5

Ferrite [mechanical properties assumed similar R.T. 5 Semibrittle 6, 20

to spinel (MgALOi] or brittle

>1600 °C 5 Semibrittle

Silicon (single crystal) R.T. 5 Brittle 7

>650 °C 5 Semibrittle

Alumina (single crystal and polycrystalline) R T 2 Brittle 3, 8

>900 °C 2 Semibrittle

Boron carbide (polycrystalline) >1400 "C 9 Brittle 9

Table 2. Grain sizes and hardnesses of ceramics u^ed for this study

Material Supplier Grain size (jum) Hardness (Knoop)

MgO single crystal Norton 336

Sintered MgO IITRI 18 411

Hot pressed NiZn ferrite Honeywell 3 670

Si single crystal Mallinckrodt 950

Sintered (Lucalox) AhOi General Electric 38 1800

Hot pressed AI2O3 AVCO 2 2000

Cold pressed and sintered (AD999) AI2O3 Coors 2 2100

Hot pressed B4C AVCO 8 2800

318
0 1 '

1

Figure 1. Photomicrographs of the cutting faces of

(a) a 4.6 grit vitrified bond alumina wheel and

(b) a 100 grit, resinoid bond diamond wheel.

chanical sawing and polishing to shape. The

magnesium oxide crystals were annealed 1 hour

at 1200 °C in air in order to pin all existing

dislocations and to place the material in a uniform starting condition with respect to mechanical behavior [10].

3.2. Grinding Parameters

Samples were ground on a DoAll Model

D618-7 high speed precision surface grinder.

All grinding experiments were performed in

the down (climb) grinding mode with the cutting face of the wheel moving in the same direction as the workpiece. Up grinding was not

studied since some materials shattered. This

problem was particularly severe with magnesium oxide crystals.

The wheel type is an important parameter in

any grinding operation. Two wheels have been

used in these studies; a 46 grit vitrified bond

alumina wheel (Norton type AA-46-H8V40)

and a 100 grit resinoid bond man made diamond wheel (DoAll type DIAI-MDIOONIOOB1/4B3). The alumina wheel is 17.8 cm in diameter and 12.7 mm wide and the diamond

wheel is 15.2 cm in diameter and 3.18 mm wide.

The wheels have significantly different structures as shown in figure 1. The structure of the

alumina wheel is relatively open and many

voids exist in the cutting surface. The cutting

face on the diamond wheel, on the other hand,

is relatively closed and consists of diamond

abrasive points protruding from a continuous

matrix.

Ground ceramic surfaces were characterized

as a function of the machining parameters,

wheel depth of cut (d) and workpiece feed

rate (F). The spindle speed was fixed at 3540

rpm. Grinding was carried out both dry and

wet using water containing water soluble oil

in the ratio of 80:1 as a coolant.

The alumina wheel was dressed before each

surface was ground. The diamond wheel, on the

other hand, was dressed only at the beginning

of the program. It was observed that this wheel

did not readily load up during grinding. When

grinding debris did appear on the wheel it was

removed by wiping with a rag moistened in

methanol.

Magnesium oxide, ferrite, and silicon were

ground with both the alumina and diamond

wheels. The harder ceramics were only machined with the diamond wheel.

The rate of material removal is a convenient

parameter describing the severity of the grinding operation and is defined as the product

(dVb) where b is the width of the area ground

during each pass. Feed rates ranged from

0.042 to 4.2 cm/s (1 to 100 in/min) and depths

of cut ranged from 12.7 to 50S^m (0.0005 to

0.020 in). Typical rates of material removal

employed with each wheel are listed below.

Wheel b(mm) V (cm/sec) Rate of Material Removal

(cm'/sec)

46 grit

alumina

12.7 25.4

25.4

50.8

0.042

0,42

4.2

1.36 X 10 ' — low

1.36 X 10-' — moderate

2.72 X 10-' — high

100 grit

diamond

3.18 25.4

25.4

50.8

0.042

0.42

4.2

0.34 X 10-' — low

3.4 X 10-' — moderate

6.8 X 10-' — high

319
4. Grinding Damage in Semibrittle

Ceramics

4.1. Magnesium Oxide Single Crystals

As might be expected single crystal semibrittle ceramics undergo plastic deformation

during grinding. However, it is found that the

extent of deformation depends on the ability of

the material to respond to the change in shape

imposed by the grinding wheel. For this reason, it is necessary to distinguish between

magnesium oxide crystals most favorably oriented for shear; i.e., those ground on {100}

planes in <100> directions (see figure 2a) and

other orientations less favorably oriented for

shear (see figure 2b). Further distinctions will

be made between subsurface and surface damage.

a. [100] Surfaces Ground in <100> Directions

Previously published results [11] have shown

that the nature and extent of grinding damage

in magnesium oxide crystals ground on {100}

surfaces in <100> directions is a strong function of both wheel type and rate of material

[ooi]

CLEAVAGE

Q FACES

Figure 2. Schematics showing the orientations in

which magnesium oxide crystals were ground.

(a) {100} <100> orientation {100} cleavage

faces and {110} slip planes favorably oriented

for shear are indicated, (b) {110} <110> and

{120} <120> orientations.

removal. Coolant was also noted to have an

effect at high rates of material removal.

Subsurface damage was studied on {100}

planes cleaved transverse to the ground surface

and etched to reveal dislocations. Surface damage was characterized using conventional optical and replica electron microscopy techniques. The results are summarized below.

(a) Subsurface Damage. In all cases it was

found that surface grinding resulted in the

formation of a discrete, highly deformed layer

immediately adjacent to the ground surface

that could be observed as a dense band of dislocation etch pits. Examples of subsurface

grinding damage in crystals ground wet at low

rates of material removal on {100} surfaces in

<100> directions with the alumina and diamond wheels are shown in figure 3. The un- eveness of the damage layer in the crystal

ground with the diamond wheel is attributed

to the uneven nature of the impact of the protruding abrasive points on the cutting surface

a ' . . • , 100 ym

Figure 3. Photomicrographs showing subsurface

grinding damage on etched {100} cleavage

faces of magnesium oxide crystals ground wet

on {100} surfaces in <ilOO^ directions, (a) 46

grit alumina wheel, d = 25.4^to, V = 0.055

cm/s. (b) 100 grit diamond wheel, d = 25.4/tm,

V = 0.042 cm/s.

320
of the diamond wheel. Cracks extending into

the subsurface deformed layer were never observed in crystals ground in the {100} <100>

orientation with either wheel.

The depth of the damage layer increased

with material removal rate when the alumina

Figure 4. Photomicrographs of {100} surfaces of

magnesium oxide crystals ground wet in a <C100> direction with a i.6 grit alumina wheel

at a low and high rate of material removal.

(a) Unetched, d = 25.4/<.m, V = 0.021 cm/s. (b)

Unetched, d = 254Mm, V = 0.34 cm/s. (c) etched, d = 254aito, V = 0.34 cm/s.

wheel was used. Measurements of the average

depth of the layer in crystals ground wet with

the alumina wheel ranged from about 30/xm

at low rates of material removal to 100/i,m at

high rates of material removal. In crystals

ground dry with the alumina wheel the depth

of damage ranged from about 30/xm at low

rates of material removal to 240ju.m at high

rates of material removal. When the diamond

wheel was used, the depth of damage remained

relatively independent of grinding condition

also converging to 30/i,m at low rates of material removal.

(b) Surface Damage. Crystals ground wet

with the alumina wheel exhibited two surface

characteristics depending on the rate of material removal. At low rates, (d = 25.4;u.m, V =

0.042 cm/sec), material was removed by brittle fracture resulting in a rough cleavage

faceted surface (fig. 4a). When the rate ex- ceeded that at which swarf could be expelled by

the wheel, the voids in the open alumina wheel

(fig. 1) loaded up. This prevented fresh cutting

points on the wheel from impacting the surface

of the workpiece. The frictional component of

the grinding forces and the temperature at the

surface increased to the point where material

was removed by plastic flow. The resulting sur- face became extremely smooth and burnished

as shown in the electron micrograph in figure

4b. This type of surface was observed at even

the lowest rates of material removal when

crystals were ground without coolant using the

alumina wheel.

The regular array of fine cracks present on

the surface, shown in figure 4b, resulted from

thermal stresses set up when the hot, as-ground

surface was quenched following passage of the

wheel. Etching enhanced the appearance of the

cracks as shown in figure 4c. The thermal

cracks were quite shallow and were never

found to extend into the subsurface damage

layer. The shallow depth of thermal cracking

is a strong indication of the steepness of the

stress and temperature gradients existing un- der the wheel-workpiece interface during grinding.

The diamond wheel did not load up wet or

dry, at any rate of material removal. Grinding

with the diamond wheel always resulted in the

typically rough surface shown in figure 5. The

white areas on the micrograph are {100} cleavage facets lying parallel to the ground surface.

b. {100} Surface Ground in Other Directions

The effect of crystal orientation on the na- ture and extent of grinding damage in magnesium oxide crystals is being examined to help

relate single crystal observations to machined

polycrystalline ceramics. Also, as mentioned

earlier, changing the orientation of single crystals leads to configurations less favorably or321
GRINDING

100

Figure 5. Photomicrograph of the as-ground {100}

surface of a magnesium oxide crystal ground wet

with a 100 grit diamdond wheel in a -ClOO^ direction, (d = 12.7/iTO, V = 0.42 cm/s).

iented for plastic flow. In this series of experiments {100) crystal surfaces (i.e., the same sur- face as in section 4.1.a) were ground in <110>

and <120> directions. In this case, all grinding was carried out wet. Again, the results

with both alumina and diamond wheels were

compared ; however, grinding with the alumina

wheel caused the specimens to shatter at all

but the lowest rates of material removal.

(a) Subsurface Damage. The subsurface

damage in crystals ground on {100} surfaces in

directions other than <100> was essentially

the same both in nature and extent to that

observed on {100} surfaces ground in <100>

directions. That is to say a discrete plastically

deformed surface layer extending approximately 30 microns beneath the surface was

GRINDING

100 m

Figure 7. Etched {100} cleavage faces showing subsurface grinding damage in magnesium oxide

crystals ground wet on [ilO] surfaces in <ill0y

directions, (d = 25.4/tm, V = 0.042 cm/s). (a) 46

grit alumina wheel, (b) 100 grit diamond wheel.

[ 100 ]

GRINDING

DIRECTION

[ 010 ]

Figure 6. As-ground {100} surface of a magnesium oxide crystal ground wet

at a low rate of material removal in a <C110~> direction with a k6 grit

alumina wheel {d = 25.4/iW, V = .042 cm/sec).

322
observed with the alumina wheel at low rates

of material removal and with the diamond

wheel at all rates of material removal. The

depth of damage of crystals ground with the

diamond wheel appeared to be relatively insensitive to both rate of material removal and to

the direction of grinding.

(b) Surface Damage. The surfaces of these

crystals ground with the alumina wheel at low

rates of material removal were generally rough

and contained large amounts of cleavage fracture. A typical example of a crystal ground on

a {100} surface in a <110> direction is shovni

in figure 6. It is interesting to note the strong

{100} fracture morphology on the surface.

At intermediate and high rates of material

removal with the alumina wheel crystals maAlumina 25.4 0.042

Alumina 25.4 0.21

Alumina 127.0 0.042

c. {110} Surfaces Ground in <110> Directions and

{120} Surfaces Ground in <120> Direction

The next step in this study of the effects of

orientation on grinding damage in magnesium

oxide crystals involved a change to different

crystal faces. Specifically grinding was performed on the {110} and {120} planes in

<110> and <120> directions respectively.

Again, all grinding was carried out wet. Specimens with these surfaces were carefully prepared beforehand by grinding, polishing and

annealing.

As figure 2b shows, {100} cleavage surfaces

exist perpendicular to the ground surface and

parallel to the grinding direction for both

orientations. After grinding, specimens were

cleaved on this plane and etched to determine

the nature of the subsurface damage.

In agreement with the experience of section

4.1.b crystals could be successfully ground with

the alumina wheel only at the lowest rates of

material removal (d = 25.4/Am, V = 0.042

cm/sc). On the other hand, the diamond wheel

could be used at all rates of material removal

without shattering the crystals.

(a) Subsurface Damage. The subsurface deformation in these crystals was quite similar in

nature and extent to crystals ground in other

orientations. Again, a discrete, heavily deformed layer existed under the ground surface.

chined in both <110> and <120> directions

shattered. Nevertheless examination of the

ground areas of the shattered remnants always

showed them to be burnished. Thus, for these

orientations, as soon as the alumina wheel

loaded and the grinding frictional forces increased, the crystals were not able to respond

to the necessary change in shape by plastic

deformation. Instead they shattered.

It is interesting to note that the transition

to shattering was more dependent on the feed

rate than on the depth of cut. This is evident

from the observations listed below, where it can

be seen that a five-fold increase in material

removal rate (proportional to dV) resulted in

shattering when it was achieved by increasing

the feed rate but not when it was achieved by

increasing the depth of cut.

Grinding Crystal

direction shattered

1.07 X 10-* <110> No

<120> No

5.33 X 10-* <110> Yes

<120> Yes

5.33 X 10-* <110> No

<120> No

Typically, this layer was uniform in crystals

ground with the alumina wheel and uneven in

crystals ground with the diamond wheel.

The most striking difference from previous

observations, however, was the extent of subsurface cracking. Damage layers in crystals

ground on {100} and {110} surfaces were

routinely found to contain extensive arrays of

subsurface cracks. This is shown for crystals

ground at a low rate of material removal in the

{110} <110> orientation with the alumina and

diamond wheels in figure 7. The cracks lie predominantly along {100} traces in both micrographs and extend to the bottom of the damaged layer in the crystal ground with the diamond wheel. Also note the tendency for the

cracks in samples ground on {110} faces with

the diamond wheel to slant preferentially in

the direction that the wheel travels over the

surface. Presumably this is due to the nature

of the stress state at the wheel-workpiece interface.

Crystals ground on {120} faces also exhibited

extensive subsurface cracking. The slip and

crack geometry with respect to the ground sur- face in these crystals was more complicated as evident in figure 8. It is an interesting point that the damage

layers in crystals ground in the {110} <110>

and {120} <120> orientations with the diaWheel d(u.m) V(cni/sec) dV(cm^/sec)

323
Figure 8. Etched {100} cleavage face showing subsurface grinding damage in a magnesium, oxide

crystal ground wet on a {120} surface in a <^120^ direction with a 100 grit diamond wheel.

(d = 25.4Mm, y = 4.2 cm/s).

mond wheel using heavy cuts (50.8 and 254/i,m)

contained considerably fewer and shallower

cracks than samples given 25.4/xm cuts.

(b) Surface Damage. {110} and {120} sur- faces ground with the alumina wheel at low

rates of material removal were rough and produced by multiple cleavage faceting. At higher

rates the specimens shattered, but the surfaces

of the fragments indicated that burnishing had

occurred prior to shattering.

{110} and {120} surfaces ground with the

diamond wheel were typically rough and similar to that shown in figure 4 with the exception

that {100} cleavage faces parallel to the ground

surface were, of course, not observed.

4.2. Magnesium Oxide and Ferrite

Polycrystals

The effects of material removal rate and

wheel type on the nature of the surfaces of

polycrystalline semibrittle ceramics ground

using coolant have been examined for magnesium oxide and hot pressed Ni-Zn ferrite

(see table 1). As expected from the results described above on magnesium oxide crystals, the

alumina wheel could only be used successfully

at low rates of material removal (d = 25.4/xm,

V = 0.042 cm/sec). At higher rates of material

removal the samples shattered. These materials, on the other hand, could be ground quite

successfully at all rates of material removal

with the diamond wheel.

a. Subsurface Damage

No measurements or observations have yet

been made of the subsurface damage introduced by grinding these polycrystalline materials. The depth of the subsurface damage in

polycrystalline magnesium oxide machined

with a diamond saw has been studied by Evans

and Davidge [12] who noted that the maximum depth of flaw (cracks or grain pullouts)

penetration depended on grain size. The depth

of damage was observed not to exceed two

grain diameters in material having grain sizes

less than about 100ju,m and did not exceed one

grain diameter in larger grain size material.

Whether the machining was carried out wet or dry was not brought out in the paper.

b. Surface Damage

Examination of the ground surfaces of fragments from shattered specimens of polycrystalline magnesium oxide and ferrite ground wet

with the alumina wheel indicated that extensive burnishing and thermal cracking had oc- curred in every case. This observation agrees

qualitatively with those made on single crystals.

The as-ground surfaces of magnesium oxide

polycrystals ground wet at a number of material removal rates with the diamond wheel

showed little or no evidence of plastic flow or burnishing in agreement with the observations

made on magnesium oxide single crystals. The

major features of as-ground surfaces of semibrittle ceramics ground wet with a diamond

wheel are illustrated in figure 9 which is an

electron micrograph of a replica of the ground

surface of a magnesium oxide polycrystal.

Three types of surface feature are evident in

figure 9, exposed grain boundaries, exposed

grain boundary porosity, and brittle transgranular fracture. The top and bottom grains

at the left side of the figure exhibit brittle

transgranular fracture while the rest of the

figure is composed of exposed grain boundaries

resulting from grain pullout. An interesting

feature of this micrograph is the slip lines in

the grains exposed by the grinding operation.

The presence of slip lines indicates the regions

Figure 9. Electron micrograph of a replica

of the surface of polycrystalline magnesium oxide ground 50.8fxm in successive

12.7/um cuts wet with a 100 grit diamond

wheel.

324
immediately adjacent to the machined surface

of a semibrittle ceramic are plastically deformed. Similar features were observed on

replicas of ground ferrite surfaces.

With regard to the effect of the severity of

the machining treatment on the nature of

ground surfaces of magnesium oxide and ferrite polycrystals, it was observed that an increased rate of material removal resulted in

more grain pullout. Grain size can also affect

the amount of pullout. In their observations on

diamond sawed magnesium oxide polycrystals,

Evans and Davide [12] observed that large

grain samples exhibited less pullout than small

grain material.

5. Grinding Damage in Brittle

Ceramics

5.1. Silicon Single Crystals

In these experiments silicon crystals were

ground on {111} faces in <110> directions

with the 46 grit alumina and 100 grit diamond

wheels at various rates of material removal.

All grinding was carried out wet.

a. Subsurface Damage

In order to assess the nature and extent of

grinding damage beneath the ground surface

of silicon crystals an attempt was made to

detect subsurface deformation by the etch pit

technique. Following grinding, the crystals

were cleaved on the {111} plane intersecting

the {111} ground surface along a line lying

parallel to the grinding direction. The cleavage

faces were then etched using the Dash etch

[13] and examined for evidence of dislocation

etch pits. A number of attempts failed to re- veal any dislocation etch pits that could positively be identified with the ground surface and

it was concluded that the deformation, if any,

produced by surface grinding either does not

extend to measurable depths beneath the

ground surface or cannot be detected on fill}

cleavage faces by the etch pit technique. Stickler and Booker [14] have shown conclusively by

transmission electron microscopy that discrete

dislocation arrays can be introduced as far as

14/x.m below the {111) surface of silicon crystals

by abrading with No. 240 SiC paper. These

arrays are associated with individual cracks

and cannot be detected by etch pit techniques

because of their close (lOOA) spacing. Thus, it appears likely that there is a subsurface damage layer in silicon, but its extent is considerably less than observed in magnesium oxide

crystals.

b. Surface Damage

Silicon crystals ground wet with the alumina

and diamond wheels exhibited quite different

Figure 10. Photomicrographs of {111} surfaces of

silicon single crystals ground in <^110^ directions

wet with a id grit alumina wheel at two rates of

material removal, (a) d = 25 A/xm, V = 0.042 cm/s

(b) d = 2hAixm, V = 0.42 cm/s.

surfaces. {111} surfaces ground with the alumina wheel were burnished as shown in figure

10. The light regions on the micrographs are

islands of burnished material. Comparison of

the two micrographs in figure 10 shows that

increasing the rate of material removal results

in an increase in burnished expected.

Grinding with the diamond wheel at all rates

of material removal resulted in rough surfaces

produced by brittle, cleavage or conchoidal

fracture. Burnishing was not observed.

5.2. Alumina Polycrystals

In these experiments three kinds of polycrystalline alumina (see table 2) have been

ground with the 100 grit diamond wheel. (Naturally, grinding with an alumina wheel was

impossible.) Grinding was carried out both wet

and dry. The ground surfaces were characterized by optical microscopy and by replica electron microscopy.

325

439-921 O - 72 - 22
a. Subsurface Damage

No attempt was made to characterize the

depth of damage in polycrystalline alumina

during this study. Hockey [15], however, has

examined single and polycrystalline alumina

mechanically polished through 0.25/xm diamond

by transmission electron microscopy and has

noted that the damage extended less than 3/j,m

below the polished surface. Furthermore, the

damage consisted of a shallow, highly deformed

surface layer containing dense arrays of dislocation half-loops associated with scratches

left by the abrasive particles.

b. Surface Damage

A most surprising feature of polycrystalline

alumina surfaces ground with a diamond wheel

was that they routinely contained regions of

burnished material. Figures 11, 12, and 13

show optical and electron micrographs of the

surfaces of Lucalox, AVCO, and Coors alumina

respectively ground dry using a 25.4;um cut at a feed rate of 0.42 cm/sec.

The light areas on the optical micrographs

are burnished regions that correspond to the

regions on the electron micrographs containing

closely spaced parallel grooves. This has been

observed by others [16, 17] and appears to be

a common feature of alumina surfaces ground

with a diamond wheel.

The dark areas on the optical micrographs

correspond to regions in which material has

been removed by brittle transgranular or intergranular fracture (pullout). These regions are

self-evident on the electron micrographs.

There appears to be more extensive burnishing

in the AVCO and Coors aluminas than in

Lucalox. This could be related to the small

grain size (2ju,m) of these materials as compared with the Lucalox (38/x,m).

The surfaces of alumina samples ground wet

contained features identical to those ground

dry (i.e., burnishing, transgranular and intergranular fracture). It was generally noted,

however, that the amount of burnishing decreased when a coolant was used.

5.3. Boron Carbide Polycrystals

a. Subsurface Damage

No attempt was made to characterize the

depth of damage in polycrystalline boron

carbide.

b. Surface Damage

Optical and electron micrographs of the

surface of polycrystalline boron carbide ground

dry with the diamond wheel (d = 25.4/xm, V =

0.42 cm/sec) are shown in figure 14. The optical micrograph indicates that the major features of the ground surface are similar to

GRINDING

b I 10 ijin

Figure 11. Optical and electron micrographs of the

surface of G.E. sintered (Lucalox) alumina

ground dry with a 100 grit diamond wheel (d =

25.i/j.m, V = 042 cm/sec). Grain size SSfim.

those exhibited by polycrystalline alumina;

namely, light burnished areas and dark areas where material has been removed by brittle

fracture. The electron micrograph, however,

illustrates some notable differences. Large portions of the fracture surfaces in areas that are

not burnished are conchoidal in appearance

and have no size relationship to the grain size

of the material. In fact, the surface closely re- sembles that of Pyrex glass ground with a similar wheel [16]. Material appears to have been

removed as flakes lying parallel to the ground

surface. Although portions of the surface show

slight indications of grain boundaries (see

arrow in figure 14) clear cut evidence of grain

pull-out has not been observed.

Again, it was noted that boron carbide sur- faces ground both dry and wet contained identical features, but that those ground with a

coolant contained fewer and smaller regions of

burnished material.

6. Discussion of Results

Results of this assessment of surface damage

introduced during surface grinding on various

ceramic materials are summarized in tables 3,

326
1

b L

5 wa b 1 5 urn

j

Figure 12. Optical and electron micrographs of the

surface of AVCO hot pressed alumina ground dry

with a 100 grit diamond wheel (d = 25Afim, V =

0.42 cm/s). Grain size 2/im.

Figure 13. Optical and electron micrographs of the

surface of Coors AD999 cold pressed and sintered

alumina ground dry with a 100 grit diamond

wheel, (d = 25.Vm, V = 0.42 cm/s). Grain size

2fim.

Table 3. Summary of observations made on ground magnesium oxide crystals

Orientation Wheel Rate of a. Nature of ground surface

material removal b. Mechanism of material removal

(100) <100> Alumina

(100) <110>

(100) <120>

Diamond

Alumina

Diamond

Low a. Rough

b. Brittle fracture

Moderate a. Smooth and burnished. Thermal

cracking,

b. Plastic flow

All a. Rough

b. Brittle fracture

Low a. Rough

b. Brittle fracture

Moderate Workpiece shattered

All a. Rough

b. Brittle fracture

Nature of subsurface damage

Discrete plastically deformed

layer.

No subsurface cracks.

Discrete plastically deformed

layer.

No subsurface cracks.

Uneven plastically deformed

layer.

No subsurface cracks.

Discrete plastically deformed

layer.

No subsurface cracks.

Uneven plastically deformed

layer.

No subsurface cracks.

(110) <110>

(120) <120>

Alumina Low a. Rough

b. Brittle fracture

Moderate Workpiece shattered

Diamond All a. Rough

b. Brittle fracture

Uneven plastically deformed

layer.

Subsurface cracks.

Uneven plastically deformed

layer.

Subsurface cracks.

327
to be of two types ; short range and long range.

The short range stresses act at the wheelworkpiece interface and arise from the direct

interaction of the cutting points with the material. Long range stresses arise from the fact

that the workpiece is essentially being forced

between the grinding wheel and the bed of the

machine. Long range stresses are at a minimum when the wheel is removing material effi- ciently, but if the grinding wheel should become loaded and operate inefficiently, the long

range stresses increase. They must then be re- laxed in some manner by deformation of the

specimen or the grinding system, if not the

workpiece or tool will shatter. Relief of long

range stresses by plastic deformation of the

specimen is rarely observed in ceramics since

they are not easily deformed at low temperatures. It is for this reason that many of the

specimens in this study shattered when machined with the alumina wheel at moderate re- moval rates.

At this stage it is not possible to make any

quantitative statements concerning the relative

magnitudes of the short and long range stresses. Efforts are, however, being made to monitor

them. In discussing the origins of surface damage in machined ceramics it is best for the

moment to consider each material separately.

Figure 14. Optical and electron micrographs of the

surface of AVCO hot pressed boron carbide

ground dry with a 100 grit diamond wheel, (d =

25Afim, V = 0.42 cm/s). Grain size 2/ttm.

4 and 5 for magnesia single crystals, silicon

single crystals and polycrystalline ceramics re- spectively. It can be appreciated that a great

variety of surface conditions can be generated

varying from smooth burnished surfaces to

rough fractured surfaces. The process of material removal responsible for these surfaces is

complex. It is a consequence of the interplay

between factors relating to the grinding operation itself, such as wheel type and rate of material removal, and to the microstructure and

mechanical properties of the workpiece.

It has been pointed out [11] that the stress

fields under a grinding wheel can be considered

6.1. Magnesium Oxide Single Crystals

Magnesium oxide is an example of a semibrittle ceramic which shows limited room temperature and extensive high temperature plasticity. Grinding magnesium oxide with an

alumina wheel leads to interesting observations

at moderate and high removal rates. Under

these conditions, the alumina wheel loads up

and fresh cutting points are prevented from

impacting the workpiece, the frictional forces

and therefore, the temperature at the interface

increases. Because magnesium oxide is completely plastic at high temperatures (table 1)

material removal occurs entirely by plastic flow

leaving a completely burnished surface. Thermal quenching after the machining operation is

responsible for the fine cracking also observed

on the surfaces (see figure 4).

Table 4. Summary of observations made on ground silicon crystals

Orientation Wheel Rate of

material removal

a.

b. Nature of ground surface

Mechanism of material removal

Nature of subsurface damage

(111) <110> Alumina All a. Mixed. Burnished regions and Not detectable by etch pit

regions containing brittle techniques.

fracture.

b. Plastic flow and brittle

fracture.

Moderate Workpiece shattered.

Diamond All a. Rough.

b. Brittle fracture.

328
Table 5. Summary of observations made on ground polycrystalline ceramics

Material Wheel Rate of

material removal

Magnesium Oxide Alumina Low

Moderate

Diamond All

Ferrite Alumina Low

Moderate

Diamond All

Lucalox sintered

alumina

Diamond All

AVCU not pressed

alumina

Diamond 411

Coors AD999 cold

pressed and

sintered alumina

Diamond All

AVCO hot pressed

boron carbide

Diamond All

a. Nature of ground surface

b. Mechanism of material removal

a. Rough. Some burnished areas.

b. Grain pullout, transgranular fracture, and plastic flow.

Sample shattered.

a. Rough.

b. Grain pullout and transgranular fracture.

a. Heavily burnished, some rough areas.

b. Plastic flow, grain pullout and transgranular fracture.

Sample shattered.

a. Rough.

b. Grain pullout and transgranular fracture.

a. Rough. Some burnishing.

b. Grain pullout, transgranular fracture, and plastic flow.

a. Rough. Some burnished areas.

b. Grain pullout, transgranular fracture, and plastic flow.

a. Rough. Some burnished areas.

b. Grain pullout, transgranular fracture, and plastic flow.

a. Rough. Some burnished areas.

b. Conchoidal fracture and plastic flow.

Magnesium oxide crystals oriented for machining over {100} surfaces in <100> directions provide a rare opportunity for a ceramic

to relieve long range grinding stresses by plastic deformation (section 4.1.a). It has been

pointed out that for this orientation the crystal

is ideally oriented and can achieve the neces- sary change in shape by shear over {110}

<110> slip systems. It is significant that as soon as the orientation is shifted slightly away

from this ideal orientation the necessary shape

change cannot be accomplished by shear and

the crystals shatter as soon as the v^heel loads

up and burnishing commences (sections 4.1.b

and 4.1.C).

When material is being removed efficiently, as is the case with the alumina wheel wet at

low removal rates and the diamond wheel wet

or dry at all removal rates, the machined sur- face is generated by the initiation and propagation of {100} cleavage cracks due to the impact of the cutting points. This is attributed to

the low impact resistance of magnesium oxide.

For {100} machined surfaces, cleavage leads to

flaking off of material and cracks do not penetrate the surface (section 4.1.a) ; for {110}

and {120} machined surfaces however cleavage

planes are inclined to the surface and subsurface cracking is produced (sections 4.1.b and

4.1.C).

The effect of the grinding fluid is difficult to assess from these limited observations. Clearly,

the fluid prevented burnishing in crystals

ground with the alumina wheel in the {100}

<100> orientation at low rates of material

removal by acting to enhance the removal of

swarf thereby preventing wheel load up. On the

other hand, material was efficiently removed by

brittle fracture with the diamond wheel at all

rates of material removal whether a fluid was

used or not. In this case, the fluid apparently

did not play a role. There is no doubt that the

environment can affect the fracture strength

and grinding efficiency of many materials (see

Westwood, this volume) but from our results

it does not appear that water containing soluble oil has that effect on magnesium oxide.

A study of the effects of different fluids on

the grinding characteristics of magnesium oxide is currently underway in our laboratory.

6.2. Silicon Single Crystals

Silicon is an example of a material which is

brittle at room temperature but can deform on

five independent slip systems at high temperatures. Correspondingly, when silicon crystals are ground with the alumina wheel, the sur- faces are partially burnished (see figure 10)

due to high temperature flow. The extent of

burnishing increases with the rate of material

removal until eventually the crystals shatter.

Shattering occurs because silicon cannot relieve

long range stresses by plastic flow. Normally,

the long range stresses must be relieved in

some other manner. The rough patches of conchoidal fracture in figure 10 indicate that material is being removed in a brittle manner by

329
impact in addition to burnishing. It is presumed that by this semiefficient method, the

long range stresses are relaxed.

Machining silicon with the diamond wheel

results in a rough surface generated entirely

by impact fracture. Again this is attributed to

be low impact resistance of this material in

single crystal form.

6.3. Polycrystalline Magnesium Oxide

and Ferrite

For reasons indicated in section 6.1. polycrystalline magnesium oxide cannot undergo

plastic deformation to relax long range stresses

and thus as soon as the alumina wheel loads up

and burnishing starts, the specimens shatter.

With efficient machining conditions, i.e., with

the diamond wheel, workpiece integrity is

maintained. The machined surfaces of polycrystalline magnesia and ferrite are rough and

contain predominantly intergranular fracture

(section 4.2.). This is attributed to the low impact resistance of these materials.

Subsurface cracking was observed in magnesium oxide crystals ground in all but the

"ideal" {100} <100> orientation and is an ex- pected feature of machined polycrystalline

semibrittle ceramics. This has been found to be

the case in polycrystalline magnesium oxide by

Evans and Davidge [12]

The presence of subsurface dislocations also

appears to be a general occurrence in semibrittle ceramics. Evidence of a subsurface deformed layer in machined polycrystalline ferrite has, in fact, recently been noted [2, 21].

6.4. Polycrystalline Alumina

Alumina is an example of a brittle ceramic

which does not have the necessary slip systems

to become completely plastic at high temperatures.

The surprising feature of the polycrystalline

aluminas studied here is that burnishing is observed as well as transgranular and intergranular fracture on surfaces machined with the

diamond wheel. (Semibrittle ceramics do not

show burnishing when machined with the diamond wheel.) These observations are attributed

to the fact that alumina has a higher impact

resistance than magnesia. Consequently, effi- cient material removal does not occur by fracture alone, there is in addition localized plastic

flow due to the high stresses existing under the

diamond cutting points.

Alumina can be deformed at room temperature by abrasive particles [15] and diamond

hardness indenters [18] so the presence of limited plastic flow is not surprising. What is surprising, however, is that large regions are deformed. Possibly the extra degree of freedom

afforded by the free surface allows gross deformation even if less than five independent slip

systems are operating. Furthermore, the localized grinding temperatures are expected to

be high thereby enhancing the deformation

process. Gielisse and Stanislao [19] have, in

fact, noted white hot illumination to accompany

the impact of a diamond particle on a grinding

wheel with an alumina surface.

More fracture and pullout was observed in

the larger grain size Lucalox than in smaller

grain size AVCO and Coors material. This presumably reflects the decrease in impact resistance with increase in grain size in alumina.

Note that this is opposite to observations on

magnesium oxide cut with a diamond saw [12]

which indicate the amount of pullout to de- crease with increasing grain size.

The effects of wheel dressing on the nature

of ceramic surfaces ground with the diamond

wheel was not examined in this study. Since

dressing increases the density of sharp edges

on the abrasive particles in the cutting face of

the wheel, the incidence of brittle fracture

would no doubt be enhanced. This would de- crease the amount of burnishing observed. The

use of a grinding fluid did decrease the incidence of burnishing in brittle ceramics ground

with the diamond wheel. Whether this effect is

due to surface flushing, interface cooling, or other factors is not known.

6.5. Boron Carbide

Boron carbide surfaces are also extensively

burnished by diamond machining (see figure

14) again indicating a high impact resistance

and tendency to flow under the high short

range stresses imposed by the diamond points.

Relaxation of long range stresses occurs by

the removal of flakes of material by brittle

fracture. It is significant that the fracture sur- faces of these flakes in this material are conchoidal and bear no clear relationship to the

grain structure. This is further indication of

the high intergranular strength and resistance

to impact of this material. The effect of a coolant on the nature of the ground surface of

boron carbide is identical to that observed in

the experiments on alumina.

Alumina and boron carbide surfaces ground

with the diamond wheel contained patches of

burnished material and patches where material had been removed by brittle fracture.

This observation allows one to speculate on the

sequence of events occurring at the wheel workpiece interface during a grinding pass. Two

possibilities exist. On the one hand, material

may be initially removed efficiently by brittle

fracture and then patches on the surface load

up periodically with grinding debris which

burnish. On the other hand, during grinding

330
burnishing could be occurring over the whole

wheel workpiece interface as a result of the

high short range stresses set up due to the

combined effect of debris in the wheel and high

mechanical strength of the workpiece. Burnishing is only a surface effect. To relieve long

range stresses larger amounts of material must

be removed. This happens when portions of the

burnished surface flake off. This process is en- hanced by the large temperature and stress

gradients existing in the workpiece during

grinding. When a patch of burnished material

flakes off a patch where material has been

removed by brittle fracture remains thus ac- counting for the observations made in this

study on ground brittle ceramics. These flakes are quite large in comparison with the theoretical chip size further indicating they are a re- sult of the long range stress fields in the workpiece. No doubt judicious dressing procedures

and light machine settings could decrease or

eliminate the burnishing observed in this study

if necessary. It is not clear, at the present time,

however, if the presence of burnished patches

on the machined surface of a brittle ceramic

is detrimental.

7. Conclusions

On the basis of the work reported here there

appears to be three broad categories of behavior during ceramic machining.

First, there is the burnishing behavior observed on soft, plastically deformable ceramics

ground with a loaded wheel, exemplified by

{100) magnesia single crystals. Under these

conditions the surface of semibrittle ceramics

is generated entirely by plastic flow (although

there are fine surface cracks due to thermal

quenching)

Second, there is the brittle fracture surface

observed on brittle (or semibrittle) ceramics of

low impact resistance, exemplified by single

and polycrystalline magnesia and ferrite

ground with a diamond wheel. Under these

conditions the surface of these materials is

generated entirely by brittle transgranular and

intergranular fracture.

Third, there is the mixed, burnished and

fractured surface observed on brittle ceramics

of high impact resistance ground with a diamond wheel exemplified by polycrystalline

alumina and boron carbide. In these materials

the extremely high local short range stresses

remove material by plastic flow while the long

range stresses are relaxed by brittle intergranular or conchoidal fracture.

Further research is concentrating on substantiation of the ideas expressed here, with

emphasis on direct measurement of the stresses

and temperatures involved and better characterization of the ground surfaces produced.

The author is indebted to Dr. R. J. Stokes

for invaluable discussions and to D. Woodward for able experimental assistance during

the course of this research. The continued interest of Dr. C. H. Li, Director, Honeywell

Corporate Research Center, and Dr. A. Diness,

Office of Naval Research, is gratefully ac- knowledged. This work was supported by the

Office of Naval Research.

8. References

[1] Ceramic processing, pp. 134-173 (Publication

1576, National Academy of Science, Washington, D.C., 1968).

[21 Stokes, R. J., p. , these Proceedings.

[3] Stokes, R. J., Mechanical behavior of polycrystalline ceramics, Ceramic Microstructures, Ed.

R. M. Fulrath and J. A. Pask, pp. 379-405,

(John Wiley and Sons, Inc., New York, New

York, 1968).

[4] Stokes, R. J., Fracture of ceramics, Fund. Phen.

Mat. Sci. 4, 151-175 (1967).

[5] Day, R. B., and Stokes, R. J., Grain boundaries

and the mechanical behavior of magnesium

oxide, Mat. Sci. Res. 3, 355-386 (1966).

[6] Choi, D. M., and Palmour III, H., Fractographic

evidence of multiple slip in deformed hotpressed spinel, Mat. Sci. Res. 3, 473-482 (1966).

[7] Pearson, G. L., Read, W. T. and Feldmann, W. L., Deformation and fracture of small silicon

crystals, Acta Met. 5, 181-191 (1957).

[8]Wachtman, J. B., Jr., and Maxwell, L. H., Plastic

deformation of ceramic-oxide single crystals:

II, J. Amer. Ceram. Soc. 40, 377-385 (1957).

[9] Fitzgerald, L. M., The hardness at high temperature of some refractory carbides and borides,

J. Less-Common Metals 5, 356-364 (1963).

[10] Stokes, R. J., Thermal mechanical history and

the strength of magnesium oxide single crystals: I, Mechanical tests, J. Amer. Ceram. Soc.

48, 60-67 (1965).

[11] Koepke, B. G., and Stokes, R. J., A study of grinding damage in magnesium oxide single crystals,

Journ. Mat. Sci. 5, 240-247 (1970).

[12] Evans, A. G., and Davidge, R. W., The strength

and fracture of fully dense polycrystalline

magnesium oxide, Phil. Mag. 20, 373-388

(1969).

[13] Dash, W. C, Dislocations in Si and Ge crystals,

Prop, of Elemental and Compound Semi-conductors, Metallurgical Society Conf. Vol. 5, p. 195 (Interscience Publishers Inc., New York,

1960).

[14] Stickler, R., and Booker, G. R., Surface damage

on abraded silicon specimens, Phil. Mag. 8, 859-876 (1963).

[15] Hockey, B. J., p. , these Proceedings.

[16] Sedlacek, R., Processing Ceramics-Surface Finishing Studies, Final Technical Report under

Contract No. N00019-67-C-0494, Stanford

Research Institute, Nov. 1969.

[17] Ida, I., Arai, Y., and Inamori, K., Grinding

characteristics of ceramics for circuit components with a metal bonded diamond wheel,

presented at the 1968 General Meeting of

C.I.R.P. at Birmingham, Great Britain.

[18] Hurley, G. F., Observation of plastic deformation in sapphire base ceramics, Met. Trans.

1, 2029-2031 (1970).

[19] Gielisse, P. J., and Stanislao, J., Material re331
moval phenomena in ceramics, Final Technical

Report under Contract No. N00019-69-C-0131,

Univ. of Rhode Island, Dec. (1969).

[20] Newey, C.W.A., and Radford, K. C, Plastic deformation of magnesiiim aluminate single

crystals, Anisotropy in Single Crystal refractory

compounds, Vol. 2, Ed. F. Vahldiek and S. Mersol, pp. 321-338 (Plenum Press, New York,

New York 1968).

[21] Knowles, D., The effect of surface grinding upon

the permeability of manganese-zinc ferrites,

J. Phys. D: Appl. Phys. 3, 1346-1351 (1970).

Discussion

Miller: Why did you select a 100 mesh

wheel? Were you deliberately trying to establish certain stresses?

Koepke: We used a 100 grit wheel to produce

surfaces representative of rough or semi-finish

grinding operations.

Unidentified: You emphasized, or pointed

out at least, the effect of plastic flow. There

can be a tremendous difference between con- ventional and climb grinding in that respect.

Which one did you use?

Koepke: In all cases we used climb grinding,

i.e., the cutting face of the wheel was moving

in the same direction as the workpiece. For

this reason the chips could not be examined.

Since up-grinding resulted in shattering of the

workpiece in many cases, we felt the compressive stresses set up during climb grinding

would aid in preventing workpiece fracture.

This, in fact, turned out to be the case. Unidentified: How much or to what extent

can one attribute shattering to chip crowding?

Koepke: When a grinding wheel loads up

the grinding forces rise very rapidly. In our

experiments shattering almost always occurred

under a loaded wheel. The 100 grit diamond

wheel did not readily load up. As a result, un- reasonably high rates of material removal were

needed to shatter a workpiece with this wheel.

national bureau of standards special PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

332
Observations on Mechanically Abraded Aluminum

Oxide Crystals by Transmission Electron Microscopy

Bernard J. Hockey

National Bureau of Standards, Washington, D.C. 20234

Use of the argon ion-bombardment thinning technique has made possible the

examination of the near surface regions of mechanically abraded aluminum oxide

by transmission eletron microscopy. Observations on diamond-polished (0.25/xm).

alumina-polished (0.3/xm), and diamond-ground specimens have shown that subsurface damage as well as surface damage is typically produced.

Specifically, mechanical polishing introduces relatively high densities of dislocations to a depth of approximately l/xm from the original surface. The dislocations are generally in the form of half-loops and are clearly associated with surface

scratches produced by individual abrasive particles.

The magnitude of residual surface stresses and the irregular surface topography

produced by grinding necessitated the removal of at least 2-4fim from the original

surfaces. At this depth, both the nature and extent of subsurface damage in polycrystals varied from grain-to-grain. Most grains contained either tangled dislocation

arrays or microtwins (either basal and rhombohedral). A number of grains, however, were found to be completely free of damage and may correspond to regions below

fracture surfaces which are apparent in observations of ground surfaces by scanning

electron microscopy.

Key words: Abrasion; aluminum oxide; dislocations; ion-bombardment; microtwins;

scanning electron microscopy; subsurface damage; transmission electron microscopy.

1 . Introduction

Despite the extensive use of mechanical abrasion in the fabrication and surface finishing of

ceramic materials, there have been few studies

on the exact nature and extent of damage produced. Such studies are of considerable value,

not only because the residual damage can be

related to the basic mechanisms involved during abrasion, but because of the dependence of

mechanical behavior on surface perfection.

Abrasive damage is generally classified into

two types; (1) surface damage, consisting of

cracks and flaws (scratches, notches, "pullouts", etc.) and (2) subsurface damage, consisting of lattice defects, such as dislocations

and microtwins. Because most ceramics can be

classified as either "semi-brittle" or "completely-brittle", surface-type damage is commonly produced to an extent which depends

upon the particular material and abrasive ac- tion employed. Subsurface damage, on the other

hand, is generally thought to occur in "semibrittle" materials for which plastic deformation is known to occur at normal ambient

temperatures. The existence of extensive subsurface damage in these materials as a result

of abrasion has been well demonstrated by observations on MgO [1, 2]\ In contrast, previous

studies on scratched [3, 4], fully abraded [5], or "wear" surfaces [6, 7] of aluminum oxide,

which is considered "completely-brittle", have

produced only limited evidence suggesting at

most the generation of isolated dislocations or

' Figures in brackets indicate the literature reference at the end

of this paper.

microtwins during abrasion. The results obtained in these studies, while informative, were

based exclusively on optical examination of as- abraded or etched abraded surfaces.

For this reason, transmission electron microscopy has been employed in the present study

to show directly the full nature and extent of

subsurface damage produced in aluminum oxide

by abrasion. Three surface finishing treatments

were used in the preparation of specimens.

These were: (a) grinding with a resin-bonded

wheel (325 diamond grit), (b) diamond polishing (0.25 /xm diamond grit), and (c) alumina

polishing (0.30 fim alumina powder). Examination of the near surface regions by transmission electron microscopy was made possible by

using ion bombardment [8] to produce the re- quired thinned sections. This technique was

previously used to thin aluminum oxide crystals

in a controlled manner so that the regions sur- rounding and underlying hardness indentations

could be examined by transmission electron

microscopy [9].

2. Experimental Procedure

The bulk of the observations were made on

high purity, large-grained polycrystalline aluminum oxide, which had been obtained commercially in the form of 1/8 in diameter rods.

Mechanically polished specimens 50-75 /xm

thick, which were suitable for subsequent thinning and examination in the electron microscope, were readily prepared by sectioning the

rods into disks, grinding the surfaces flat, and

then polishing both surfaces of the disks using

333
progressively finer diamond compounds. For

diamond-polished specimens, the final surface

preparation consisted of polishing with 1/4

fxm diamond, while alumina-polished specimens

were additionally polished with 0.3 /xm alumina

powder. Thin disks with one ground surface

were prepared by grinding one side of the cut

disks with a 325 diamond grit resin-bounded

wheel and diamond polishing the opposite surface, as above, until the final thicknesses of the

disks were 50-75 /xm.

Observations on the damage produced by

diamond-polishing (0.25 ixm) single crystal

specimens with (0001) surface orientations

were also made. These specimens were prepared from slices cut from Verneuil-grovim

boules and were mechanically polished in the

same manner as the polycrystalline specimens.

In certain cases, the (0001) sections were chemically polished using the hot phosphoric acid

treatment described by Tighe [10] and then

lightly abraded again using 0.25 /xm diamond.

Examination of the damage associated with

the different types of surface finish required

thinned sections containing the original sur- face or regions at various depths below the

surface. This was accomplished by using argon

ion bombardment [8], which has been previously demonstrated as a suitable thinning technique for a number of ceramic materials [9,

11-13] and is now in general use. It is particularly well suited for surface studies on these

materials since the thinning rate from each

surface of the specimen can be controlled independently and the removal of material from

one surface does not alter the opposite surface.

3. Results

3.1. Residual Surface Stresses

Some indication of the nature and strength

of the residual surface stresses introduced by

mechanical polishing was found during the

preparation of thinned specimens by ion bombardment. Unlike specimens used in other

studies, which had been chemically polished,

annealed, or thinned simultaneously from both

surfaces, these specimens spontaneously developed radial cracks when the thickness of the

central area was reduced to approximately

1 fim. Fortunately, the cracks were generally

arrested within the thicker surrounding regions. After further thinning until a central hole

appeared, it was found that the thin, electrontransparent regions surrounding the hole were

curled or bent with the original mechanically

polished surface always convex. This final configuration resulted because the thinned regions

originally bowed-out in the direction which

made the abraded surface concave. It is thus

apparent that mechanical polishing with either

fine diamond or alumina abrasives leaves relatively large surface stresses which are compressive in nature. The extent to which cracks

were produced and the thinned regions curled

indicated larger stresses were produced by diamond polishing. Grinding introduced still

larger residual surface stresses. The magnitude

of these stresses necessitated the removal of at

least 2 fim from the ground surface during the

early stages of the final thinning process in

order to produce whole electron microscope

specimens.

As a consequence of the direction of curling

in polished specimens, some material was un- avoidably removed from the original surfaces

during the later stages of thinning. It was only

within relatively thick regions, where observations with 100 kV electrons were difficult, that

regions containing the original surface or immediately below could be examined. In addition,

the curling of the thinned regions severely re- stricted analysis of observed features.

3.2. Subsurface Damage Produced by

0.25 fim Diamond Abrasive

Observations (including selected area diffraction) in regions of diamond-polished specimens which contained the original surface

served only to confirm the existence of intense

lattice strains. Within these regions there was

even difficulty in distinguishing the individual

grains in the polycrystalline material. Removal

of approximately 2000A resulted in marked

improvement in diffraction contrast. Observations from these regions clearly showed that

relatively high densities of dislocations are produced in aluminum oxide by diamond polishing.

As seen in figure 1, the majority of dislocations are confined to narrow, randomly oriented.

Figure 1. Dislocations prodrxced in polycrystalline

aluminum oxide by mechanical polishing with 0.25 fim

diamond abrasive.

334
linear regions, while the connecting areas contain a relatively low density of dislocations.

These regions of high dislocation density obviously define the traces of surface scratches

produced by the individual abrasive particles.

As a result, the local dislocation density depended upon the size and number of scratches.

Certain areas contained such high densities of

dislocations that the individual dislocations

were difficult to resolve. Differences in the ex- tent of damage associated with the same scratch

in different grains and also between two

scratches of different orientation within the

same grain were also noted. This was due, in

part, to the fact that the majority of dislocations generated appeared to be basal half-loops, as seen in figure 1. This result is consistent with

previous bulk deformation studies [14, 15, 16],

which have established that slip occurs more

easily on the <1120> (0001) systems in sapphire, while the nucleation of half-loops from

the surface is consistent with the mode of

deformation. There were, nevertheless, large

numbers of dislocations, which due to their

particular configurations within the foils, appeared to lie on nonbasal slip planes. The fact

that nonbasal dislocations are also generated

was confirmed by observations made on diamond-polished single crystal specimens with

(0001) surface orientations. In figure 2, the

dislocations associated with a scratch on basal

plane surface not only appear to lie on an

inclined set of parallel planes, but were found

to remain in contrast for all three 3030 type

reflections unlike basal dislocations. Observations on basal plane sections also showed that

high densities of basal dislocations can be

generated. In every case, however, basal disFiGURE 2. Nonbasal dislocations associated with a

single scratch on the (0001) basal plane surface of

sapphire. Scratch and associated damage are typical of that produced by polishing with 0.25 /xm

diamond.

Figure 3. Basal dislocations associated with a relatively deep scratch on the (0001 ) basal plane sur- face of sapphire.

locations were only found associated with

scratches that were larger than those typically

produced by 0.25 ij.m diamond abrasives, as seen

in figure 3.

No attempt was made to precisely measure

the depth to which subsurface damage ex- tended in these specimens. It was found, however, that removal of about 1 /xm from the sur- face eliminated nearly all traces of subsurface

damage produced by 0.25 pm diamond polishing.

3.3. Subsurface Damage Produced by

0.3 Micron Alumina Abrasive

Observations made in regions of aluminapolished specimens which contained the original surface were quite similar to those made

on diamond-polished specimens. The presence

of severe lattice strains again made it impossible to distinguish individual defects. As before, removal of approximately 2000A from the

original surface resulted in a marked improvement in diffraction contrast. Observations in

these regions clearly showed that the extent of

damage produced by alumina polishing was

much less than that by polishing with diamond

of a similar particle size. At this depth below

the original surface, it was estimated that

nearly half of the grains examined were either

dislocation-free or contained only a relatively

low density of individual dislocations.

In the remaining regions, the extent of subsurface damage varied locally. Certain areas were found to contain such large numbers of

scratches and associated dislocations that they

appeared quite similar to areas found in diamond polished specimens, figure 1. These regions, however, were generally located near

pores in which the abrasive grit may have ac335
cumulated during polishing. Other areas were

found which contained relatively few scratch

traces and associated dislocations, as seen in

figure 4. The reason for the uneven distribution of surface scratches and therefore subsurface damage in these specimens is considered to be due to either a wide variation in

alumina particle size or contamination of the

abrasive. It is not considered likely that subsurface damage produced during prior diamond

polishing was still present, since more than

25 fim was always removed by polishing with

alumina powder during the preparation of the

specimens.

In addition to a variation in the extent of

damage in various regions of alumina-polished

specimens, there was also a variation in the

density of dislocations retained within individual grains. Quite often dislocations in one

grain were accumulated near the grain boundary, with no dislocations present within the adjacent grains as seen in figure 5. This feature

was not apparent in diamond polished specimens and may be associated with the "relief

polishing" produced when alumina is polished

with alumina powder.

3.4. Subsurface Damage Produced by

Grinding With a Diamond Resin-Bonded

Wheel

As discussed previously, the extremely irregular surface topography of ground polycrystalline specimens, due in part to a high incidence of fracture, figure 6, together with the

existence of large residual stresses, made it impossible to prepare electron microscope specimens that contained regions of the original

surface or close to it. For these specimens, it

Figure 5. Irregular distribution of dislocations introduced in polycrystalline aluminum oxide by mechanical polishing with 0.3 fim alumina powder.

was necessary to remove at least 2 nm from

the ground surface. The polycrystalline material used, however, contained few grains

whose linear dimensions did not exceed 10 ^m,

so that the observations were made on the or- iginal surface grains.

Examination of these foils revealed an irregular distribution of subsurface damage which

varied in both nature and extent from grain to

grain. A number of grains were found to be

dislocation free and apparently corresponded to

regions below the obvious fracture surfaces seen in figure 6. It is also possible that due to

the highly anisotropic nature of plastic deformation in aluminum oxide, certain grains were

unsuitably oriented for plastic flow even to a depth of 2 fivn. The majority of grains ex- amined within these specimens were, however,

found to have suffered plastic deformation by

Figure 4. Subsurface damage produced within the near surface regions of polycrystalline aluminum

oxide by mechanical polishing with 0.3 fim alumina

powder.

Figure 6. Scanning electron micrograph of diamondground surface of polycrystalline aluminum oxide.

336
Figure 7. Tangled dislocation arrays produced in polycrystalline aluminum oxide by diamond-grinding.

either slip or twinning. The relatively high

density of dislocations in the form of tangled

arrays, shown in figure 1, illustrates the extent

of damage retained within grains in which slip

had occured. Of the two modes of deformational twinning which are known to occur in

aluminum oxide, basal twinning [16, 17] was

by far the more prevalent. An example of a grain containing thin basal microtwins, which

were identified by selected area diffraction, is

shown in figure 8. The presence of both dislocations lying in the boundaries of basal twins

and those apparently generated by the internal

stresses associated with these twins has been

previously documented [9].

Despite the fact that rhombohedral twinning

[18] occurs more readily than basal twinning

in bulk deformation tests [19, 20], very few

rhombohedral microtwins, such as the one seen

Figure 8. Basal microtwins produced within surface

grain of polycrystalline aluminum oxide by diamondgrinding.

Figure 9. Rhombohedral twin, A-A', produced in polycrystalline aluminum oxide by diamond-grinding

Twin appears associated with grain boundary parting along B-C-D.

in figure 9, were found. Figure 9 also illustrates a particularly interesting phenomenon found

within these foils, namely the association of

plastic deformation with grain boundary parting. In the example shown, the rhombohedral

twin A-A', appears unmistakenly associated

with grain boundary fracture along B-C-D.

Although the sequence of events cannot be directly determined from observations such as

this, it would appear more reasonable to as- sume that grain boundary parting occured due

to the development of large internal stresses at

the twin-grain boundary intersection which

could not be relieved by slip or twinning within

the adjacent grain. Accordingly, plastic deformation even though limited to the immediate

surface regions, may be responsible for the

nucleation of penetrating cracks which more

directly limit the strength. Heuer [18] has previously reported that twin-twin intersections

may have been responsible for the nucleation of

failure producing cracks in sapphire tested at

various temperatures.

4. Discussion

Unlike previous studies, the present observations by transmission electron microscopy have

shown that the abrasion of aluminum oxide

does lead to considerable subsurface damage in

the form of dislocations and microtwins as well as surface type damage. Although this damage

has been found to be confined to only the immediate surface regions (approximately 1 fim

for surfaces mechanically polished with 0.25 fim

diamond) the results indicate that plastic deformation does play an important role in the

removal of material by abrasion. Furthermore,

the presence of such high densities of lattice

337
defects near the surface can be expected to

significantly alter the subsequent mechanical

behavior of abraded bodies. At high temperatures, where bulk plastic deformation has been

observed in this material, the surface regions

of abraded crystals can either act as ready

sources of dislocations and twins, as workhardened layers, or as both. At present, there

is no data on the dynamic behavior of dislocations in aluminum oxide at low or intermediate

temperatures (<0.5 Tm) and it is difficult to

ascertain the role of subsurface damage at

these temperatures. It has been shown, however, that chemical [5. 21] and flame-polishing

[21-23], sputtering [24], and prior annealing

[25, 26] do lead to substantial strengthening at

all temperatures. Each of these treatments re- sults in either the removal of subsurface damage or the relaxation of the internal stresses

associated with the high densities of defects as

well as the elimination of surface cracks and

general improvement in surface topography.

Aluminum oxide has also been shown to display

environmental efi'ects, such as static fatigue

[27, 28] , and there is little doubt that the high

densities of lattice defects produced by abrasion contribute significantly due to stress corrosion.

The fact that subsurface damage is produced

in this material in which bulk plastic deformation has not been observed below 900 °C is con- sidered to be a consequence of the nature and

magnitude of the stresses developed under the

individual abrasive particles. As is the case of

static indentation, the load is applied locally

and the stress fields developed contain large

hydrostatic components. Under these conditions, the extremely large shear stresses re- quired for dislocation generation can be developed since full relation of the load cannot

occur by fracture. Surface heating is not con- sidered to be significant in light abrasive actions, such as polishing, but most certainly

plays an important role in cutting and grinding

operations. The fact that penetrating arrays of

dislocations and microtwins, typical of bulk

deformation, are found in the near-surface re- gions of ground specimens but not in polished

specimens is considered to be due in part to the

difference in surface heating produced by the

two operations.

5. References

[1] Cutter, I., and McPherson, R., Examination of

abraded MgO by x-ray diffraction line broadening, Phil. Mag. 20, 489-494 (1969).

[2] Koepice, B. G., and Stokes, R. J., A study of

grinding damage in magnesium oxide single

crystals, Office of Naval Research Technical

Report HR—69—283: 5—26, July 1969.

[3] Palmour, H., Kriegel, W. W., and DuPlessis, J. J., Microbrittleness anisotropy in thermally etched

sapphire, pp. 313-328 in The Mechanical Properties of Engineering Ceramics (Interscience

Publishers. Inc., New York, 1961).

[4] Levengood, W. C, Quantitative aspects of twinning deformation in sapphire and ruby single

crystals, pp. 93-97 in Crystal Grovrth, Proceedings of an International Conference on Crystal

Growth, Boston, 1966 (Pergamon Press, Inc., New York, 1967).

[5] King, A. G., Chemical polish and strength of

alumina, pp. 529-538 in Materials Science Research, Vol. 3 (Plenum Press, Inc., New York,

1966).

[6] Steijn, R. P., On the wear of sapphire, J. Appl.

Phys. 32, 1951-1958 (1961).

[7] Duwell, E. J., and Butzke, H. C, Microscopial

observations of wear surfaces of sapphire, J. Appl. Phys. 35, 3385-3390 (1964).

[8] Paulus, M., and Reverchon, P., Dispositif de

bombardment ionic pour preparations micrographics, J. Phys. Rad. 22, 103 A-107 A (1961).

[9] Hockey, B. J., Plastic deformation of aluminum

oxide by identation. J. Amer. Ceram. Soc. 54,

223-231 (1971).

[10] Tighe, N. J., Jet thinning device for preparation

of AhO, electron microscope specimens, Rev.

Sci. Instr. 35, 520-521 (1964).

[11] Tighe, N. J., and Hyman, A., Transmission elec- tron microscopy of alumina ceramics, pp. 121-

136 in Anisotropy in Single-Crystal Refractory

Compounds (Plenum Press, Inc., New York,

1968).

[12] Tighe, N. J., Microstructure of fine-grain ceramics, pp. 109-133 in Ultrafine-Grain Ceramics

(Syracuse Univ. Press, Syracuse, New York,

1970).

[13] Tighe, N. J., and Hockey, B. J., Ion thinning of

electron microscope specimens; pp. 375-380 in

Record of the lOth Symposium on Electron,

Ion, and Laser Beam Technology (San Francisco Press, San Francisco, California, 1969).

[14] Wachtman, J. B., and Maxwell, L. H., Plastic

deformation of ceramic-oyide single crystals,

J. Am. Ceram. Soc. 37, 291-299 (1954).

Plastic deformation of ceramic-oxide single

crystals, II, J. Am. Ceram. Soc. 40, 377-385

(1957).

[15] Kronberg, M. L., Dynamical flow properties of

single crystals of sapphire, I, J. Am. Ceram.

Soc. 45, 274-279 (1962).

[16] Kronberg, M. L., Plastic deformation of single

crystals of sapphire: basal slip and twinning.

Acta. Met. 5, 507-524 (1957).

[17] Veit, K., Artificial deformation and transpositions

in sapphire, Neues Jahrb. Min., Beil-Bd., 45,

121-148 (1921).

[18] Heuer, A. H., Deformation twinning in corun- dum, Phil. Mag. 13, 379-393 (1966).

[19] Stofel, E., and Conrad, H., Fracture and twinning

in sapphire (alpha-ALOs Crystals), Trans.

AIME, 227, 1053-1060 (1963).

[20] Conrad, H., Janowski, K., and Stofel, E., Additional observations on twinning in sapphire

(alpha-AUO,, crystals) during compression.

Trans. AIME 233, 255-256 (1965).

[21] Mallinder, F. P., and Proctor, B. A., Preparation

of high-strength sapphire crystals, Proc. Brit.

Ceram. Soc. 6, 9-16 (1966).

[22] Morley, J. G., and Proctor, B. A., Strengths of

sapphire crystals. Nature 196, 1082 (1962).

[23] Mallinder, F. P., and Proctor, B. A., The strengths

of flame-polished sapphire crystals, Phil. Mag.

13, 197-207 (1966).

[24] Rice, R., private communication.

[25] Heuer, A. J., and Roberts, J. P., The influence of

338
annealing on the strength of corundum crystals, Proc. Brit. Ceram. See. 6, 17-27 (1966).

[26] Davies, L. M., The effect of heat treatment on

the tensile strength of sapphire, Proc. Brit.

Ceram. Soc. 6, 29-35 (1966).

[27] Charles, R. J., and Shaw, R. R., Delayed Failure

of Polycrystalline and Single-Crystal Alumina,

General Electric Co. Report 62-RL-3081M

(1962) . [28] Charles, R. J., Static fatigue: delayed fracture;

pp. 467-519 in Studies of the Brittle Behavior

of Ceramic Materials, Aeronautical Systems

Div. Report No. ASD-TR-61-628, Ft. II, (1963) .

Discussion

Heuer: Do you think that microhardness

anisotropy in ALOa, which gives rise to rehef

polishing when using ALOa abrasives on AI2O3,

was responsible for the heterogenity you observed in damage due to the abrasives.

Hockey: I mentioned this possibility in the

paper. In alumina powder polished polycrystalline specimens, there were areas in which

grains containing dislocations were surrounded

by defect-free grains. This was never observed

in diamond-polished specimens, which do not

exhibit relief-polishing.

Caveney: What are your estimates of the

depth of damage in the grinding experiments

and what are the grinding conditions?

Hockey: With regard to the grinding conditions, all I can say is that each specimen ex- amined was ground in the same manner using

a 325 diamond grit resin-bonded wheel. Also,

no attempt was made to measure the depth of

subsurface damage, although this is certainly

possible using the techniques described. My

rough estimate of the minimum depth to which

damage extended would be 10-15 /xm. The fact

that large numbers of twins are produced by

grinding and are undoubtedly the most far

reaching defects indicates the extent of damage

will be highly anisotropic.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

339

Acoustic Emission Monitoring of Surface-Damaged

Ceramic Materials

L. R. Bunnell, J. C. Crowe, and P. E. Hart

Battelle Memorial Institute, Pacific Northwest Laboratories

Richland, Washington 99352

The acoustic emission response of Lucalox, fused silica, single crystal alumina,

Pyroceram, and soda-Iime-silica glass to surface damage is reported. Low temperature

annealing at 200 °C reduces the response levels; annealing at 400 °C removes all of

the induced acoustic activity.

Key vvrords : Acoustic emission ; glass ; Lucalox ; Pyroceram ; surface damage.

1 . Introduction

Elastic waves produced in a solid by energy

released as the material deforms or fractures are termed acoustic emissions. These waves can

be detected by a high-sensitivity ultrasonic sensor and, after suitable signal conditioning,

can be used as an indication of crack growth

or flaw formation. The technique has been applied to metals with considerable success, both

for flaw detection in pressure vessels [1]^ and

for weld monitoring. [2] It has also recently

been applied to crack growth detection in ce- ramics. [3]

Intuitively, ceramics surface damaged at

room temperature should be in a condition of

high surface energy and should contain defects which are unstable at elevated temperatures. In order to evaluate the utility of acoustic emission monitoring for surface damage in

ceramics, a series of polycrystalline ceramics

and glasses were tested for changes in acoustic

emission levels upon immersion in hot water

following surface damaging and low temperature annealing.

Samples of polycrystalline AUOs^ (10 and 30

micron grain size), single crystal AI2O3, fused

SiOa,^ glass-ceramic * and soda-lime-silica

(window) glass were treated as described below:

• The specimens, which were all disc-like,

approximately 6 cm^ area x 2 mm thick,

were annealed ^ to remove damage introduced by sample shaping (usually by diamond-sawing).

• Samples were damaged on one face by wet

grinding on Di-Met brass wheels charged

with 9 ixm (Damage #1) or 63 /xm (Damage #2) diamonds.

• Samples at Damage Level #2 were heated

in air to either 200 °C or 400 °C, held at

* Figures in brackets indicate the literature references at the end of this paper.

* Lucalox, General Electric Co.

' General Electric transparent fused SiO,.

* Corning Glass Co.

' Annealing temperatures: Lucalox 1500 "C, Si0„ 1050 °G, Pyro- ceram 800 °C. Glass 500 °C.

temperature for six hours, and furnace

cooled.

• A PZT (lead zirconate-titanate) sensor with a primary response at 600 KHz was

bonded to each sample on an undamaged

edge. The transducer was connected to the

monitoring system which recorded count

rate and total counts.

• The sample was immersed in water at

90 °C, and emissions occurring within four

minutes were recorded.

Figure 1 shows results (total recorded emissions versus time) for the soda-lime-silica

glass. The annealed sample produced no

emissions, while the two damaged samples both

produced emissions. The specimens receiving

the severest damage (as supposed from the

abrasive size) produced the most emissions.

Some decrease in emissions was associated with

the 200 °C anneal, while emission output

dropped to 0 after the 400 °C anneal.

16,000

14,000

12,000

10,000

^ 8000

< 6000

4000

2000

DAMAGE »2

DAMAGE #2+200 C AS -ANNEALED AND DAMAGE #2 DAMAGE #1 I +400°C, OEMISSIONS

2 3

TIME, MINUTES

Figure 1. Acoustic emission response (total) versus

time for a soda-lime-silica glass given the indicated

treatments.

341

439-921 O - 72 - 23
These same trends were followed by all of

the other materials, although results did not

always coincide with the supposed severity of

damage as in figure 1. For instance, in both

Lucalox samples, the sample damaged with 9

fjm diamonds produced more emissions than

that damaged with 63 /xm diamonds.

The paucity of data precludes a detailed analysis of the results; however, the experiment

shows that effects of surface damage can be

detected by acoustic emission monitoring. Measurements on annealed samples showed that

emission resulting from surface damage can be

altered by heat treatment at lower temperatures than would normally be imagined to cause

any annealing effects. The nature of the emission sources, their interaction with water, and

the effect of low-temperature heat treatments

is not understood.

The results indicate that acoustic emission

monitoring is extremely sensitive to surface

damage effects in ceramics. While sensitivity

does not per se guarantee utility, it is hoped

that further experiments will reveal the nature

of the events being detected and thereby give

a greater insight into the effect of surface finishing on ceramic materials.

2. References

[1] Hutton, P. H., Integrity Surveillance of Pressure

Systems by Means of Acoustic Emission, BNWLSA-2194, October 1969. Published in the Proceedings of the First International Conference on Pressure Vessel Technology Part II by the

American Society of Mech Eng.

[2] Jolly, W. D., Acoustic emission exposes cracks

during welding process, published by the Welding Journal Vol 48, p 21-29, January 1969.

[3] Romrell, D. M., and Bunnell, L. R., Monitoring of

Crack Growth in Ceramics by Acoustic Emission, BNWL-SA-3064, published in Materials

Evaluation, December 1970.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing, Froceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

342
V. MECHANICAL AND OTHER EFFECTS OF FINISHING

Session Chairmen

W. H. Rhodes, AVCO Corporation

J. B. Wachtman, Jr., National Bureau of Standards

Effects of Surface Finishing on Mechanical and Other Physical Properties of

Ceramics

Robert J. Stokes ^

Honeywell Research Center, Hopkins, Minnesota 53343

This paper reviews the effects of mechanical finishing operations on the physical

properties of ceramics. Ceramic machining results in a defective surface containing

cracks, dislocations, point defects and residual stresses. The relative significance of

these defects depends on the physical property of interest. The mechanical properties

of single crystals are sensitive to dislocations (semibrittle) and surface cracks

(brittle) ; the mechanical properties of polycrystals are sensitive to surface cracks;

the electrical properties of semiconductors are sensitive to surface trapping sites; magnetic, piezoelectric and optical properties are particularly sensitive to residual

stresses. To optimize physical properties these defects must be eliminated by mechanical lapping, chemical etching or thermal annealing.

Key words: Ceramics; cracks; dislocations; electrical properties; electro-optical; ferrite; machining; magnetic properties; mechanical properties; optical properties; resi- dual stress.

1 . Introduction

There is no doubt that ceramics are finding

increased application because of their unique

physical properties. Until comparatively re- cently it has been sufficient to fabricate ceramics of special shapes by direct methods such as

slip casting or sintering techniques. However,

modern technology emphasizes speed, density

and accuracy and this in turn demands ceramic

materials vi^ith precise control over composition

and topography. In the past ten years, homogeneous and uniform compositions have been

attained through advances in primary fabrication techniques, and dimensional tolerances

have been attained through advances in sec- ondary fabrication techniques, namely machining, grinding and lapping.

Another modern trend in the application of

ceramics involves the fabrication of components with smaller and smaller dimensions. The

result of this trend is to increase the surface

to volume ratio to the point where the properties of surface material are of overriding significance. Thus it becomes of extreme importance to consider the effects of surface finishing

on the critical physical properties for which

ceramics may be chosen.

' Manager, Metallurgy Department.

' Figures in brackets indicate the literature references at the end of this paper.

The purpose of this paper is to consider how

surface finishing operations affect the different

physical properties of ceramics. In order to

do this it will be necessary to review briefly

first the features which are associated with a machined ceramic surface. The consequences of

these features on the various physical properties will then be considered, using direct ex- perimental illustrations where available.

2. Features of the Mechanically

Finished Surface

A number of other papers in this volume and

elsewhere consider the nature of grinding damage in considerable detail and we do not propose to elaborate here. The major features as- sociated with a finished surface are listed in

table 1 according to their origin. As Koepke

[1]^ has pointed out, the surface topography

generated by machining operations depends on

the abrasive wheel being used, the machining

parameter and most important of all, the mechanical properties of the material being machined. Accordingly the surface features depend on whether material is being removed

by brittle fracture or plastic flow and on the

interior structure of the ceramic.

343
2.1 Features Generated by Brittle Fracture

First we consider those features generated

when the machining conditions are such that

material is removed by the propagation of brittle fracture. The fracture mode may be transgranular (i.e.,—cleavage or conchoidal) or intergranular depending on the material. When

the fracture propagates essentially parallel to

the finished surface the result is a general sur- face roughness determined by the faceting of

the fracture surface. The scale of this roughness is obviously related to the microstructure

and has dimensions approximately (Z/2, where

d is the grain size.

When the fracture propagates essentially

perpendicular to the finished surface, fine

transgranular or intergranular surface cracks

are left in the surface. Figure 1 illustrates

transgranular surface flaws introduced into

a magnesium oxide single crystal by a surface

grinding operation. These flaws have a depth

of 150 ixm.

The dimensions of surface cracks are determined by material properties and microstructure and by the severity of the machining operation. In polycrystalline ceramics the surface

cracks typically run until they intercept a grain

boundary and thus have dimensions approximately equal to d, the grain size. Direct measurements by Evans and Davidge [2] indicate

that diamond sawing polycrystalline magnesium oxide produces intergranular surface

cracks with depth approximately one or two

grain diameters for small grain size (less than

100 jam) material and less than one grain diameter for larger grain size material. Rice [3]

has pointed out that the depths of surface

cracks may be related- to the distribution and

Figure 1. Cleavage cracks introduced into the surface

of a magnesium oxide single crystal by a semi-rough

grinding operation. Note layer of etch pits cor- responding to dislocations also in surface layer

(X250).

amount of porosity. In porous material a sur- face crack is likely to intercept and be arrested

by pores thus reducing its dimension.

Another surface feature associated with brittle crack propagation is "grain pull-out," which

is experienced in final lapping and polishing

stages. If earlier rough machining stages have

caused intergranular cracks to propagate one

or two grains beneath the surface, the grains

may not actually separate from the material

until a later stage of the surface finishing operation giving rise to the undesirable "grain

pull-out" craters illustrated in figure 2.

2.2. Features Generated by Plastic Flow

Second we consider features due to plastic

flow. These may consist of either the dislocations and point defects generated immediately

beneath the finished surface or those surface

topographical features associated with the re- moval of material by plastic flow.

One of the consequences of abrasive machining operations is the generation of a defective

layer immediately beneath the finished surface.

The depth of penetration depends on the mechanical strength of the material and the na- ture of the abrasion. The cross section through

the magnesium oxide crystal in figure 1 illustrates the fact that in this soft semi-brittle

material a semi-rough finishing operation generates a layer of extensive plastic deformation

approximately 100 /xm deep as revealed by the

etch pits. The depth of this layer for different

machining conditions has been studied by

Koepke and Stokes [4].

In harder materials the plastically deformed

Figure 2. Polished surface of TViioM Zn^.M ferrite illus- trating surface roughness due to grain pull-out, grain

boundary porosity, and intragranular porosity. Surface etched to reveal grain boundaries (X440).

344
layer is not so deep. For example, in silicon,

hand polishing with silicon carbide under pres- sure forces dislocations 20 fj.m beneath the sur- face [5], whereas polishing alumina with diamond particles injects dislocations only 2-3 fim

deep [6] . In polycrystalline materials the depth

of penetration varies with the orientation of

the individual grains.

Under certain machining conditions [1] the

plasticity of the ceramic material may be sufficient for the finished surface to be generated

entirely by plastic flow. The surface topography

is then smooth and burnished. In some materials the extensive plastic flow at the wheel

workpiece interface can result in excessive heating and the subsequent thermal quenching gives

rise to fine surface cracks as described elsewhere [4].

The steep temperature, pressure and plastic

flow gradients at the wheel workpiece interface

can cause phase transformation or recrystallization with preferred grain orientation in the

finished surface layers [7].

One of the most significant consequences of

irreversible changes occurring in finished sur- face layers is the generation of residual stress.

When the surface layer responds plastically to

the stresses under the machining tool then it experiences a residual compressive surface

stress which must be balanced by the tensile

stresses in the interior. It is important to note

that the average magnitude of the internal

tensile stress depends on the volume of material available. As the specimen shrinks in size

and the finished surface to volume ratio increases, the internal tensile stress increases.

2.3. Features Generated by Exposure of

Internal Structure

Another set of features indirectly associated

with finishing operations arises from the cross sectioning of a sintered block to expose its internal structure. The surface topography generated by final lapping and polishing onerations

should be perfectly flat and smooth. However,

this is rarely the case in polycrystalline ceramics. There are regions where grains have separated from the surface; these may be due to

"grain pull-out" as described earlier or, if due

to the relief of internal stresses remaining

from fabrication, "grain pop-out." There are

other regions where pores are intercepted by

the grinding plane as illustrated in figure 2. The overall surface topography is related to

the microstructure, i.e.,—grain size, pore size

and distribution.

2.4. Summary

The principal features associated with a finished surface are reviewed in table 1. There

Table 1. Features of the mechanically finished surface

Surface features Origin

1. Features generated by brittle fracture

a. General surface roughness Transgranular and intergranular

fracture approximately parallel

to finished surface.

b. Surface cracking Transgranular and intergranular

fracture approximately perpendicular to finished surface.

c. Grain pull-out Intergranular rupture of isolated

surface grains.

2. Features generated by plastic flow

a. Surface dislocations Short range and long range stress- es imposed by the grinding wheel.

b. General surface burnishing Plastic flow and heat generated at the wheel workpiece interface.

c. Crystalline reorientation Plastic flow at the wheel workpiece

interface.

d. Phase transformations Temperature gradients at the wheel

workpiece interface.

e. Residual stresses Plastic flow and chemical gradients.

3. Features generated by exposure of interior structure

a. Pores Within grains or along grain

boundaries.

b. Grain boundaries

c. Grain pop-out Relief of internal stresses.

345
are surface topographical features such as sur- face roughness, grain pull-out and pores and

subsurface features such as cracks, dislocations,

point-defects and residual stress. The relative

extent of these features depends on the nature

of the finishing operation and the quality and

properties of the material. Which feature is of

primary concern depends on the physical property being used.

To identify the sensitivity of physical properties to certain features of the surface condition it is necessary to consider the results of

simple, definitive experiments. We will now do

this where possible for mechanical, magnetic,

electrical and optical properties of ceramics.

3. Effects of Surface Finishing on the

Mechanical Properties of Ceramics

Much has already been written on the effects

of surface condition on the mechanical properties of ceramics. We shall only present a brief

review here. It will be found that dislocations

and cracks represent the most critical features

of the surface condition. For our purpose it will

be convenient to subdivide ceramics into those

which can undergo limited plastic deformation

(so-called semibrittle ceramics) and those which

cannot (brittle ceramics) [8]. A further distinction will be made between the behavior of

single crystal and polycrystalline material.

3.1. Semibrittle Ceramics—Single Crystals

The distinguishing feature of semibrittle

ceramics is that while they can deform plastically they are none the less extremely notch

sensitive. In this instance there is a large differ- ence in mechanical response of single crystals

depending on whether they are perfect or contain surface defects (dislocation loops) or contain surface cracks.

Crystals which are perfect (i.e.,—whiskers

or in some instances, notably MgO, bulk crystals which have been chemically polished to

remove all surface dislocations) deform elastically up to stress levels approaching their

theoretical cohesive strength. However if a single dislocation loop is injected into such a crystal by the slightest mechanical abrasion then

the high strength is lost and the crystal deforms

plastically at its yield stress (see table 2). This

probably represents the most striking demonstration of the sensitivity of a ceramic material

to its surface condition, and illustrates the

significance of the presence of fresh surface

dislocations [9].

If mechanical abrasion of a semibrittle ce- ramic crystal results in surface cracks, the

cracks are extended under stress until they

achieve critical dimensions for brittle fracture.

The mechanism of crack extension has not been

completely resolved, however it definitely involves plastic flow [10, 11]. The important

point about this mechanical behavior is that

the fracture strength is approximately equal

to the yield strength, and is not determined by

the original depth of the crack.

Sensitivity to surface cracks depends strongly on temperature and deformation strain rate.

Semibrittle crystals containing surface cracks

have been found to remain notch sensitive and

brittle within 50 °C of the melting point when

deformed under impact [12].

In view of these mechanical responses to sur- face condition it is reasonable to conclude that

semibrittle single crystals will always be weak

Table 2. Effect of surface condition in the strength of magnesium oxide

Single crystals Strength

Mechanical

behavior

Chemically polished

Mechanically abraded

(MN/m')

1,105

35

(KPSI)

160

5

Elastic

Plastic

Bi-crystals

Chemically polished

Mechanically abraded

760

40

110

6

Elastic

Brittle

Polycrystals

(i) Hot pressed-as received (5-30/itm)

Chemically polished

Mechanically abraded

195

130

28

19

Brittle

Brittle

(ii) Hot pressed-annealed (75 ixm)

Chemically polished

Mechanically abraded

130

100

19

15

Brittle

Brittle

(iii) Sintered (10-20 /tm)

Chemically polished

Mechanically abraded

100

100

15

15

Brittle

Brittle

346
Table 3. Effect of surface condition on the tensile strength of alumina single crystals

Fracture strength

Percent

Surface condition {MN/m') (Kpsi) Theoretical

Perfect (whiskers) 15,900 2,300 61

Flame polished 7,350 1,100 29

Chemically polished (borax) 6,860 1,000 26

Mechanically polished 1,040 150 4

(annealed in 0=)

Mechanically polished 590 90 2

(centerless ground)

As machined 440 60 2

and brittle following mechanical finishing. The

only way in which this situation can be modified is by chemical polishing to completely eliminate the surface defects introducing by the

machining operation.

3.2. Brittle Ceramics—Single Crystals

The distinguishing feature of brittle ceramics is that they cannot deform plastically at low

temperatures. Thus whether surface dislocations are introduced or not by a finishing operation makes no difference to the mechanical re- sponse of these crystals. What does matter is

whether surface cracks are introduced. There

is then a large difference in mechanical fracture

strength depending on surface finish as indicated in table 3.

In the as-machined condition alumina single

crystals deform elastically up to their fracture

stress of 440 MN/m^ Presumably the surface

cracks introduced by machining act as Griffith

cracks and propagate catastrophically at this

stress level. Applying the Griffith relationship

where o-f is the fracture stress (440 MN/m^)

E is the elastic modulus (3 x 10' MN/m^

y is the fracture surface energy

(^ lOJ/m^), and

c is the surface crack depth

gives a value for c of approximately 10 /xxn,

which is reasonable.

Any process which reduces the dimensions of

the surface cracks or eliminates them raises

the fracture stress as illustrated by the different surface treatments in table 3.

3.3. Semibrittle Ceramics—Polycrystals

The sensitivity of semibrittle crystals to sur- face condition extends also to single crystals

containing a grain boundary, i.e.,—simple, bicrystals. This can be seen in table 2. Again a chemically polished bicrystal free from surface

dislocations behaves elastically up to very high

stress levels whereas the defective bicrystal

fractures at a stress corresponding to the single crystal yield strength. The onset of slip

from surface dislocations in a bicrystal results

in internal crack nucleation because the plastic

strain discontinuity cannot be accomodated

across the grain boundary [8].

Attempts have been made to extend this contrast between chemically polished and defective surfaces to more typical polycrystalline

material. The degree of success depends on the

microstructure of the material as indicated in

table 2. Fully dense hot pressed material does

show a significant drop in strength when sur- face dislocations are introduced. In other less

dense materials internal defects mask any

sensitivity to this kind of surface condition.

The strength of machined polycrystalline material is dominated more by surface cracking.

Accepting the idea that the depth of surface

cracks is related to the grain size, then the

fracture strength should vary with the grain

size, d, according to a modified Griffith relationship in which the crack depth, c, is replaced by

the grain size and

CTf = K^d-^/' (2)

where K, is a constant dependent on the material microstructure. When the grain size and

thus the surface crack depth is small, the applied stress exceeds the stress to initiate slip

before fracture occurs; slip initiates and ex- tends cracks at stress levels below the Griffith

stress. The fracture then follows the Fetch re- lationship

CTf = cTi + Kod-^/^ (3)

where o-,, is the stress to initiate slip. For a machined polycrystalline ceramic this gives

rise to the two stage fracture strength—grain

size relationship first noted by Carniglia [13].

When semibrittle polycrystalline ceramics

are chemically polished to remove surface

cracks then fracture will always be initiated by

slip and the strength—grain size relationship

follow equation (3). This has been demonstrated by Evans and Davidge [2] in their experimental work on the effects of surface finishing

on the fracture strength of polycrystalline magnesium oxide. A replot of their data is included

in figure 3.

347
350 -

J/2 , 1/2, l/d (l/^im )

Figure 3. Fracture strength of polycrystalline

magnesium, oxide as a function of grain size

and surface condition. (After Evans and

Davidge-Eef. 2).

Other treatments designed to remove the

surface cracks introduced during machining

likewise enhance the fracture strength at a given grain size. Harrison [14] showed that

fracture strength increased progressively as the machined surface was first lapped with

finer and finer polishing grits, then annealed

and finally chemically polished to give the best

crack free, stress relieved specimen condition.

3.4. Brittle Ceramics—Polycrystals

If there is any connection between flaw size

and grain size, the fracture strength of polycrystalline brittle ceramics should follow the

modified Griffith relationship of equation (2).

However, Carniglia [13] has shown that a two

stage relationship fits the available data better,

and has suggested that fracture is initiated by

plastic flow at small grain sizes. Unfortunately,

the extrapolated value for a, of 300 MN/m^ is

far below the tensile strength of single crystals

(see table 3) and there is no data on chemically

polished crack free material to support the hypothesis. In short, the experimental studies on

the effects of surface condition on polycrystalline alumina are neither as comprehensive nor

as definitive as those on polycrystalline magnesia.

Thermal and chemical treatments of machined alumina have been found to enhance the

fracture strength a small amount [15]. While

the improvements are significant as indicated

by table 4, they are not as spectacular as in

magnesia.

Table 4. Effect of surface treatment on the bend

strength of polycrystalline alumina

Fracture strength

Surface condition MN/m' Kpsi

Chemically polished (borax)

Mechanically polished

(annealed in air)

Plasma flame polished

Mechanically finished

(ground)

540 78.5

470 68.1

440 63.5

440 64.0

3.5. Summary

The effects of surface finishing on the mechanical properties may be accounted for in

terms of surface dislocations, surface cracks

and residual stresses. Which feature is the

most significant depends on whether the materials are brittle or semibrittle. The magnitude

of the effect depends on the microstructure,

i.e.,—grain size and porosity. Minimization of

surface damage by mechanical, thermal or chemical means generally results in an increase

in strength.

4. Effects of Surface Finishing on the

Magnetic Properties of Ceramics

This subject has not been discussed so ex- tensively in the literature as has mechanical

properties. The reason for this is that residual

stress represents the most critical feature of

the surface condition and it is not until specimen dimensions become very small that the

residual stress becomes large enough to exert a measurable influence. As indicated in the

introduction the use of small components

represents a fairly recent development in the

technical application of ceramics and thus the

deterioration in magnetic response is only just

becoming a problem receiving attention.

4.1. Effects of Stress on Magnetic Properties

The effect of stress on the magnetic behavior

of materials is well recognized [16]. The magnitude and sense of the changes induced depend

on the magnetostriction coefficient. When the

magnetostriction coefficient is positive then the

application of a tensile stress permits easier

domain vector rotation. This enhances the intensity of magnetization and thus the permeability for a given field. The remanence is also

increased resulting in a squarer B-H loop. When

the magnetostriction coefficient is negative the

application of a tensile stress makes the B-H

loop flatter. For a compressive stress these effects are reversed.

The effects of stress on the magnetic properties of certain ferrites have been studied in

detail and the results are in agreement with

the general statements made above [17, 18, 19].

The fractional change in permeability "'V/x due

348
to an applied stress (a) has been estimated to

be [17]

9 A (7 JU, 40 tt/^ (4)

where A is the magnetostriction coefficient

and / is the intensity of magnetization.

4.2. Effects on Machining on Magnetic

Properties

If the effects induced by ceramic machining

are due to residual stresses left by the machining operation then there should be certain con- sequences. First, it should be possible to re- move the effects by etching away the highly

stressed surface which the internal material is

supporting, second it should be possible to minimize the effects by a stress relief heat treatment, and third the magnitude of the effects

should be dependent on the surface to volume

ratio. Each of these consequences has been demonstrated at one time or another on a variety of

ceramic magnetic materials, namely garnets,

Ni-Co, Ni-Zn, Mg-Zn and Mn-Mg ferrites [17,

18, 19].

Figure 4 shows the effect of machining and

annealing on the shape of the B-H loop for a

toroid cut from the Nio.se Zno.64 ferrite shown

in figure 1. The machined (lapped and polished) material has a flatter loop than the an- nealed material due to a decrease in permeability, remanence and magnetic saturation. This

particular ferrite has a negative magnetostriction coefficient and the changes are consistent

with the existence of residual tensile stresses.

This observation is significant since tensile

stresses exist in the interior of the finished

ceramic to balance the compressive stresses

left in the surface. Thus the effects of surface

finishing or magnetic properties are primarily

due to changes induced in the volume of the

material. Such changes will be larger as the

(a)

As-machined

and polished

(b) Annealed

Ni-Zn Ferrite

0.35

O30

0.25

E

i/>

?=, 0 20 m

UJ5

- 0.15 m

CD

0.10

Q05

-o~. ANNEALED

GROUND

ANNEALED

GROUND

J I I I I

I I I I 10 SURFACE-TO-VOLUME RATIO, THOUSANDS x m"' 40

Figure 4. B-H loop for Nii.m Zru.M ferrite toroid.

(a) As-machined and polished

(b) Annealed

Figure 5. Variation of Br and Bs with specimen size for

Nio.K Zrio.ei ferrite in the as-machined and annealed

conditions.

volume available to balance the surface stress

decreases, as illustrated for remanence (Br)

and magnetic saturation (Bg) in figure 5 for

the same NiZn ferrite. Note that annealing

causes B, and Bg to undergo greater recovery

as the toroid becomes thinner and the surface

to volume ratio iscreases. The fact that annealing does not completely restore these magnetic properties to their true bulk value indicates there is some irreversible surface damage

which cannot be removed, this is probably

to be associated with surface cracking.

For many applications of magnetic ceramics

surface cracking and surface roughness are of

great significance. In high frequency multiple

gap magnetic recording heads for example, gap

alignment and gap definition are extremely

important if the flux density is to remain uniform over the total width of the head. The C

and I sections of the heads must be lapped

flat before bonding and the edge definition must

be retained during final shaping of the head.

Dimensional tolerances of 0.1 ixm are essential

and the occurrence of grain pull-out or the incidence of porosity in the gap region are generally unacceptable.

5. Effects of Surface Finishing on

the Electrical Properties of Ceramics

Ceramics offer a wide variety of electrical

properties. Depending on their composition and

crystal structure they may be either semiconductors, dielectrics, insulators or piezoelectrics.

Effects of surface condition on the electrical

properties have not been studied directly, nevertheless they are extremely important.

349
5.1. Effects of Surface Finishing on

Semiconductor Properties

In order to indicate how surface condition

might affect the electrical properties of ceramic

semiconductors we have to draw on observations made on the elemental semiconductors,

Ge and Si. The drastic effects of mechanical

surface damage, produced by sawing and lapping was one of the earliest effects recognized

in germanium and silicon transistor research.

Indeed, various chemical etchants have been

developed specifically to remove mechanically

damaged material and this is an important as- pect of silicon technology [20].

The most critical features of the machined

surface are the surface defects, i.e., —dislocations, vacancies and cracks. Their primary effect is to provide trapping sites for the positive

and negative charge carriers, resulting in their

recombination and annihilation. Once annihilated the charge carriers can no longer contribute to the establishment of a current or

potential difference essential in a device operation.

The effectiveness of trapping sites as re- combination centers may be measured in terms

of the surface recombination velocity, i.e.,

the recombination rate per unit surface area. A sensitive method for determining the surface

recombination velocity is the measurement of

the photoelectromagnetic (PEM) voltage generated when a surface is illuminated while inserted in a magnetic field. The lower the sur- face recombination velocity the greater the

density and separation of the positive and

negative charge carriers and thus the higher

the PEM voltage.

Figure 6 illustrates the normalized value of

the PEM voltage (F^) in germanium as a function of the material removed by etching. It can

be seen that different surface finishing operations cause damage of varying depth with the

sand blast treatment being the worst, it damages a layer 32 jxm deep.

Surface recombination velocity is a critical

parameter in the detectivity (D*) of semiconductor radiation detectors depicted in figure

7. Detailed analysis [21] shows that.

(5)

where /i is a complicated function of charge

carrier mobility (fx), lifetime (y) and concentration (n) and /a is a function of the front

(i.e.,—the illuminated) and back surface re- combination velocity (S).

In the photoconductive mode, an electric field

(E) is apDlied during radiation (there is no

magnetic field, i.e.,—B = 0) and the detector

current must be maximized. In this instance

both the front (Sx) and the back (S.) recombination velocities should tend to zero. Thus the

whole surface must be chemically etched to

eliminate mechanical damage.

In the photoelectric-magnetic (PEM) mode,

on the other hand, a magnetic field (B) is applied during radiation (there is no electric

field, i.e.,—E = 0) and the detector PEM voltage must be optimized. In this instance the

charge carriers generated at the front surface

drift down the concentration gradient towards

the back surface. As they drift they separate

to establish the PEM voltage. Charge carriers

reflected at the back surface will be deflected

by the magnetic field in the opposite sense,

generating a voltage tending to cancel out the

PEM voltage. Thus to maximize the PEM voltage reflection must be avoided at the back sur- face by the generation of a charge carrier sink

there. To achieve this, recombination (So) is

deliberately increased by abrading the back surface. This is surely a rare instance where sur0 2 4 6 8 10 12 14 16 18 20 22 24 26 28 30 32 34 36 DEPTH OF MATERIAL REMOVED BY ETCHING, COMPUTED FROM WEIGHT LOSS (MICRONS

)

Figure 6. Effect of surface condition on the PEM voltage of germanium.

(After Buck and McKim-Ref. 20).

350
RADIATION

1 1

+

0 B - MAGNETIC FIELD M E - ELECTRIC FIELD

Figure 7. Diagrammatic representation

of radiation detectors. (Photoconductive mode B = 0, Photoelectromagnetic

mode E = 0).

face damage is used to enhance the performance of a device!

We are not aware of comparable studies on

the effects of surface finishing on the electrical

properties of oxide semiconducting ceramics.

However, the introduction of dislocations and

point defects by machining will certainly generate a variety of trapping sites in the surface

causing effects similar to those noted above.

5.2. Effects of Surface Finishing on Other

Electrical Properties

The most prominent problems associated with

surface finish on dielectrics and insulators are due to surface cracking. The absorption of

conducting materials on the rough surfaces

can result in premature electrical breakdown.

It is for this reason that most insulating ceramics are glazed to avoid the build up of contamination [22],

In the case of piezoelectric ceramics it is to

be expected that the major effects of surface

condition will be those associated with residual

stresses in much the same manner as described

for magnetic materials. The magnitude of the

effect will depend on the piezoelectric coefficient.

Residual stresses have been shown to affect the

temperature at which the ferroelectric phase

transition takes place in barium titanate and

quartz [23]. Obviously, annealing treatments

must be developed for reproducible behavior

in piezoelectric devices, whether used for voltage generation or ultrasonic drivers.

6. Effects of Surface Finishing on the

Optical Properties of Ceramics

Ceramics find many applications by virtue

of their unique optical properties in lasers, light

modulators and electro-optic devices. Surface

finishing is an important aspect in the fabrication of such devices. There are obvious effects of surface roughness on reflectivity and

emissivity although these have not been studied

systematically [22] . Residual stresses also have

obvious effects on the local refractive index.

Precision optical materials must be carefully

annealed to relieve residual stresses.

Recent developments of fully dense hot

pressed lead zirconate—lead titanate (PZT)

ferroelectric ceramics offer unique possibilities

for light beam addressable data storage applications [24], Internal stresses induced by the

application of local electric fields cause light to

be reflected at internal grain boundaries and

modulate the intensity of transmitted light.

Any residual stresses from mechanical finishing operations must therefore be minimized or even eliminated by annealing for successful

operation of such devices.

7. Summary

As ceramics find new applications requiring

precise shape and surface finish it has become

necessary to understand the consequences of

surface finishing operations on the physical

properties being exploited. Ceramic machining

results in a defective surface containing cracks,

dislocations, point defects and residual stresses.

The significance of these defects depends on the

particular physical property of interest. Mechanical properties are particularly sensitive

to surface cracking; electrical properties are

particularly sensitive to surface dislocations

and point defects; magnetic, piezoelectric and

optical properties are particularly sensitive to

residual stress. Special precautions or treatments are necessary to avoid or eliminate sur- face defects. The finished component may need

to be machined, lapped and polished in a precise

sequence; the finished surface may need to be

chemically etched; or the finished component

may need a thermal stress relief anneal to

overcome the deleterious consequences of ce- ramic machining. While this subject has in the

past received a general qualitative appreciation

of the need to improve surface condition in the

application of ceramics, it lacks detailed fundamental studies on the precise magnitude, na- ture and extent of the changes in physical

properties.

The author has benefited immeasurably from

discussion with Drs. J. R. Kench, B. G. Koepke,

and P. W. Kruse. In particular, the results and

analysis of the effect of surface condition on

magnetic behavior in figures 4 and 5 are due

to Dr. J. R. Kench; the studies of the surface

features associated with ceramic machining

are due to Dr. B. G. Koepke; and the information on radiation detectors is due to Dr. P. W.

Kruse.

351
8. References

[1] Koepke, B. G., these Proceedings.

[2] Evans, A. G., and Davidge, R. W., Philosophical

Magazine 20, 373-388 (1969).

[3] Rice, R. W., Paper presented at 72nd Annual

Meeting, American Ceramic Society, Philadelphia, May 1970.

[4] Koepke, B. G., and Stokes, R. J., Journal Materials Science 5, 240-247 (1970).

[5] Stickler, R., and Booker, G. R., Philosophical

Magazine 8, 859-875 (1963).

[6] Hockey, B. J., these Proceedings.

[7] Cutter, I., and McPherson, R., Philosophical

Magazine 20, 489-94 (1969).

[8] Stokes, R. J., Fundamental phenomena in the

materials sciences, Vol. 4, pp. 151-175, (Plenum

Press, Inc., New York, N. Y., 1967).

[9] Stokes, R. J., Trans. Am. Inst. Min. Metall.

Engrs. 224, 1227-1237 (1962).

[10] Stokes, R. J., and Li, C. H., Fracture of Solids, pp. 289-313, (John Wiley and Sons, Inc., New York,

N. Y., 1963).

[11] Clarke, F.J.P.C., Sambell, R.A.J., and Tattersall,

H.G., Philosophical Magazine 7, pp. 393-413

(1962).

[12] Johnston, T.L., Stokes, R.J., and Li, C.H., Philosophical Magazine 4, 1316-24 (1959).

[13] Carniglia, S.C., Journal Am. Ceram. Soc. 48,

580-83 (1965).

[14] Harrison, W.B., Journal Am. Ceram. Soc. 47, pp. 574-79 (1964).

[15] Gruszka, R.F., Mistier, R.E., and Runk, R.B.,

Bulletin Amer. Ceram. Soc. 49, 575-79 (1970).

[16] Bozorth, R.M., Ferromagnetism, p. 620, (D. Van

Nostrand Co., Inc., New York, N.Y. 1963).

[17] Ratheneau, G.W. and Fast, J.F., Physica, 21, pp. 964-70 (1955).

[18] Stern, E. and Temme, D., I.E.E.E. Transactions on Microwave Theory and Techniques 13, pp. 873-4 (1965).

[19] Knowles, J.E., Brit Journal Physics, D: Appl.

Phys., 3, 1346-51 (1970).

[20] Buck, T.M., and McKim, F.S., Journal Electrochem. Soc. 103, 593-97 (1956).

[21] Kruse, P.W., McGlauchlin, L.D., and McQuistan,

R.B., Elements of Infrared Technology, pp. 324-

345 (John Wiley and Sons, Inc., New York,

N.Y., 1962).

[22] Kingery, W.D., Introduction To Ceramics, p. 524,

p. 647, (John Wiley and Sons, Inc., New York,

N.Y., 1960).

[23] Bogardus, E.H. and Roy, R., Journal Am. Ceram.

Soc, 48, 205-7 (1965).

[23] Land, C.E., Sandia Laboratories Report, SCR-67-1219, (October 1967).

Discussion

Rice: One brief comment, twinning can cause

some of the same surface sensitive effects that

slip can and therefore this should also be considered in defining whether a material is totally

brittle or semibrittle.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2^, 1970, Gaithersburg, Md. (Issued May 1972).

352
Strength Effects Resulting from Simple Surface Treatments

H. P. Kirchner, R. M. Gruver, and R. E. Walker

Ceramic Finishing Company, State College, Pennsylvania 16801

In well made polycrystalline alumina ceramics subjected to external loads, fracture

originates, in almost every case, at surface flaws rather than at volume flaws. This

susceptibility to surface flaw failure has been reliably established in experiments in which compressive surface layers were used to obtain substantial improvements in

strength. Knowing that the failures occur at surface flaws, the artificial introduction

of surface flaws can be used to obtain an understanding of the ways in which surface

flaws affect the strength. Additional information can be gained by using simple

treatments to change these artificial surface flaws and by observing the eff'ect of these

treatments on the strength.

In the present investigation artificial flaws were introduced by single point

tools, abrading, thermal shock, and abrasive machining. The flaws were treated by

refiring, chemical and flame polishing, chemical etching, glazing and prolonged storage

in various environments.

The changes in the flaws were characterized by microscopy and to a limited

extent by profilometry. Both flexural and tensile strengths were measured. The

effects of the treatments are discussed in terms of changes in the average strength

and variations in the distributions of the individual strengths.

Key words: Abrasive machining; alumina ceramic; chemical polishing; crack healing; delayed fracture; glazing; humidity; refiring; strength; surface flaws; surface

treatments; thermal shock.

1 . Introduction

Present evidence indicates that there are several classes of flaws that affect the strength of

polycrystalline ceramics. These include discontinuities in the elastic properties, chemical

flaws and residual stresses. The experiments

described in this investigation deal primarily

with the first type, discontinuities in the elastic

properties, and especially those flaws in which

the discontinuity consists of missing material

as in the case of surface roughness or cracks.

The effects of the other classes of flaws were

investigated to a limited extent in delayed

fracture and annealing experiments.

In well made polycrystalline alumina ceramics subjected to external loads, fracture originates, in most every case, at surface flaws rather

than at volume flaws. This susceptibility to

surface flaw failure has been reliably established in experiments in which compressive

surface layers were used to obtain substantial

improvements in strength. Knowing that the

failures occur at surface flaws, the artificial

introduction of surface flaws can be used to

obtain an understanding of the ways in which

surface flaws affect the strength. Additional

information can be gained by using simple

treatments to change these artificial surface

flaws and by determining the effect of these

treatments on the strength.

It is usually accepted that the as-fired sur- faces of ceramics have the best integrity, in

the sense of absence of gross flaws. Therefore,

only limited improvement in strength should be

expected to result from polishing to improve

the smoothness. For example, the specimens

with an average flexural strength of 51,500 psi

(36.2 X 10* kg/m^) in the as-fired condition

were polished. From 0.0001 to 0.0003 in of

material was removed. The amount of material

removed was minimized in order to retain as much as possible of the high integrity skin. The

average flexural strength of the polished material was 56,100 psi (39.5 x 10« kg/m"), an

improvement of nine percent.

In the present investigation, artificial flaws

were introduced into the as-fired surfaces by

scoring with single point tools, abrading, thermal shock, and abrasive machining. The flaws

were treated by refiring, chemical and flame

polishing, chemical etching, glazing, prolonged

storage in various environments, and so forth.

The changes in the flaws were characterized

by microscopy and to a limited extent by profilometry. The effects of the treatments are

discussed in terms of changes in the average

flexural strength and variations in the distributions of the individual strengths.

The principal material used in this investigation was a 96 percent alumina body \ extruded

to form rods that ranged from 0.125 to 0.198 in

in diameter. The average grain size was about

5 micrometers.

2. Strength Effects Resulting from

Simple Surface Treatments

2.1. Humidity, Refiring, and Storage

Humidity is well known to decrease the

strength of alumina ceramics if it is present in

1 ALSIMAG #614, American Lava Corporation.

353
testing environment. The average flexural

strength of the 96 percent alumina decreases

from 48,400 psi (34.0 x 10" fca/m^) when

tested in 18 percent relative humidity to 41,700

psi (29.3 X lO" kg/m^) when tested under

water. This difference represents a decrease of

about 14 percent. Clearly, it is advisable to

make comparative measurements under controlled humidity conditions or at least at the

same relative humidity.

Distribution curves for the flexural strengths

under four different humidity conditions are shown in figure 1. The shapes of the distribution curves are quite alike, but the curves for

84 percent relative humidity and testing under

water do have somewhat greater scatter than

the curves for 18 percent and 32 percent relative humidity.

Storage under humid conditions seems to

have little or no effect on the flexural strength.

For example, soaking in water for 67 hours,

followed by testing in the laboratory atmosphere, gave results similar to those expected

for as received controls.

When specimens were stored in the laboratory for 300 days, a decrease in strength was

observed. This decrease is attributed to the

difference in humidity when the specimens were

originally tested (about 10%) and when they

were tested after storage (about 45%).

Humidity reduces the observed strength of

polycrystalline alumina. If similar material is

loaded almost to the fracture stress in an en- vironment containing water, the material will

fracture within a short time. However, if

water is not present the material will withstand

the load for very long periods of time, if not

forever. This phenomenon involving water is

called static fatigue or delayed fracture.

Fiexural Strength -psi in thousands

10 20 30 40

1.0

50 60

Under water (new lot, .137 in. diameter) Average 41,700 psi (29.3 x 10' Kg/m^ ) 84% relative humidity Average 43.300 psi (30.4 x 10^ Kg/m^)

32% relative humidity Average 46,600 psi (32.8 x lo' Kg/m^ ) 18% relative humidity Average 48,400 psi (34.0 xlO^Kg/m^)

o X o X + • +•

o X o X

10 20 30

Flexural Strength-Kg/m^ in millions

The effect of immersion in water on delayed

fracture performance is shown in figure 2 in

which this performance is compared with that

of similar material tested in a laboratory atmosphere. The observed strength at short time

is similar to that expected from figure 1, but

after 10,000 seconds the strength difference

narrows only to about 2,000 psi (1.4 x 10®

kg/m^). Based upon these results, it is evident

that the delayed fracture curves for increasing

humidity will consist of a family of curves between the two curves in figure 2 (and other curves at higher stresses and lower humidities).

Refired may have several effects on the sur- face of alumina ceramics. Thermal etching may

lead to new stress concentrators. Surface flaws

may be healed. In addition, annealing may

cause reduction of both large scale and localized residual stresses. When 96 percent alumina

rods are refired, the average flexural strength

increases. These results are presented in figure

3.

If more than one type of surface flaw contributes to the observed failures and if refiring

fails to affect at least one of these types of

flaws, then one would expect some of the weak55

50

45 30L

20

• Laboratory Atmosphere

o Under Wafer

X Log Averages

30 E

25

100 1,000 Time -Seconds

10,000 100,000

Figure 2. Delayed fracture of 96 percent alumina.

1000 1100 1200 1300 1400

Refiring Temperature - "C

1500 1600

Figure 1. Distribution curves for 96 percent alumina

at various relative humidities.

Figure 3. Flexural strength of 96 percent alumina

rods refired to various temperatures.

354
est specimens to remain weak after refiring.

In that case, the scatter of the strengths will

increase. Distribution curves for 96 percent

alumina rods before and after refiring at

1500 °C for one hour are given in figure 4. The

scatter of the refired strengths is much greater

than the scatter of the as received strengths.

Some of the weakest specimens do remain weak

after refiring.

Refined specimens were tested before and after storage for 300 days. The distribution

curves are presented in figure 5. Again the

small decrease in average flexural strength is

attributed to the difference in the humidities

when the specimens were tested. The distributions of the strengths remain the same and

there seems to be no definite tendency to re- vert to the less scattered distribution characteristic of the specimen before refiring.

2.2. Flaws Caused by Abrasion

The average strength and the distribution of

the individual strength values can be varied by

abrasion or proof testing, or by combinations

of these treatments. Abrasion, by milling the

Flexural Strength -psi in thousands ,„20 30 40 50 60 70 80

1 1 1

• As Received Controls

1 1 • o

• o Averoge 50.600 psi • o (35.6 X 10^ Kg/m2 ) • o -

o Refired I500''C. One Hour

• o

Average 56.300 psi • o

(39.5 X 10^ Kg/m^)

• o • o

•

o o

• o

•

•

•

o

o

o

•

o • o •

o •

1 1

o

1 1 20 30 40 50

Flexural Strength - Kg/m^ in millions

Figure 4. Distribution of flexural strengths of

96 percent alumina before and after refiring at 1500 °C for one hour.

Flexural Strength-psl in ttiousands 20 30 40 50 60 70 80

Refired to I500°C

^

• Start, Average 56.300 psi ©

(39.5 X I0« Kg/m^) „ ,

o After 300 days. Stored in Laboratoryo • Average 54,200 psi o •

(38.1 X 10^ Kg/m^) o •

30 40

Flexural Strength - Kg/m^ in millions

samples in abrasive grain causes randomly or- iented scratches. These scratches decrease the

average strength and make the strength distributions more uniform. Apparently, abrasion

by fine grain size abrasives introduces small

flaws into the surface of the samples. The flexural strengths of the stronger specimens in the

group are reduced more than those of the

weaker specimens. The flaws introduced into

the surfaces of the weaker specimens are not

very severe relative to those already present so that the weaker specimens are weakened only

slightly if at all. Proof testing removes the weaker specimens

from the distribution. This raises the average

strength of the remaining specimens and de- creases the scatter of the individual strength

values. If specimens are proof tested at a relatively high stress, the remaining specimens will

not fail at stresses substantially lower than the

proof test level because those with severe sur- face or volume flaws were removed from the

distribution. Subsequently, the specimens can

be abraded to reduce the strength of all of the

specimens below the stress at which they were

proof tested. In this condition none of the specimens will exhibit volume flaw failure.

96 percent alumina in the "as received" condition was proof tested at 45,000 psi (31.6 x

lO^kg/m-) removing about 50 percent of the

specimens from the distribution and then

abraded to reduce the strength of the remaining specimens below the proof test level. The

abrasion process. Ten minutes in 240-mesh B4C,

was more severe than expected causing a substantial reduction in strength. Even so, a much

more uniform strength distribution with a

standard deviation of 1700 psi (1.2 x 10^

kg/m^) and coefficient of variation of 4.2 percent, was obtained. The distribution curve is

presented in figure 6 where it is compared with

the results for "as received" and "as received"

Flexural Strengfh-psi In thousands 30 10 50 60

• As Received. Averoge 46.100 psi. ( 32.4 x 10^ Kg/m^)

o As Received, and Abraded Average 41.800 psi (29.4 x 10^ Kg/m^

X As Received. Proof Tested of 45.000 psi. Abraded, Average 40,600 psi (28.6 X 10= Kg/m^)

Figure 5. Distribution of flexural strengths of

refired 96 percent alumina before and after

storage for 300 days.

Flexurol Strength -Kg/m^ in millions

Figure 6. Flexural strength of 96 percent alumina in as received and abraded condition.

355
and abraded specimens. The proof testing and

abrading reduces the scatter almost in half.

The abrasion treatment by itself is quite effective in reducing the scattering of the flexural

strength data and resulted in a standard deviation of 2200 psi (1.55 x Wkg/m^).

Similar reductions in scatter have been observed as a result of abrasion followed by tensile strength measurement.

Flaws were also introduced by abrading the

specimens with coated abrasives. After several

exploratory treatments, two treatments were

selected for further investigation. One of these

treatments consisted of using 220-mesh SiC

abrasive paper to abrade the surface of the

specimen while it rotated in a lathe. A fixed

load was applied by hanging a weight on the

paper. In the other case, 100-mesh SiC abrasive paper was used and only the center of the

rod in a strip one inch wide was abraded. The

scratches caused by these abrasion treatments

are perpendicular to the stresses. The results of

these experiments are presented in table 1. Abrasion with the 220-mesh SiC causes only a

slight reduction in the strength compared with

the "as received" material. This decrease in

strength is more than regained by refiring. The

abraded material that was refired is slightly

stronger than the refired controls. On the other

hand material that is refired and then abraded

is weak like the as received and abraded rods.

Since the gross surface roughness is unlikely

to be much affected by refiring, it is evident

that the flaws that affect the strength and that are introduced by abrasion are microscopic in

character and are healed by refiring. Glazing of

the abraded material further increases the

strength to 76,500 psi (53.8 X 10" kg/m^). This

glaze has a lower expansion coefficient than the

alumina and is in compression after firing. Because of the large difference in elastic modulus

between the glaze and the alumina body, the

gross roughness in the alumina surface as a

result of abrasion still constitutes flaws with

large associated stress concentrations. However, the high strengths observed after glazing

indicate that these are not critical flaws. Apparently, the flaws that weaken the samples

after abrasion are very small and are healed

by refiring and subsequently protected by the

glaze.

Abrasion by the 100-mesh SiC caused very

severe gouging of the surface. The flexural

strength decreased 12,400 psi (8.7 x 10^

kg/m^) to 30,000 psi (21.1 x 10« kg/m^). However, the strength was regained on refiring and

increased substantially when the abraded specimens were glazed.

Other specimens were abraded and the delayed fracture performance under water was

determined. These results are presented in

figure 7. This curve is very similar to that of

the unabraded specimens. Therefore, the available evidence does not show any substantial effect of abrasion on delayed fracture.

2.3. Flaws Made Using Single Point Tools

Two methods were used to make surface

flaws using single point tools. The first method

involved making a single scratch on the alum100 1,000 Time - Seconds

10,000 100,000

Figure 7. Delayed fracture of 96 percent alumina,

abraded by S^O-mesh silicon carbide for 10 minutes

and tested under water.

Table 1. Flexural strength of abraded alumina rods treated by refiring and glazing

Treatment Refiring conditions Average flexural strength

Temp.

°C

Time

hours

No.

specimens psi

kg/m'

X10-"

As received 5 42,400 29.8

Abraded, 220 mesh SiC 5 40,000 28.2

Refired 1500 i 5 50,200 35.3

Abraded, 220 mesh SiC, refired 1500 1 5 51,600 36.3

Refired, abraded, 220 mesh SiC 1500 1 5 41,100 28.9

Abraded, 200 mesh SiC, glazed 1500 1 5 76,500 53.8

Abraded, 100 mesh SiC 5 30,000 21.1

Refired i500 "i 5 52,300 36.8

Abraded, 100 mesh SiC, refired 1500 1 5 51,600 36.4

Abraded, 100 mesh SiC, glazed 1500 1 5 72,100 50.7

356
Table 2. Flexural strength ' 96 per cent alumina rods with axial or circumferential scratches (Rods 0.125 in

diameter)

Axial scratch Circumferential scratch

Load on diamond pointocratcn

width

Average

flexural strength

Scratch

width

Average

flexural strength

grams psi kg/m^-t

XlO um

kg/m^-t

psi XlO

110 20 48,800 34.3 20 42,400 29.8

220 25 50,800 36.7 30 25,700 18.1

420 30 47,500 33.4 35 24,500 17.2

720 30 45,000 31.6 35 20,700 14.5

220 "59,400 41.7(b) 62,900 (b) 44.1(b)

' 50 per cent relative humidity

•> After refiring at 1500 °C for one hr.

ina rod using a diamond point (75° coned

wheel dresser, 1/4 carat) subjected to various

loads. These rods were scratched in axial

and circumferential directions. The flexural

strengths of specimens scratched by the diamond point subjected to various loads, were

measured. In both cases, the scratches were

located in the region of maximum tensile stress

during testing. The widths of the scratches and

the flexural strengths are given in table 2. The

flexural strengths are plotted against applied

load in figure 8. The circumferential scratches

reduce the flexural strengths much more than

the axial scratches. Although the widths of the

circumferential scratches at a given load seem

slightly greater than the width of the axial

scratches, the difference does not seem sufficient

to explain the difference in strength.

Specimens scratched under a load of 220

grams were refired at 1500 °C for one hour. As

a result of refiring, the average flexural

strengths were higher than the strengths before scratching, indicating that the critical

flaws associated with these scratches were

healed by refiring.

100 200 300 400 500 600 700 800

Lood on Diomond Point - qroms

Figure 8. Flexural strength of 96 percent alumina

rods vs. load on diamond point.

The second method involved scoring the

alumina rods with a diamond tool while they

were held in a lathe. The specimens were scored

to various depths. In this case the cuts were

several times as wide as they were deep so that the cut itself was not expected to have a

large stress concentration factor but smaller

cracks resulting from surface damage during

the scoring may have serious stress concentrations. Four grooves were cut in the test section

of each sample. The grooves were approximately one-eighth inch apart. The flexural

strengths were measured and are presented in

figure 9 for various depths of cuts. The samples

always broke at one or the other of the outermost grooves. Using these results, a depth of

cut of 0.002 in was chosen and scored specimens were treated by refiring and by glazing.

These flexural strength results are presented in

table 3. Note that in each case the scored specimens are only slightly weaker than the com-

.002 .003 Flow Deplh- in. Figure 9. Flexural strengths of 96 percent

alumina rods scored to various depths.

357

439-921 O - 72 - 24
Table 3. Flexural strengths of scored " ALSIMAG #6H rods, refired or glazed

(0.15 in diameter rods)

Treatment

conditions

Treatment Temp.

°C

Time

hours

No.

specimens

Average

Flexural strength ^

psi

kg/m'.e

XIO

As received 5 46,800 32.9

Scored to 0.002 in 5 28,900 20.3

Refired, cooled with kiln 1500 1 5 51,400 36.1

Scored, refired, cooled with kiln 1500 1 5 46,900 32.9

Glazed, cooled with kiln 1500 1 5 62,400 43.8

Scored, glazed, cooled with kiln 1500 1 5 57,300 40.2

' Turned on lathe with diamond tip cutting tool to 0.002 in. Four point loading on a 2 in span.

parable unscored specimens indicating that the

treatments are effective in either healing the

cracks or preventing them from acting to cause

failure.

2.4. Flaws Made by Thermal Shock Treatments

Flav^s were also introduced by thermal shock

treatments. In this case the specimens v^ere

quenched in room temperature water from a temperature 225 °C above room temperature.

This treatment is sufficient to reduce the

strength of the rods to about 25 percent of the

"as received" value. The thermal shock treatment leads to very thin cracks that are observable by dye testing. They form as a result of

tensile stresses induced in the surface layer by

quenching.

Dye tested specimens, profilometer traces and

optical and scanning electron micrographs v^^ere

studied in an attempt to characterize these

cracks. The area density of the cracks increases

strongly v^^ith increasing quenching temperature. At the lowest quenching temperature at

which cracks form, all of the cracks are circumferential. As the quenching temperature increases, some cracks bend to join neighboring

cracks. At high quenching temperatures, a checkered pattern of cracks is formed. When

the specimens were dye tested and then broken

at the large cracks, the dye was observed to

have penetrated to a depth of as much as onethird of the sample diameter (0.040 in for

0.125 in diameter samples). The depth and

sharpness of these cracks accounts for the great

loss in strength.

Photomicrographs of a quenched specimen

are presented in figure 10. In this case the

specimen was polished, thermally etched for

five minutes at 1600 °C, and thermally shocked

by quenching in water at 25 °C from an oven

at 325 °C. Figure 10a, a scanning electron

micrograph at 3000X shows a transgranular

thermal shock crack about one micrometers

wide running through alumina grains five to

twelve micrometers in diameter. The crack

seems to be filled with debris. The thermally

etched grain boundaries are one-half to one

micrometers wide. The grains have flat tops as a result of polishing.

Figure 10b, 1000 x, shows a larger area of

the specimen and provides an indication of the

concentration of pores, pullouts, etc. Figure 10c, an optical photomicrograph at

450 X shows the thermal shock crack cutting across many grains. The microstructure of the

body, including the variations in grain size and

shape, porosity, and relationships of alumina

grains to matrix material are evident.

The surface profiles of the thermally shocked

96 percent alumina were studied in an attempt

to detect the cracks. The measurements were

made at American Lava Corporation using a TALYSURF-. One specimen (fig. 11a) was

polished, thermally etched, thermally shocked

(A T = 300 °C) and repolished. Dye testing

showed an average of a major thermal crack

every 0.045 in across the surface. The trace in

figure 11a shows a drop of more than 0.00001

in on the average of every 0.007 in across the

surface. Assuming a hemispherical tip on the

0.0001 in diameter stylus, a crack or hole two

micrometers in width is required to show this

deflection. The thermal shock cracks are not

wide enough and do not occur frequently

enough to account for the large deflections

shown in the figure. Since grain boundary relief formed by thermal etching is also quite nar- row and occurs more frequently than the large

deflections, these deflections cannot be ac2 Measurements made under the direction of Mr. R. D. Dillender.

358
A. 3000 X

B. 1000 X

C. 450 X

Figure 10. Ninety six percent alumina polished, thermally etched and

thermally shocked.

counted for by grain boundaries. Therefore, it seems likely that these deflections represent

rather wide, shallow voids caused by pores,

pullouts, etc. Figure lib shows a trace on the same sam-

-H 1^ 0.002" ^ 0.000005"

4 Microinch (CLA) Finish

A. Polished. Thermolly etched, Thermally Shocked

( AT = SOO'C ) ond Repolished

B. Polished , Thermally etched and Thermally Shocked (AT = 300*)

C. Polished

Figure 11. Talysurf traces for thermally

shocked ALSIMAG No. 6I4. 96 percent

alvmina.

pie in the thermally etched condition before

repolishing. The number of small peaks in this

trace is consistent with the expected frequency

of grain boundaries. In this case, major thermal shock cracks indicated by dye testing oc- curred every 0.080 in on the average across

the surface. Again, the large deflections occur

too frequently to be accounted for by the thermal shock cracks alone.

Figure 11c shows the trace for a polished

sample. After dye testing, this sample showed

no flaws. The large deflections are slightly more

frequent than they are in (a) but the traces are quite similar.

The thermal shock treatments are very effective in degrading the strength of alumina. This

decrease in strength is the result of formation

of very narrow, deep cracks. The surface profiles do not show evidence of these cracks.

Measurements of surface roughness alone are not sufficient to evaluate the relationship between the quality of a surface and strength.

The cracks are thought to lead to severe stress

concentrations. Treatments with 76 percent

HoSO-i + 12 percent HF or 85 percent orthophosphoric acid (H3PO,) were used in an attempt to round these cracks and thereby reduce

their stress concentrating effect. These experiments were unsuccessful, although some im359
provement in strength of the "as received"

specimens was observed as a result of chemical

polishing. The flexural strength results are presented in table 4. Refiring of the thermally

shocked specimens healed the cracks (although

they remain visible at high temperatures because of increased emissivity) and the specimens became stronger than they were originally and comparable in strength to the "as

received" specimens that were simply refired.

As received 5

Reheated with gas torch 5

Thermally shocked 5 Thermally shocked 4

Thermally shocked, reheated with 5

gas torch to heal cracks

(a) Four point loading on a 2 in span.

(b) Some low values in group indicated partial rehealing

The flaws in the thermally shocked specimens can also be healed by gas-oxygen flame

from a welding torch. The results are presented

in table 5. This treatment might be useful for

repairing small parts that are accidentally

thermal shocked during use.

2.5. Abrasive Machining

The evidence presented in the previous sec- tions shows that even though it is possible to

psi X.10

41,800 29.4

41,000 28.8

10,400 7.3 10,900 7.7 41,600 (b) 29.2

Table i. Flexural strengths of ALSIMAG #61i rods with flaws introduced by thermal

shock treatments (a) (Rods 0.125 in diameter)

Treatment

conditions

Treatment Temp. Time

hours

No.

specimens

Average

flexural strength (b)

As received controls

Treatment in 76 per cent H2SO4 + 12 per

cent HF, s.

Treated in H3PO4, 5 min.

Thermally shocked

Thermally shocked, treated in 76 per

cent H2SO4, 12 per cent HF, 15 s. min.

Thermally shocked, treated in H3PO4 5 min.

Refired

Refired, treated in H3PO4, 5 min.

Thermally shocked, refired

1500

Refired, treated in 76 per cent H2SO4 + 1500

12 per cent HF, 15 s.

1500

1500

Thermally shocked, treated in 76 per 1500

cent H2SO4 + 12 per cent HF, 15 s., re- fired

Thermally shocked, treated in H3PO4, 5 1500

min, refired

1

1

1

1

1

5

5

5

5

4

5

5

5

5

5

psi

41,800

42,200

43,800

10,4,00

11,800

11,800

53,300

55,500

52,300

53,900

55,100

55,400

kg/m

XIC

29,4

29.6

30.8

7.3

8.3

6.0

37.4

39.0

36.7

37.8

38.7

38.8

(a) Thermal shocked by quenching in water from 225 °C above room temperature

(b) Four point loading on a 2 in span.

Table 5. Flexural strength of 96 per cent alumina rods, thermally shocked, cracks

healed by gas torch (Rods 0.125 in diam.)

Treatment

No.

specimens

Average

flexural strength (a)

360
vary the strength of alumina ceramics by

abrading, the resulting strength is not related

primarily to the large scale surface roughness

but depends mainly on some kind of microscopic damage. Although no comprehensive investigation of the effect of abrasive machining

on the strength has been made at Ceramic Finishing Company, a limited amount of practical

experience has supported the opinion that careful machining of 96 percent alumina has very

little effect on the strength.

96 percent alumina rods were machined by

two different methods. In one case tensile test

specimens with a necked down diameter of

about 0.125 in and 3 in long were machined

from alumina rods 0.198 in in diameter by 7

in long. The strength of these tensile specimens was measured in flexure by four point

loading on a 2 in span. In the other case, the

diameter of rods 0.125 in in diameter was re- duced in the center to about 0.110 in by forming an arc with a 4 in radius of curvature.

The principal precaution observed in machining these specimens was to rotate the wheel so that the scratches formed by the abrasive were

parallel to the axis of the rods and parallel to

the expected tensile stresses when the rods

were tested in flexure. The machining was done

with a Regular Rimlock Industrial Blade ^ 8

in in diameter by 0.040 in wide and rotated at

900 rpm. The alumina rod was rotated under

the blade at 6 rpm. The coolant was a 40:1

solution of oil * in water.

The distributions of the flexural strengths of

the machined specimens are compared with the

distribution curve of the as received material

in figure 12. The average flexural strengths

and the distributions are very similar.

Machined specimens are also strengthened by

retiring. The distribution for machined and reFleKurol Strength - 40 50

psi in thousands 60 70

•

1 1 As received controls1 1 1

Averoge 50,600 psi (35.6 K 10^ Kg/m2) a*

As machined, ore samples o •©

Average 49,700 psi - uom

( 35.0 X 106 Kg/mZ) CD • - o As machined , tensile m m Samples tested in <D • Flexure di • Averoge 49,100 psi. O) • -

(34.6 K IO^Kg/m2 ) CD • CDS

OCI —

CMC o«o

oaG o<] -

co

o c

1 1

c:

>•

1 1 20 30 40 50

Flexuro Strength - Kg/m^ in millions

Figure 12. Distribution of flexural strengths of 96

percent alumina before and after machining.

fired samples is presented in figure 13. Except

for the weakest specimens, the distribution is

very similar to that of the specimens that are simply retired. At the low strength end of the

distribution the machined and refired samples

are stronger than the specimens that were simply refired. This observation may indicate that a larger fraction of the critical flaws in the

machined samples are capable of being healed

by retiring.

Machined specimens were stored in the laboratory for 300 days. The distribution of the

flexural strengths of these samples is compared

with the distribution before storage in figure

14. The distributions are very similar. Since

the measurements after 300 days were made at

substantially higher humidity than was present before storage one would expect the aver- age strength to be lower than was observed.

Possibly the machined specimens are less sensitive to changes in humidity, during testing,

than the as received specimens.

Machined and refired rods were stored in the

30

Flenurot Strength - psi in thousonds 40 50 60 70 T 1 r

• Retired, ISOCC , one hour 4veroge 56,300 psi . ( 39.6 « 10' Kg/m2)

o futochined and refired 1500 °C one tiour Averoge 59,300 psi. (41.7 » 10' Kg/m^)

• o

• o

30 40 50

FleKurol Strength - Kg/m^ in millions

Figure 13. Distribution of flexural strengths

of 96 percent alumina machined and refired.

Flexural Strength-psi in thousands 20 30 40 50 60 70 -r

Machined

• Slort Averoge 49.700psi (350 x lo' Kg/m^ ) o After 300 Days In Laboratory Averoge 48.400 psi (34.0 X 10' Kg/m^l

' Felker Manufacturing Company.

* 0-214 U.S. Army.

20 30 40 50

Flexural Strength - Kg/m^ in millions

Figure 14. Distribution of flexural strengths

of machined alumina before and after stor- age for 300 days.

361
laboratory for 300 days and then tested. The

results are presented in figure 15. Again, the

distribution of the strengths remained essentially unchanged.

3. Summary and Conclusions

Humidity variations in the test environment

can cause variations of at least 14 percent in

the measured flexural strength of 96 percent

alumina. Therefore, comparative measurements

should be made under controlled humidity or at

the same relative humidity.

Refiring increases the average flexural

strength of 96 percent alumina. The scatter of

the individual strength values increases substantially. In a storage period of 300 days, the

average strength and the distribution of

strengths showed little or no tendency to revert

back to the as received values.

Abrasion by loose grains can be used to de- crease the scatter of the individual strengths.

Abrasion by coated abrasives decreases the

average flexural strength. This decrease is

more than recovered on refiring even in the

case of abrasion v^^ith 100 mesh SiC paper

v^^hich causes substantial roughness.

Scratching and scoring with single point diamond tools decrease the average flexural

strength. Axial scratches are less damaging

than circumferential scratches. Refiring and

glazing are effective in healing the critical

flaws.

Thermal shock (quenching into water) has

the most severe effect on the strength. Again,

the loss of strength is more than recovered on

refiring.

Although no extensive investigation of the

effect of abrasive machining on the strength of

96 percent alumina has been made at Ceramic

Finishing Company, the available evidence indicates that careful machining has little effect

Flexural Strength-psi in thousonds 20 30 40 50 60 70 80

1.0

.4

.2

0

1111 1• o Machined and Refired, I500**C, one hour • o Start-Average 59.300 psi (41.7 x lO^Kg/m^) • o

After 300 Days in Laboratory- ^ ^

om —

Average 59.900 psI (42.1 x 10^ Kg/m^) ^ ^

o •

• o

• o

• o

• o

• o

• o

«> m

om

o • o •

o •

1 1 1 r 20 30 40 50

Flexural Strengtti - Kg/m^ in millions Figure 15. Distribution of flexural strenpths of machined and refired alumina before and after storage

for 300 days.

on either the average strength or the distribution of the individual strengths. Limited evidence from experiments not described in this

paper indicates that fine grained, hot pressed

alumina is somewhat more susceptible to loss

of strength as a result of surface damage.

Therefore, the flexural strength of the damaged material can be increased by careful

polishing. This susceptibility may be grain

size dependent so that this increar e in strength

may not be observed in coarse gr-'iined alumina.

Present evidence indicates that there are several types of critical flaws that may act to

cause failure. One class of these flaws, con- sidered in very general terms, involves discontinuities in the elastic properties of the

material which lead to stress concentrations.

In some cases these flaws may involve missing

material such as small cracks, pullouts or pores

that open to the surface. The importance of

these flaws was demonstrated by the weakening that occurred when the alumina was

scratched, abraded or thermal shocked.

Another class of critical flaws are chemical

flaws.^ The importance of these flaws is indicated by the delayed fracture properties of

alumina. These flaws are not the same as those

above because, when delayed fracture tests on

96 percent alumina are interrupted and the

short time strength is measured, the strength

is the same as that of material that has not

been subjected to the long time static stress.

(In other words, it is not merely a matter of

crack grov^rth under the influence of stress cor- rosion.)

Still another class of flaws that may be important in some cases is the class consisting of

the localized stresses caused by thermal exuansion anisotropy. The critical nature of these

stresses is evident in those materials that show

increases in strength as these stresses are re- lieved by increasing the testing temperature.

In examining these types of flaws that are known to be of critical importance and considering other types of flaws that may be important, it is evident that only the first type of

flaw is affected by the way material is removed

in shaping a 96 percent alumina body. Even

though considerable care is taken in grinding

and polishing, and even assuming that subsurface pores do not become critical flaws when

they come to the surface, the best that can be

hoped for in strength improvement by grinding

and polishing of 96 percent alumina is to shift

the distribution of critical flaws so that other

discontinuities in elastic properties, chemical

flaws, and localized stresses account for the

large majority of the fractures.

' Schmitz, G. K., Gulden, M. E., and Metcalfe, A. G., A Study

of Stress Corrosion of Glass Fibers, Part II, Effect of Glass Composition, Solar Final Report Contract N00019-69-C-0145 (Feb.

1970).

362
This research was sponsored by the Naval their associates at Ceramic Finishing CornAir Systems Command. The authors are pany.

pleased to acknowledge the contributions of

Discussion

Heuer: I would like to make one comment

about your interpretation of interrupted tests.

I think that one can take the view that stress

corrosion or static fatigue can be the result of

the growth of subcritical cracks. Clearly, if you

have a population of "Griffith cracks" in the

sample, fracture in a short term test will result

from their propagation. If you impose a lower

stress on the material, fracture is not observed

until there is growth of the smaller cracks to

the critical size. If this interpretation is correct, one would not expect to see any effect of

"pre-fatiguing" on the subsequent short-time

tests. KiRCHNER: I don't regard this matter as being settled by any means, but I'm thinking

mainly of the work that's being done at Solar

where they have studied stress corrosion of

fiber glass. They find flaws generated in fiber

glass where apparently there was no inherent

flaw to begin with. Furthermore, these flaws

are very periodic rather than random.

Wiederhorn: Regarding the experiment at

Solar, I don't know if the information is transferable. What they showed was that the cracking in their fibers was a result of the ion

exchange—hydrogen ions in the water with the

sodium ions in the glass. You might not have

this situation.

KiRCHNER : I think that's right, but I'm trying to make a point that perhaps the whole

thing is a good deal more complicated than

people thought.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).

363

The Effect of Grinding Direction on the Strength of Ceramics

R. W. Rice

Naval Research Laboratory, Washington, D.C. 20390

It is shown that grinding bars in a direction parallel with the tensile axis has

limited effects on their strength, while grinding perpendicular to the tensile axis may

have no effect or can reduce strengths as much as 50 percent. In single crystals, the

reduction in strength generally increases with the hardness of the material, but can

also depend on crystallographic orientation. In polycrystalline bodies, the reduction

also depends a great deal on grain size, with the effect generally increasing with

decreasing grain size. Also the effect tends to be greater in polycrystalline bodies that are weaker than normal. The effect of grinding direction on strength is attributed

to stress concentrations due to grinding stria, whose continuity and depth generally

increase with hardness. The severity of these stria tends to increase as the amount of

plastic deformation in the machined surface decreases. The amount of plastic deformation is shown to vary approximately inversely with hardness. The effect of varying strength with grinding direction on the interpretation of various tests in discussed.

Key words: Carbides ceramics; fracture; grain size; grinding; hardness; mechanical

testing; nitrides; oxides; strength.

1 . Introduction

Grinding is the most widely used method of

ceramic machining; almost always being the

final finishing operation itself, or the step preceding final finishing. Some surfaces, e.g., flat

ones, generally allow grinding in several, and

often any, directions. Other surfaces, e.g., cylindrical ones, generally make it more difficult, or impossible, to grind in more than a few, or one, directions. Despite the fact that the

option or constraint of grinding direction might

affect the resultant strength under different

stresses, grinding direction-strength relations

have apparently not been studied. That such a

study may be important is shown by Mould

and Southwick's [1] ^ findings that hand abrasion of glass with emergy cloth (150 grit) perpendicular to the tensile axis results in

strengths averaging about 50 percent lower

than when abraded parallel with the axis of

subsequent tensile stressing.

This paper examines some simple grinding

direction-strength relations in a variety of ce- ramics, and shows that significant strength differences can occur, depending on the material

and its microstructure.

2. Experimental Technique

The materials studied, shown in table 1, were

diamond sawn into slabs (except for MgO crystals which were cleaved and the fine grain

ZrOa which was received as fired bars).^ The

slabs were then ground in one direction on one

side using a (Boyar Schultz) surface grinder

with an 8 in dia. x 0.5 in wide (20.2 x 1.27

cm) 100-grit metal bonded diamond wheel op-

' Figures in brackets indicate the literature references at the end

of this paper.

crating at 1725 rpm. A water soluble oil-water

mixture was used as a coolant. The depth of

cut was 0.05^ mm per pass with a feed rate of

approximately 5 cm/min. The second side of

the slab was then ground the same, except that

the direction was 90° to the direction used on

the first side. Polycrystalline and single crystal bodies of each type of material were generally ground together as a group. Some different single and polycrystalline materials were

also ground together. The slabs were then diamond savm (except MgO crystal slabs which

were cleaved) into bars whose cross sections

were nominally 3.8 x 1.8 mm, with the bar

thickness being the slab thickness. The direction of sawing (or cleaving) was parallel with

one of the grinding directions. The edges of all

bars were then beveled by sequentially hand

sanding (dry, and normal to the bar axis) on

120, 400, and 600 grit SiC papers.

For comparative purposes, some ALOs specimens were dry sanded with 150-grit SiC paper.

Individual bars, or narrow slabs were sanded

parallel to the tensile axis on one side and perpendicular on the other side. In order to assure

that only the effect of sanding was observed,

the polycrystalline bodies with previously asfired surfaces and the single crystals were gas

polished [2].

Ambient mechanical testing of bars was on

an Instron machine in 3-point flexure on 1.27

cm span with a 1.27 mm/min head travel rate.

The fracture surfaces, and some ground surfaces, were examined by standard optical and

replica electron microscopy. Grain size was

taken as the linear grain intercept on fracture

surfaces.

^ These bars were ground touching one another to approximate a

plate.

365
Table 1. Materials studied

Material Fabrication Designation Source

(A) Glass

Glassy carbon Pyrolysis Vitreous carbon Beckwith Carbon Corp., VanNuys,

California

Soda lime glass Plate glass Lustraglass double

strength

American St. Gobain Corp., Kingsport,

Tenn.

(B) Single crystals

Al=Os Czochralski Substrate material Union Carbide Co., Union NJ

Verneuil Split boule. Gem material "Walter E. Johnson, Morgan Hill, Calif.

TiO= Verneuil Rutile, gem material w w //

MgA1.04 Verneuil Gem material (approx. 3

MgO : 3.5 AUG,)

DJEVA, Montrey, Switzerland

MgO Shull melting (N)* Norton Co., Worcester, Mass.

MgO ShuU melting High purity (H.P.)* Semi-Elements, Saxonburg, Pa.

(C) Polycrystalline bodies

SiC Sintering KT Carborundum Co., Niagara Falls, N.Y.

B,C Hot pressing Norton Co., Worcester, Mass.

Si,N. Hot pressing (Plessey Co., England) Courtesy of Dr. Lindley, Admiralty

Materials Lab., England

MgAl=04 Hot pressing By author

MgAhO, Sintered Transparent Courtesy of D. Sellers of Coors

Porcelain Co., Golden, Colo.

AW, Sintering Lucalox General Electric Co., Cleveland, Ohio

AhO, ATD-998-A Courtesy of J. Flobeck, Coor Porcelain

Co., Golden, Colo.

AhO, Sintering AD-94 Coor Porcelain Co., Golden, Colo.

AUO3 Sintering Alsimag 614 American Lava Co., Chattanooga, Tenn.

AUO3 Hot pressing By author

AUO3 Sintering Courtesy of J. Bailey, American Lava

Co., Chattanooga, Tenn.

ZrO,

ZrOa

Hot pressed,

^ annealed > Sintered J

Zyttrite

Courtesy of K. Mazdiaysni, Air Force

Materials Lab., Wright-Patterson

AFB, Ohio

ZrO^ Sintered Fine grain (Zircar) Courtesy of A. Naumann, Union Carbide Corp., Tarrytown, N.Y.

*Author's designation

3. Results

The results for glass and oxide crystals are shown in table 2, which also indicates the general relationships between the tensile axis and

crystal directions.^ The orientations of the

sapphire specimens from the split Verneuil

^ Note the designation of the Ti0„ boule axis as the "c" axis is based on the suppliers information. The tensile direction is given

as approximately the 110 direction since (a) the specimens were ground on a plane which was 90° ± 5° to the common

cleavage plane and in a direction that was ± 10° to the inter- section of this cleavage plane and the plane of grinding; and (b) reference 3 reports that the only plane of cleavage observed at room temperature is 110 cleavage.

boules is shown schematically in figure 1, along

with summary strength data. The data on bars

cut normal to the boule axis show more scatter

because they are from three different boules.

Comparison of individual specimens indicates

that the strength difference due to the grinding

effect may be 30 percent or more. The results

for polycrystalline specimens are shown in

table 3. Fractures were examined to determine the

fracture origins (where feasible), compare

fracture features (e.g., mirrors on glass and

366
Table 2. The effect of grinding direction on the bend strength of glass and single crystals

Strengths, as-ground

Material

Typical naruness

vickers

Number

01

tests

Parallel to

tensile axis

Perpendicular to tensile axisx^ercenx

decrease

kg/mm

10 ' psi 10' N/m' 10 ' psi 10 ' N/m'

Soda Lime Glass 550 7 14±1 9.7 10±1 6.9 27

Glassy Carbon 150-300 7 18±4 12 14±3 9.7 12

Sapphire 1 (AhO,) ' 7 48±9 33 24±7 17 50

Sapphire 3 (AhO,) ' 9 48±9 33 26±6 18 47

Sapphire 1, 2, 3, (AUO3)' 7 75 ±18 52 58 ±13 40 *22 (30)

Sapphire 0° ' 2400 4 50±6 35 38±3 26 24

Sapphire 30° ' 4 55±5 38 52±11 36 5

Sapphire 60° ' 4 61±2 42 59±6 40 3

Sapphire 90° ' 4 48±3 33 46±5 32 6

Spinel ' 1 5 34 + 2 23 21±1 14 40

^ 1800

Spinel4 38±3 26 20±3 14 47

TiO^" 1800 3 13±1 9 11±1 7.6 15

MgO (N)' 800 7 23±4 16 23±3 15 4

MgO' (H.P.) 600 4 20±3 14 20±3 14 0

' Sapphire from boules 1 and 3 were cut parallel to the boule axis as shown schematically in figure 1.

' Sapphire specimen from boules 1, 2 and 3 were cut perpendicular to the boule axis as sketched in figure 1.

' The nimiber of degrees refers to the angle between the normal to the tensile surface and C axis. The tensile

direction is <1010>.

* AUOa-rich spinel, approx. 3 MgO to 3.5 AI2O3, tensile axis was parallel to the boule axis.

' The same AliOs-rich spinel with the tensile axis pei-pendicular to the boule axis.

'Tensile axis perpendicular to the boule (c), approx. <110> direction.

' Tensile axis is <100> direction.

" See text.

single crystals), and measure grain sizes on the

polycrystalline specimens. As shown in table

4, many origins could be found and most were

from the tensile surface. The greatest percentage of fractures originating from the edge were

in specimens ground parallel to the tensile axis

(i.e., the higher strength condition), especially

in glass, which is known to be somewhat sensitive to edge effects. Also, three of the four 90°

Czochralski sapphire bars ground parallel to

the tensile axis failed from the corner. Fracture

mirrors were measured on glass and single

crystals of MgO, AI2O3, and TiOa, and found

to decrease with increasing strength in accordance with other testing (which will be covered

in a later paper). While some of the edge fracture origins were in specimens failing at lower

relative strengths, others were in specimens of

at least medium or higher relative strengths.

The results of sanding trials are shown in

table 5. Examination of the sanded sapphire

fractures showed six of eight originated at the

GRINDING DIRECTION

Figure 1. Schematic representation of initial observations on the effect of grinding direction on strength

of sapphire. This sketch also indicates the orientations of bars from split boules. The tensile surfaces are {1010} and the tensile directions are within 20°

of <1120> and <2423> for specimens cut with their

tensile axes respectively parallel and perpendicular

to the boule axis.

367
Table 3. The effect of grinding direction on the strength of polycrystalline bodies

Characterization Strength As-ground

Materials

Fabrication

Designation

Approx.

percent

porosity

Approx.

grain

size /jLiii '

Number

of

tests

Parallel to tensile axis

Perpendicular to tensile axis

Percent

change

10 ^ psi 10

N/m

10 ^ psi 10 ' N/m'

SiC s 4 100 9 30±5 21 27±5 19 -10

B4C H.P. 0 30 6 56±8 39 49±5 34 -11

H.P. 0 1 3 95±9 66 52±2 36 -45

S ^Tiiipfllov^ 0 60 7 37±2 26 39±3 27 + 4

0 50 4 39±4 27 34±4 24 -13

Al-0^ S (AD-di) 7 20 3 51±2 35 48±3 33 - 6

AhO, S (A614) 6 12 4 60 ±6 41 54±4 37 -10

AI2O3 H.P. 0 9 5 66±5 46 59±5 41 -12

AljOs s 2 7 g 60±8 41 33 ±3 23 -45

AI2O3 H.P. 0 6 1 71 49 47 32 -34

MffALOi i.TJ. g xn. 12 V4 s 0 100 5 31±4 21 30±4 21 - 3

M2'AL04 H.P. 1 6 4 41±2 28 26±4 18 -37

Zr02 H.P.-A. 1 15 5 48±6 33 41±5 28 -15

ZrOi s 1 10 4

on _u Q 1 Q -+- -40

ZrOa s 1 1 3 116±3 80 100 ±2 69 -14

MgO H.P.-A. 0 100 3 18±2 12 20±2 14 + 11

MgO H.P.-A. 0 35 3 27±2 19 31±1 21 * +15( + 1)

MgO H.P. (LiF)-A 0 35 3 19±2 13 27±1 19 ' +32(4-7)

MgO H.P.-A. 0 10 3 39±3 27 28±8 19 -28

MgO H.P. (LiF) 0 3 3 35±2 24 20±5 14 -43

' S: Sintered; H.P. Hot pressed; (LiF) means hot pressed with LiF; -A means annealed in air after hot pressing

(before grinding).

^ Designation : Materials are listed in the same order as in table 1 ; however, some of the designations are repeated

here for clarity.

' Note the grain sizes listed are the average of the larger grains, which control strength. The matrix grain size,

which probably is more important in the grinding effect, is less (often half) of the value shown.

* Corrected per figure 6 as noted in the text.

tensile surface, and only two from the specimen edge.

While an exhaustive characterization of the

surfaces was not feasible, some clear trends

were evident. Grinding stria (fig. 2, 3) were

readily visible with the unaided eye on sapphire

and to a lesser extent on glass (least on the

glassy carbon), but were generally visible on

the spinel crystals only with the microscope,

while such stria could not be seen optically on

TiOo or MgO crystals. A similar trend was seen with the polycrystalline materials, but there

was an important superimposed effect of grain

size. Thus, grinding stria could be seen with the

unaided eye on all the ZrO; bodies ^ and the

' Note, in the case of the fine grain Zr02 bodies, the stria were

not as pronounced, but were quite clear because of the very glossy nature of the ground surface.

finer grain AI2O3, MgAloO^ and SigNi bodies,

and to a lesser extent on BjC, but not on the

fine grain MgO. The stria were clearly defined

under the optical microscope on the SiC and

the medium grain ALO., bodies, but could not

be clearly defined on the largest grain AI2O3.

Only limited sporadic sections of stria were

observed optically on the large grain spinel.

On the MgO, stria were moderately clear optically on the fine grain body, and not defined on

the largest grain body. The decreasing clarity

of the stria appeared to result from both a de- crease in the depth of the grooves of which

they consist, and a decrease in the continuity

due to surface spalling, grain pull-out, and re- lated removal. Sample electron microscopy cor- roborated these results, and more clearly

368
Table 4. Fracture origins of specimens ground

Parallel to the tensile axis Perpendicular to the tensile axis

Material

No. of

tests

No. of

origins

located

No. from

tensile

surface

No. from

specimen

edges

No. of

tests

No. of

origins

located

No. of

origin

from

tensile

surfaces

No. of

origins

specimen

edges

A. Glass

Soda^me 7 7 4 3 7 7 5 2

Glassy carbon 7 7 5 2 7 7 1 6

B. Single crystals

AW, (1) 8 8 6 2 7 7 6 1

AI2O3 (3) 9 6 3 3 9 6 6 0

AUO3 (1, 2, 3) 7 6 3 3 8 5 4 1 AUOa (Czochralski) 16 16 11 5 16 16 14 2

MgAUO, (1) 5 5 3 2 5 5 5 0

MgA1.04 (2) 5 5 3 2 5 5 5 0

TiO, 3 3 3 0 3 3 2 1 MgO (N) 8 8 7 1 6 6 6 0

MgO (H.P.) 4 4 3 1 5 5 5 0

C. Polycrystals

SiC 9 0 9 0

B4C 6 5 5 0 6 4 4 0

Si^N, 3 0 3 0

MgAUO, 5 5 3 2 5 5 3 2

MgAUO^ 4 2 2 0 4 3 3 0

AUG. 32 0 32 0

ZrOj 11 8 5 3 11 9 5 4

MgO 15 5 4 1 15 8 7 1

' Tensile axis parallel to the boule axis.

^ Tensile axis perpendicular to the boule axis.

Table 5. The effect of standing direction on the bend strength of oxides

No. of Strength Strength Percent

Material Grain size tests sanded parallel sanded perpendicular reduction

1000 psi 10 ' N/m

1000 psi 10' N/m'

AI2O3 Single crystal 4 50±8 34 30±1 21 -38

AUG, 20 5 53 ± 37 53± 37 0

AUG, 7 4 56±3 39 52±3 36 -7

' Note the grain size is the average of the larger grains since they control strength. The matrix grain size, which

probably is more important in this sanding effect, is less (often half) of the value show^n.

showed the interruption of stria by spalling

and related processes (fig. 3).

Optical examination of the sanded polycrystalline AI2O3 showed that striations were

lower in density, continuity, and probably in

height than on ground specimens. Sanded gas

polished sapphire also showed substantially

fewer striations, but their continuity and

depth more closely approached that of their

ground counterpart than did the polycrystals.

369
Figure 2. Examples of grinding Striation. Grinding was left to right. (A) Polycrystalline SisN*; (B)

Sapphire; (C) Polycrystalline B4C; (D) Large grain

polycrystalline MgAUOi; (E) Rutile crystal; (F)

MgO crystal.

4. Discussion

4.1. The Effect of Grinding Direction on the

Strength of Glass and Oxide Crystals

Table 2 clearly shows that there can be a

pronounced effect of grinding direction on the

strength of soda lime glass and some oxide crystals. However, except for some sapphire orientations, it is clear that for the crystalline

materials, the decrease in strength generally

decreases with decreasing hardness, becoming

zero for the softest MgO. This is consistent

with other observations and with the trend of

the grinding stria as follows. The author has

shown that (a) the depth of slip from machining increases with decreasing yield stress

[4], and (b), the yield stress (Y) of ceramics

is generally one-third of the hardness as in

metals [5]. Further, in NaCl [6] and MgO

[7] , it has been observed that the length of slip

bands surrounding hardness indents are inversely proportional to the hardness (H). This

is generally consistent with the above results,

and suggests that the depth of slip from machining may be a function of 1/H, and hence,

1/y. Brace [8] has shown that even mild sandC

Figure 3. Sample electron micrographs of

samples ground left to right. (A) Sapphire;

(B) soda lime glass; (C) fine grain polycrystalline MgO.

ing of fine grain NaCl results in a workhardened layer of up to 1 mm in thickness. The

author has shown that grinding of MgO nor- mally introduces slip to two to three times the

depth that sanding does [4], indicating that

similar increases should occur in NaCl. Using

this information and machining data for MgO,

TiOj (see Appendix), and sapphire [9], a definite trend is seen in figure 4. Decreasing grain

size decreases the depth of slip in machined

MgO, so data for single crystal NaCl would be

expected to be higher than that based on

Brace's [8] fine grain NaCl, indicating even

better agreement. The author has shown elsewhere [4] that data for CaO and for LiF indicate that these materials which are progressively softer than MgO develop greater depths

of slip from machining than MgO. On the other

hand, ZrOo and MgAloO,, which are progressively harder than MgO, indicate less depth of

370
machining effects than MgO. These observations corroborate the trend shown in figure 4. The increasing clarity, and hence continuity

and depth of grinding stria with increasing

hardness, can be generally related to the above

slip depth variations as follows. The impact

and frictional energy of grinding can be dissipated by fracturing, deforming or heating of

the surface material of the piece being machined (with heating aiding deformation). Increasing the amount of deformation in the

material decreases the energy available for

fracturing the surface heating. The energy ab4 7 1000

Hardness

I / 2,-1

in 10 (kg/mm )

Figure 4. The depth of slip from grinding as a function of hardness.

radius; R

CONCENTRATION OF NORMAL TENSILE STRESS = 1 + 0.7

FiGXntE 5. Stress concentration due to a

surface step after Marsh (10).

sorbed in deformation is proportional to the

volume of material deformed, and hence to the

thickness of the layer of slip. As the layer of

slip decreases, plastic deformation will be more

localized around the individual grinding particles, due in part to more surface heating, re- sulting in more pronounced stria. The increasing stria clarity with increasing hardness is

thus reasonable.

The decrease in strength due to grinding perpendicular to tensile axis thus appears to be

related to the orientation and extent of striation. These striations are essentially a series of

surface steps. Marsh [10] has shown that such

steps cause stress concentrations as shown in

figure 5, which for larger step heights and

smaller root radi results in stress concentrations as great as for Griffith flaws. It is also

clear that reducing the continuity of the steps

and their angle relative to the tensile axis will

reduce the stress concentration as observed.

The stria may be supplemented by fine (apparently discontinuous) cracks at or near the root

of these steps (which would be one-half of each

striation). The fact that almost all specimens

ground perpendicular to the tensile axis failed

away from the edge of the specimen indicates

that such small cracks or variations of the step

height occur to prevent failure from the edges.

The author [4] has observed that heavy

grinding of MgO crystals increases their

strengths 30-50 percent. However, moderate

chemical polishing to give a smooth surface

but leave most of the work hardened layer increased strengths still further (up to 50%

more). Therefore minor, local steps due to

spalling and related processes can also affect

strengths. Other tests [11] showed that the

strength of ALOs and spinel ground parallel

with the tensile axis were essentially the same

as mechanically polished specimens. Polishing

of glass did improve strengths some (e.g.,

14%). Trials indicated that perpendicular

ground spinel had to be polished quite extensively to eliminate the effect of previous grinding direction on strength.

Thus, while the introduction of oriented

Griffith flaws might have been expected to

be the primary reason for the marked difference in strengths with grinding direction,

stress concentrations due to surface steps appears to be the main cause. Further, flaws

would be expected to often occur perpendicular

to the stria and hence to give lower strengths

for specimens ground parallel to the tensile

axis. Also, an exclusive Griffith flaw-type failure would not be consistent with microplastic

mechanisms of failure that are known to oc- cur in MgO (and indicated in fig. 6), and might

be factors in the strength of AI2O3 and TiOa.

However, some results, e.g., on polished spinel,

371
400 100 50

GRAIN SIZE IN 4

"As hot pressed 45

40

o

35 S

30 %

3

IT 25 u.o

to

20 5

MgO

Hot pressed As

ond hot annealed pressed

Specimens ground parallel to the tensile axis Specimens ground perpendicular to ttie tensile axis

20 30 40 50 60

I 70 80 90 100

Figure 6. Strength versus grain size of ground MgO.

indicated that flaws were probably also generated in some cases. The decrease in the strength of soda lime

glass observed here is only about half that re- ported by Mould and Southwick [1]. However,

their tests were run in a wet condition and for

longer times to study stress corrosion, which

probably accounts for the greater effect they

observed. It is interesting to observe that application of Marsh's [12] theory of hardness

and plastic flow in glass would approximately

bring the soda lime glass result in line with

the single crystal results. Marsh gives yield

stress for glasses of about 450 kg/mm-, which

would correspond to a single crystal hardness

of over 1000 kg/mm^

The smaller effect in the glassy carbon is

generally consistent with its lower hardness.

It is uncertain whether Marsh's theory applies

to this material [5], thus the effect may be

greater here relative to the other materials.

This might arise from the low modulus allowing greater elastic strains which may reduce

spalling and hence increase the continuity of

the striations which are otherwise fairly moderate.

The one clear exception to the above trend of

increasing stria clarity and strength difference

with increasing hardness are some of the

Czochralski sapphire specimens. The generally

softer nature of Czochralski compared with

Verneuil grown sapphire, and the variation of

hardness with crystal orientation may explain

part of the differences, e.g., the CzochralskiVerneuil differences. However, the differences

for the 30°, 60°, and 90° specimens are much

smaller than could be expected from the general hardness effect. Hockey's observation that

both hardness indenting and grinding introduce

many twins on non-basal planes (e.g., the

rhmbohedral plane) but few, if any, on the

basal (0°) plane (13), suggests a correlation.

Twinning might alter either the nature of the

stria or of possible associated flaws. However,

since strength tends to increase (by nearly 2) as the tensile surface departs from the basal

plane [14], the lower variation observed here

in the parallel ground condition may be due to

twinning limiting their strength. Twinning

would normally be at an angle to the stria, and

thus be effective in either orientation, and limit

the rise in strength with departure from the

basal plane observed in other (e.g., polished)

specimens.

4.2. The Effect of Grinding Direction on the

Strength of Polycrystals

The extent and clarity of the grinding stria

on polycrystals shows general agreement with

the concepts discussed for single crystals, but

with an important superimposed effect of grain

size. This grain size effect is due in part to

effects discussed above. First, hardness of a material does generally increase with decreasing grain size, with the effect on microhardness

being greatest at relatively fine grain sizes

[5, 15].

372
Second, the author has shown that when the

depth of slip exceeds the grain size, then further decreases in grain size result in some re- duction of the depth of slip in MgO [4]. Similar effects would be expected in other materials.

This grain size effect on slip depth is consistent

with the increase in hardness (though it may

be somewhat greater than expected from hardness increases). These both, in turn, corroborate the concept of reduction in the extent of

slip increasing the extent of striation and re- sultant decreases in strength in specimens

ground perpendicular to the tensile axis.

However, these effects do not appear to be

sufficient to explain the substantial reduction

of the directional effect in large grain bodies,

especially in the harder materials. Spalling and

grain pull-out appear to be the major factors

since they interrupt the continuity and hence

the stress concentration of the striations. Spalling probably occurs more in larger grain bodies

because they are weaker. Since such spalls are normally limited by the grain size (due to grain

boundary fracture), larger grains result in

greater interruption of the striations.

The strengths of most parallel ground polycrystals were generally about the same [11]

and sometimes possibly greater than specimens

of the same materials with other finishes (annealling, chemical or mechanical polishing).

This is illustrated, for example, by the agreement for MgO in figure 6, which was plotted

since the intermediate grain sizes shown in

table 3 were actually averages of a gradient

across the specimens. Plotting the data for the

average grain size of the different areas for

the different tests shows that the increases in

strengths for these specimens shown in table 3

are high. The general (but more uniform, as shovni in figure 6) increase in strength of perpendicular ground MgO specimens is probably

real. The narrower shape of these specimens

resulted in incomplete contact with the full

width of the wheel when ground parallel. This

may have caused differences in these specimens.

The substantially greater effect of grinding

direction on the intermediate grain size (sintered) ZrOa appears to be significant. This material is substantially weaker for its grain size

than the similar material made by hot pressing

and annealing the same raw material. Because

of the purity and density, this must result from

internal stresses or inhomogenities. (Since the

fine grain ZrO^ is made by a different process,

detailed comparison of it vdth the above two

Zr02 bodies is somewhat uncertain.) This effect

of greater decrease in strength due to perpendicular grinding in weaker bodies is also indicated in the other oxides. The finest grain size

bodies of MgO and MgAloOs. and the three finest grain size bodies of AI2O3 were also somewhat weaker for their grain size relative to

other bodies of these same materials. Thus,

part of their greater effect may be due to this

somewhat weaker character as well as their

fine grain size.

The sanding trials clearly show that the ex- tent of this directional effect is dependent on

machining conditions, but is expected to be observed to varying degrees for a variety of conditions. The reductions from sanding polycrystals are probably low because only a few fim

could be removed in a reasonable time. The re- sulting density of striations was thus low

enough so that strengths may not have been

reduced as much if the amount of material re- moved more closely approached that of grinding. On the other hand, slight uneveness of the

gas polished sapphire resulted in initial contact

only on a few areas, and a few small areas were not sanded at all. The resulting higher

local pressure is probably partially responsible

for the more severe striation and resulting

strength reduction.

4.3. Application to Other Materials and Testing

The relation to hardness (i.e., yield stress)

and grain size should make these results generally applicable to most other materials. For

glassy materials, at least those which may exhibit phase segregation [5], it appears that

Marsh's concepts of yield stress may be applicable. However, hardness will not necessarily

be a good measure of yield stress in many materials containing pores or second phases, in

which case, modification may be required.

The effect of grinding direction on strength

can have important effects on different test

techniques. Since the effect of grinding direction is a surface, and not an edge, effect, it should apply to tests that do not involve edges.

For example, in biaxial testing of plates, there

always will be a tensile component normal to

the grinding stria. Thus, comparison to flexure

testing, in which grinding will usually be parallel with the tensile axis will reflect in part

these grinding differences. Failure to recognize

possible resultant differences due to grinding

can lead to false conclusions. Similarly, solid

cylindrical rods used in uniaxial tension, and

hollow cylinders used in hoop tension tests, will

normally be ground circumferentially. The uniaxial tensile specimen will thus be ground perpendicular to the tensile axis, which can affect

their strengths. Hoop tensile tests may also be

affected, but further analysis is probably

needed.

5. Summary and Conclusions

Soda lime glass, glassy carbon, and single

crystals of MgO, TiO., MgALO^, and ALO^

were tested with surfaces as-ground parallel

439-921 O - 72 - 25 373
or perpendicular to the tensile axis. No difference in strength was found for the softest MgO.

The strengths of all other crystals ground perpendicular to the tensile axis decreased relative to the strengths of parallel ground specimens as the hardness (H) and therefore the

yield stress increases. This is attributed to the

reduction in the depth of slip from machining

(which increases as a function of 1/H), limiting deformation more to the local area around

the individual grinding particles to generate

the observed surface steps or stria. These cause

maximum stress concentration when perpendicular to the tensile axis, and hence the difference

in strengths, which can approach 50 percent in

harder materials.

Polycrystalline bodies show a similar trend,

but with an important grain size effect. Reduction of strength due to grinding perpendicular

to the tensile axis, increases with decreasing

grain size. The result is that large grain MgO

shows little reduction, or possibly an increase

in strength, as do single crystals. However, fine

grain bodies are measurably weakened. On the

other hand, while fine grain bodies of harder

materials are weakened at comparable levels to

that found in single crystals, large grain bodies are not. The effect also appears to be greater in

bodies whose strengths are generally weaker

than expected for their grain size.

Grinding of specimens by S. Slawson, and

electron microscopy by J. Robinson, are gratefully acknowledged. Significant aid from Dr. P.

Becher in orienting and testing some of the

sapphire in the latter stages of the program is

also gratefully acknowledged.

6. References

[1] Mould, R. E., and Southwick, R. D., Stren^h and

static fatigue of abraded glass under controlled

ambient conditions: II, Effect of various abrasions and the universal fatigue curve, J. Am.

Ceram. Soc. 42, 582-92 (1959).

[2] Rice, R. W., Becher, P. F., and Schmidt, W. A.,

The strength of gas polished sapphire and

rutile, this volume.

[3] Hirthe, W. M., and Brittain, J. 0., Dislocations in

rutile as revealed by the etch-pit technique, J. Am. Ceram. Soc. 45, 546-54, (1962).

[4] Rice, R. W., Machining and surface workhardening of MgO, submitted for publication.

[5] Rice, R. W., The compressive strengths of ce- ramics, to be presented at the conference on Ceramics in Severe Environments, Raleigh,

N.C., Dec. 1970. Proceedings to be published.

[6] Matkin, D. I., and Caffyn, J. E., The surface

hardness of sodium chloride containing divalent

impurities. Trans. Brit. Ceram. Soc. 62, 753-61,

(1963).

[7] Groves, G. W., and Fine, M. E., Solid solution and

precipitation hardening in MgO-Fe-0 alloys, J. Appe. Phys. 35, 3587-93, (1964).

[8] Brace, W. F., Behavior of rock salt, limestone,

and anhydrite during indentation, J. Geophy.

Res. 65, 1773-88, 1960.

[9] Hockey, B. J., Observations on mechanically

abraded aluminum oxide crystals by transmission electron microscopy, this volume.

[10] Marsh, D. M., Stress concentrations at crystal

surfaces and the embrittlement of sodium chloride, Phil. Mag. 5, 1197-99, (1960).

[11] Rice, R. W., The effect of sputtering on surface

topography and strengths of ceramics, this

volume.

[12] Marsh, D. M., Plastic flow in glass, Proc. Roy.

Soc. A-269, 420-35, (1964).

[13] Hockey, B. J., private communication.

[14] Klassen-Nekludova, M. V., and Tomilovskii, G. E.,

Corundum crystal strength in bending and compression tests as a function of crystallographic

orientation, Tr. In-ta Kyestallogr. An, SSSR,

No. 8: 5-12 (1953).

[15] Brace, W. F., Dependence of fracture strength of

rocks on grain size, Penn. State Univ. Mineral

Expt. Sta. Bull. 76, 99-103, (1961).

7. Appendix. Surface Slip in Rutile

during Grinding

TiOj crystal bars of the orientation described in the

preceding paper were ground both as described there,

and at somewhat lower removal rates. Fractures which

were roughly normal to the ground surfaces were

etched in boiling, concentrated H2SO4 for about 20

minutes after Hirthe and Brittain [3]. A variety of

behavior was observed. Some specimens showed a thin

layer (i.e. about 5 fim) of etch pits along parts of the

surface, especially near depressions into the surface,

which also frequently had short slip bands extending

from some area, as shown in figure Al. Other speci- mens showed a much more uniform and thicker (e.g., to 80 yum) layer, as shown in figure A2. Annealing in

air at temperatures of the order 1200 °C appeared to reduce the density of dislocations and result in limited

polygonization, as shown in figure A3. The depth of this was in the range of slip observed between figs. 1 and 2. The average depth of slip in all specimens was of the

order of 15 /xm. It is also interesting to note that there is a high

density of dislocations which extend considerably

M

50;im

Figure Al. TiOt crystal showing ir- regular depth of surface slip from

grinding. Etched fracture approximately normal to the ground (righthand) surface.

374
deeper around the origin of fracture, figure 2. This ture origins. The depths of these and their association was observed in several cases, and sometimes lines of with fracture indicated that they were associated with

pits suggesting slip bands were observed at such frac- the fracture process.

200 /im

i A

Figure A2. TiOt crystal showing uniform depth of surface slip from grinding. Etched fracture approximately normal to the ground (right-hand)

surface.

Figure A3. Ground TiOi crystal after annealing in air at about 1200 °C. Etched fracture approximately

normal to the originally ground (right-hand) surface. Note sub-boundaries, some of which may have

resulted from annealing the ground surface.

375

439-921 O - 72 - 26
Discussion

Donovan: What were your units of hardness?

Rice: Kilograms per square millimeter.

Donovan: Well, how did you measure the

hardness ? Rice: I took typical hardness values from

the literature. I wasn't really concerned about

making a differentiation between Knoop or Vickers hardness. Vickers hardness would

normally be a little higher, say like 10 percent

or so, but I was just taking average figures

from the literature.

Donovan: I see, you didn't measure right

on the specimen.

Rice: No, we haven't been able to do that yet.

Gielisse: How much of this is due to the

anisotropy in hardness as we know it to exist?

Rice: The anisotropy in hardness is much

much less than the effect we saw here. I'd also

like to mention that the effect of different

orientations of sapphire seem to correspond

very nicely to some work of Bernard Hockey

in that all except for the 0° orientation we

didn't see much of an effect. Bernard observes

twinning in all those orientations except the

0° one. Therefore, it suggests that this mode

of plastic deformation reduces the amount of

difference in the strength between the two

different grinding directions. However, when

you don't have significant twinning on the basal

plane you get a significant effect.

Heuer: I agree with the point that Peter

Gielisse makes. The essential point is that defects due to grinding damage that can cause

fracture can have stress concentration ratios

greater than the hardness anisotropy ratios.

Rice: That may be true. My comment was

that the difference in strength with different

grinding direction was generally greater than

the hardness anisotropy. We also see significant

effects in glass, which should be isotropic.

Thus, while hardness anisotropy may be a factor, it cannot explain the complete effect. The

correlation is still poor if we narrow the consideration to only sapphire. The effect appears

to be greater on the basal plane than the anisotropy of hardness. However, there is little or no effect on other planes where hardness

anisotropy would be expected.

national bureau of standards special pub. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).

376
The Influences of Material Removal on the Strength

and Surface of an Alumina

H. S. Starrett

Southern Research Institute, Birmingham, Alabama 35205

Southern Research is conducting- a quantitative investigation of test methods used

to evaluate brittle materials. The material for this study is a high purity, high density

alumina manufactured by Coors Porcelain Company of Golden, Colorado. The primary

scope of the total program is a study of test methods, scaling laws and surface

finish effects. The scope has been broadened to include the effects of property interactions on strength and the relationships between micro and macro characterization

and the mechanical properties.

Initial tensile and flexural data indicated that surface finish had no effect on

the tensile strength of this particular alumina. Since this was contrary to generally

accepted results, considerable effort has been expended in investigating this phenomenon. It has been postulated that the lack of an effect may have been due to subsurface damage that occurred during the grinding operations.

The investigations into surface/subsurface damage have focused into four principal areas. Microstructural characterization has been used to characterize the

surfaces of interest, to look for evidence of damage and to look for microstructural

events which may be normalizing the data. A study of fracture source distribution

has been made in order to determine the depth of damage which would have a significant effect on the flexural data. Surfaces have been and are being subjected to different treatments such as refiring in modified environments and deep lapping in

order to define the nature of the surface/subsurface damage, if it exists.

The microstructural work has not revealed any obvious damage; however, it was found that even by proceeding from a shop grind to a metallurgically lapped

surface certain features such as exposed pores, interfaces between alumina matrix

and second phase and evidence of the original ground surface did not totally disappear.

The study of the fracture source distribution provided quantitative evidence that

damage on the order of 0.005 in (approximately five maximum grain diameters)

would have a significant effect on the flexural data. Results from the surface treat- ments, deep lapping and refiring show that these different preparations did not affect

the flexural strength of the material. Other indirect evidence (statistical data) show

this material to be insensitive to surface effects within the range of surface finishes

investigated.

Key words: Alumina; surface finish; surface treatments; Weibull.

1. Introduction

Recent work on the effects of surface finish

on the strength of a ceramic yielded results

which were not consistent with the results

normally presented in the literature. This study

showed that the flexural strength of an alumina

was insensitive to surface finish effects within

the range of surface finishes and conditions

studied.

The material under investigation was a high

purity alumina manufactured by Coors Porcelain Company of Golden, Colorado. The material had the following microstructural characteristics.

1. Grain size — maximum grain diameter

was 25 microns and the average grain intercept value was 3.6 microns.

2. Porosity — maximum pore diameter was

25 to 50 microns and the pore volume was 3-5

percent.

3. Secondary microconstituent—a second

phase occurring as discrete grains and representing nominally 2 percent by volume was observed. In addition to aluminum, these grains

contained the cations sodium, calcium, magnesium and silicon.

4. Fracture mode for tensile and flexural

specimens—electron fractography indicated

that approximately 90 percent of the fracture

was intergranular. No intergranular phase was

noted. The material was evaluated to determine

whether or not it was uniform and reproducible

with respect to mechanical strength and physical properties, and thereby suitable to be used

for a comprehensive test methods study.

The outcome of preliminary tests for surface

finish effects are shown in table 1. These re- sults show that except for the as-fired surface,

the surface finish did not affect the strength

of the specimens. The ground surfaces, polished surfaces and metallurgically lapped sur- faces gave essentially the same strength values

in both flexure and tension. The nominal flexural strength was about 49,000 psi (337.8 x

lO*' N/m-). The pressed and fired surface gave

a 15 percent lower strength value of 40,820 psi

(281.4 X 10*= N/m2). This lower value can be

partially attributed to the fact that the flexural specimen could not be machined to the tol377
Table 1. Results of surface finish study on macro specimens

Surface

condition

Remarks

Flexure

psi

(10" N/m')

Tension

psi

(10' N/m=)

Pressed and fired

(150-175 rms)

Pressed, green

machined, and

fired

Ground surface

(15 rms)

(shop ground)

Polished surface

(3-4 rms)

(shop polish)

Polished surface

(lapped)

(metallurgically

lapped)

Blank 3A-10-088

No data

Blank 3A-10-087

Blank 3A-10-088

Blank 3A-10-088

Blank 3A-10-088

One-in (2.5 cm) specimens taken from

two-in (5.0 cm) specimen. Alternate

ends were evaluated as ground.

MOR = 51,930 psi

(358 X 10" N/m=)

40,820

(281.4)

47,890

(330.1)

50,620

(349.0)

48,420

(333.8)

50,820

(350.4)

48,560

(334.8)

46,220

(318.6)

erances used on the other specimens without

machining the fired surfaces.

These observations can be rationalized in

either of two general ways. The first would be

to assume that the handling, machining or sur- face preparation did not affect the material at

the surface or slightly below the surface. Then

the results could be explained from one of

several points of view. One would be that even

though the "better" specimens were metallurgically lapped, sufficient material was not

removed from the surface to do away with

damage such as rough places and geometric

discontinuities which create stress concentrations. Another interpretation might be that

after a surface has become sufficiently "good",

new flaws are exposed as material is removed,

or as a third alternative, the fractures are

initiating internally. The second approach

would be to assume that during the course of

handling, etc., the surface of the material and/

or the subsurface regions were mechanically

damaged. The presence of damage would then

offset any changes resulting from surface preparation and normalize the data.

Since the initial results were somewhat un- expected and could bear directly on later analyses made regarding the material's behavior,

a more complete understanding was sought. To

this end an analytical study of the fracture source distribution was made to determine the

depth of damage which would have a significant effect on the flexural data, surfaces were

subjected to several secondary finishing techniques in an attempt to define the nature of

any surface/subsurface damage, and microstructural characterization was used to characterize the surfaces of interest and to look

for evidences of damage or microstructural

events which could be normalizing the data.

In order to define the fracture source distribution the Weibull Distribution function was

used in the following form:

where

S = survival probability

o- = tensile stress of arbitrary spatial distribution

(To = constant

m = Weibull Modulus

V = volume subjected to tensile stress

which may be transformed to

S = exp [-^'"T"']

where

^ = normalized stress, a reference ultimate

stress divided by the average

r = Gamma Function, r( 1 + —

)

\ m/

If the sample average 5 is accepted

(13 = a/a) the transformed form of the distribution shows that there is only one parameter left to define the distribution, the Weibull

Modulus m. Intuitively it would be expected

that the value of m based on a surface effect

would be different from one based on a volume

effect.

The question of probability of fracture with378
in some specified region of a test specimen re- lates the relative size of the region in question

to the total specimen volume. For a simple ex- ample, the probability of fracture within a given region of a specimen subject to uniform

tension is found by the ratio of the given re- gion's volume to that of the whole specimen.

This process may be generalized for any given

stress distribution by finding the size of the

specimen having the same probability of fracture but subject to uniform tension only. The

ratios of such "equivalent" volumes represent

the relative frequency of fracture expected to occur in the respective volumes.

Beginning with the Weibull Distribution

S = exp

m

dV

this can be altered by choosing normalized variables describing the stress distribution and

the volume.

a = or f(|)

dV = CVr gU) <U

where

(TT = reference stress, usually maximum

tension

Vt = volume in tension

C = constant

$ = normalized position variable

thus

S = exp w) ^^^^

^1 This integral is a function of the Weibull Modulus m and the limits of integration |i and $2

so that

S = exp

m

(^) Vt Gii^, ^2, m)

This equation represents a portion of a specimen whose overall distribution integral Gt is

given by the total limits as G{u, v, m). Using

(TT as a convenient reference stress, the volume

subject to a uniform tension or having the same

probability of failure as the whole specimen

is given by

S = exp GrVr

= exp -fe)

For a subsidiary portion within the limits

(^,, ^2) the same process yields

Ss = exp

= exp

-if)

m

0-0/

The probability, F^, that fracture will initiate

in the subsidiary volume is simply the ratio of

its "equivalent" volume Veq^ to the equivalent

volume of the whole specimen Veq^ or

G(u, V, m)

For a rectangular flexural specimen the dimensionless functions are

where | is the dimensionless transverse distance from the neutral plane and the volume

under tension is

dV = Vt d^

G^„ $2, m) = J

Putting in limits for the whole specimen

0 < ^ < 1 gives

Gt (0, 1, w) ^ 0

m+l

and for a subsidiary portion 0 < ^ < ^

Gs = (0, i, m) = J (U =

0

m+ l

Now the probability of fracture initiating between the neutral axis and the fraction ^ of the

beam half-height is

F = m+l

Gt m+l 1 = e

Figure 1 shows the fracture source distribution in a rectangular bar under pure bending

for various values of m. Figure 2 shows the

same curve for a round flexural specimen.

379
^ 0.2

y /

/

f 1 /

m=l ^

« 7^

/ r

^

^^^^^^^ ^ m=11.5

J '^m=12.5

'o 0.2 0.4 0.6 0.8 1.0

Fractional Distance from Neutral Axia - C

L. _1_

0 0.010 0.020 0.030 0.040 0.050

(0) (0.0254) (0.0508) (0.0162) (0.1016) (0.1270)

Corresponding Dimensions for the Macrof lexural Specimens - Inches (cm)

Figure 1. Fracture-source distribution in pure bending

for a rectangular specimen.

0 0.8

> 0.4 4

it'

m=2 / //

//m=10y

—

0 0.2 0.4 0.6 0.8 1.0

Fractional Distance from Neutral Axis - ^

Figure 2. Fracture-source distribution in pure bending

(round specimen).

For a specimen subjected to uniform tension

(7 = or, /(I) = 1

so that

dV = 2Vr idi

.$2

For the whole specimen

Gt (0, 1, m) = 2 r m = 2-

J Q = 1

and for a subsidiary portion

Gs (0, ^, m) = 2 f m =

0

and the probability of fracture is given by

Gt

.040 .030 .020 .010 0 .010 .020 .030 .040

(. 1016) (. 0508) (0) (. 1016) (.0508)

Distance from Centerline - Inches (cm)

Figure 3. Fracture-source distribution in pure tension

for a round tensile specimen.

1.0

0 0.6

S 0.4

4J

ID H

U

0.2

UNIFO W TENSION

R

F

S

ECTANGULAR

LEXURAL

PECIMEN

0.2 0.4 0.6 1.0

Fractional Distance from Neutral Axis (Flexure) or Centerline (Tension) - 5

Figure 4. Fracture-source distribution for a uniform

tensile specimen and a rectangular flexural specimen

(pure bending) for a material with a weibull modulus

of 12.5.

380
which is independent of m. Figure 3 shows the

curve for a uniform tensile specimen. Figure

4 shows the fracture source distribution for the

tensile and rectangular flexural specimen for a

material where m = 12.5, From this plot it can

be seen that over 70 percent of the fractures

should initiate in 10 percent of the half-height

and 95 percent should occur in 20 percent of

the half-height. For the specimen used in this

study, where the half-height h = 0.050 in,

(0.1270 cm), 70 percent of the fractures should

initiate within 0.005 in (0.0127 cm) of the

tensile surface. This represents about 35 aver- age grain diameters or 4-5 maximum grain diameters. Thus the existence of surface/subsurface damage or a change in the nature of the

material within the first 0.005 in (0.0127 cm)

of the surface should have a pronounced effect

on the data.

With these results in mind, several secondary surface finishing techniques were employed

on the flexural specimens (4 point loading) in

order to gain insight into surface or subsurface damage. These techniques were not intended to change the nature of the material

or to provide additional strength, but rather

were for the purpose of determining whether

or not surface/subsurface damage existed, and

if so, to what extent. The main surfaces con- sidered were:

1. Pressed, fired

2. Pressed, green machined, fired

3. Shop grind

4. Shop grind, shop polish

5. Shop grind, metallurgical polish (lap)

Figure 5. Electron photomicrograph—as-received surface, pressed and fired. Fiducial bar equals 5.0 microns.

6. Shop grind, Hg refire

7. Shop grind, lap, Ha refire

8. Shop grind, air refire

Figures 5 through 12 are electron photomicrographs of these surfaces. Also evaluations were

performed on chemically etched surfaces. These

included specimens etched in hydrofluoric acid

Figure 6. Electron photomicrograph—as-received surface, pressed, green machined and fired. Fiducial bar

equals 5.0 microns.

Figure 7. Electron photomicrograph—15 rms surface

created by standard shop surface grinding. Fiducial

bar equals 5.0 microns.

381
Figure 8. Standard shop grind plus lapping with 15, 7, 1 and % micron diamond.

Figure 9. Electron photomicrograph—<i rms surface

developed using metallurgical laboratory lapping

techniques. Fiducial bar equals 5.0 microns.

and specimens immersed in Borax at 1500 °F

( -840 °C).

The results of the flexural evaluations on the

specimens which had undergone the several

different surface treatments are shov^n in table

2. The standard shop ground surface v^^as used

as the baseline value. Normally, however, only

from 3-10 specimens were employed to test a

particular condition so that comparisons were

Figure 10. Electron photomicrograph—Surface produced by refiring a 15 rms ground surface in hydrogen. Fiducial bar equals 5.0 microns.

Figure 11. Electron photomicrograph—Surface produced by refiring a 1 rms lapped surface in hydrogen.

Fiducial bar equals 5.0 microns.

made against as-ground specimens taken from

the same blank of material as the treated specimens. The values used for comparison are separated in the table by the virgule. Also there

was a problem in providing enough material

for all the evaluations shown here plus other

evaluations required by the continuing program.

382
In order to have sufficient stock, material below

the standards established by the majority of

the blanks has been used for some of the tests. This material has been noted as A-11*. The

average flexural strength of specimens taken

from Blanks A-11* v^^as 37,130 psi (256.0 X

10" N/m^) compared to the overall average of

48,300 psi (333.0 x 10« N/m^), Also specimens

have been evaluated which have been precision

ground, and some have been evaluated after

only being sliced from the stock. These are also

noted in the table.

From table 2 it can be seen there were no

significant changes in the strength of the material as a result of the treatments. Basically

the material remained unaffected.

The microstructural work has not revealed

any obvious damage; however, it was found

that even by proceeding from shop grind to a

metallurgically lapped surface certain features

such as exposed pores, interfaces between alumina matrix and second phase and evidence of

the original ground surface did not disappear,

see figure 9.

Table 2. Effects of surface treatment on the flexural strength of an alumina

Surface condition

Flexural strength

psi (lO'N/m^) Remarks

1. Pressed, fired 40,820

(281.4)

Surface rough and pimpled

2. Pressed, green machined, fired 43,240

(298.1)

3. Shop grind 48,300

333.0)

314 specimens

4. Shop grind, shop polish 48,420

(333.8)

3-4 runs

5. Shop grind, metallurgical polish (lap) 50,820/51,930

(350.4/358.0)

One-in (2.5 cm) specimens,

matched pairs (3)

6. Shop grind, hydrogen refire 45,820/47,710

(315.9/328.9)

7. Shop grind, lap, hydrogen refire 43,790/47,710

(801.9/328.9)

8. Shop grind, air refire 44,450/47,710

(306.4/328.9)

9. Sliced material/good grinding 43,980/50,680

(303.2/349.4)

For premium material

10. HF etch for 0.005 in (0.0127 cm), 0.004

in (0.0102 cm) removed by machining

39,800/37,130

(274.4/256.0)

Sliced material; A-11*

11. Immersed in Borax 2 minutes at 1550 °F (—840 "O

45,300/50,680

(312.3/349.4)

Three specimens

12. Baked at 950 °C one hr, placed in

vacuum, 20 hrs at 10"^ torr, 1 hr at 10"' torr, broken at 10"'

34,950/37,130

(240.9/256.0)

Sliced material; 3 specimens, A-11*

383
2. Discussion

Based on the surface finish effects data it appears that the alumina has not been affected

adversely by the grinding or handling. Of

course, it is possible that all of the treatments

were ineffective in removing or repairing sur- face/subsurface damage that may have oc- curred. To date we have been unable to determine directly whether or not damage

existed, consequently, we have been unable to

determine whether or not we have been able to

remove the damage. However, the indirect evidence indicates there was not a surface/subsurface damage effect. This judgment was based

on the following observations:

1. The various surface treatments have not

affected the flexural strength of the material.

2. The flexural (and tensile) data showed

trends and effects caused by other factors.

3. The data (tensile and flexural) showed

trends that were consistent with the Weibull

calculations based on volume. Observation 1

was discussed earlier and will not be repeated

here.

Recall that the analysis of the fracture

source distribution revealed 70 percent of the

fractures should initiate within 0.005 in (0.0127

cm) of the tensile surface. Therefore any significant damage which extended to this depth

should have a pronounced effect on the flexural

strength. In fact, it is reasonable to expect that

such an effect would easily override other effects or trends in the data. Now consider figure

13 which is a plot of flexural strength versus minimum fired thickness (the minimum dimension of the fired part from which the flexural

specimens were taken). Here it is easily seen

that there is a definite correlation between the

flexural strength and fired thickness. Figures

14 and 15 show other correlations that were

obtained using flexural data. The existence of

these trends as exhibited by the flexural data

suggests that a material response was observed

that had not been masked or normalized by

some other dominant event.

Along these same lines consider the results

shown in table 3. This is a table summarizing

the results of statistical calculations which were

performed to determine whether or not correlations existed between some of the more easily

obtained material processing parameters and

mechanical properties which can be measured

in the laboratory. Both tensile and flexural data

were employed. In this table there are nine

450

400 -

5

01

350-

ui 300

250

]

ji

-4

( — ) ( )

( )

)

>

I.

0.2 0.4 0.6 0.8 1.0 1.2 Minimum Fired Thickness in Inches1. 4 1.5 1. 8

—I

0.01 0.02 0.03

Minimum Fired Thickness in Meters

0.04 0 .05

Figure 13. Average flexural strengths versus minirnvm fired thickness.

384
Figure 14. Dimensionless stress versus porosity for

alumina showing regressive curve.

Figure 15. Dimensionless stress versus grain size for

alumina showing regression curve.

direct comparisons between the tensile and

flexural strength data. In seven of the nine

comparisons the tensile and flexural data ex- hibited the same trends ; in the other two cases,

the data were inconclusive. Thus the flexural

specimens, which are surface sensitive, are showing the same trends as the tensile specimens, which are less surface sensitive. This

adds confidence to the earlier statement that

the flexural data are revealing material re- sponses and are not being normalized by sur- face effects or subsurface damage.

The third observation was based on calculations made by assuming the material responded

in a manner consistent with Weibull's statistical theory of strength. Figures 16 and 17 show

some probability of fracture curves for this

material. Note that the value of the WeibuU

Modulus m was about 12.5 for both the tensile

and flexural data. The average flexural strength

for as-ground specimens was 48,300 psi (333.0

X 10" N/m^). Using the equations developed

earlier one could predict the tensile strength

that should be obtained with the macrotensile

specimens: ^ _,l/m

tensile strength

385
Table 3. Rank correlation tests

Number Correlation indicated

Test

Items tested

of Confidence

no. samples Z Indication level Positive Negative too weak

RC 1 Green density vs sonic velocity

(a) Tensile

(b) Flexural

RC

RC

2 Green density vs cone angle

3 Green density vs tensile strength

(a) All blanks

(b) Less 11, 12, 13

(c) Blanks 11, 12, 13

RC 4 Green density vs flexural strength

(a) All blanks

(b) Less 11, 12, 13

(c) Blanks 11, 12, 13

RC

RC

RC

RC

5 Green density vs fired density

6 Cone angle vs sonic velocity

(a) Tensile

(b) Flexural

7 Cone angle vs tensile strength

8 Cone angle vs flexural strength

RC 9 Cone angle vs fired density

(a) Tensile

(b) Flexural

RC 10 Sonic velocity vs tensile strength

RC 11 Sonic velocity vs flexural strength

RC 12 Sonic velocity vs fired density

(a) Tensile

(b) Flexural

42

158

45

98

55

43

242

188

54

37

73

191

122

293

130

293

71

193

73

193

-1.801

-6.273

-0.598

3.089

0.354

3.271

3.431

0.504

2.844

-1.628

1.095

-0.520

-4.219

-8.560

-0.290

18.691

0.855

1.653

0.548

1.680

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

X

92.82%

>99.99%

99.8%

99.9%,

99.5%

99.5%

>99.99%

>99.99%

>99.99%

90.16

90.70

RC 13 Sonic velocity vs minimum thickness

(a) Tensile

(b) Flexural

RC 14 Fired density vs tensile strength

(a) All blanks

(b) Less 11, 12, 13

(c) Blanks 11, 12, 13

RC 15 Fired density vs flexural strength

(a) All blanks

(b) Less 11, 12, 13

(c) Blanks 11, 12, 13

RC 16 Minimum thickness vs tensile strength

RC 17 Minimum thickness vs flexural strength

RC 18 Minim'om thickness vs fired density

73

182

141

94

47

315

236

79

141

304

64

1.043

1.416

4.085 X

-0.017

-0.822

8.681

2.421

0.907

-6.748

-13.530

-3.727

X

X

X

X

X

X

X

X

X

>99.99%

>99.99%

98.4%

>99.99%

>99.99%

>99.98%

= flexural strength

Vt = tensile volume in a tensile specimen

Vp = tensile volume in a flexural specimen

m = Weibull Modulus

For the macroflexural specimen, figure 18

ap = 48,300 psi (333.0 x lO'' N/m^)

Vp = 0.0075 in^ (0.1229 cm«)

For the macrotensile specimen, figure 18

Vr = 0.0013 in^ (0.0213 cm^)

so that a-T calculated to be 45,100 psi (310.9 x

10'' N/m-). The measured value was 46,300 psi

(319.2 X 10« N/m==). Figure 19 shows how the

flexural and tensile data plots on a Weibull

strength-volume curve. It can be seen here that

the prediction based on volume is quite good.

If the calculations were based on surface area

the prediction would require that the tensile

specimen be stronger since the flexural specimens had nominally 2-3 times the surface area

of the tensile specimens. Thus, this reinforces

386
-J 1 1 i

1

1

I

s.

1 — — !

• Where

S - Probability oJ Tracture

r - Guvna Function «

tenBior. ""'1 4^

B - Horrnali zed vezage o;

-.^

fj _

!

1

:/

:

- -

1

'•453

1 a& determined from the Etanda rfl deviation estimator

i

1

1 1

i 1

i 1

1

i

i

Beta

Figure 16. Probability of fracture versus normalized

stress p for the macrotensile specimens.

1

:

t

1

!

t

i

!

L - -i

5 - 1-.-'"''

i /

i

!

r - Gansna JMnction

\ - Honnaliaed Stres^ maximum tension

„

/

i

!

f

' i

1

/

j

1 / tit = i; ?-977 ned from the Etanda

1 rd deviation cstimat

1 /

i

-

0.0 0.1 6.2 0.3 0.4 0.5 0.6 0,7 0.8 0.9 l.O 1.1 1,2 1.3 1,4 l.S 1.6 1.7 1.6 1.9 - J.O

Figure 17. Probability of fracture versus normalized

stress /3 for the macroflexural specimens.

the suggestions that surface/subsurface were

not controlling the mechanical property data.

3. References

[1] Pears, C. D., Starrett, H. S., Bickelhaupt, R. E., et al., A Quantitative Evaluation of Test Methods for Brittle Materials, Technical Report

AFML-TR-69-244, Part I, March, (1970).

[2] Pears, C. D., Starrett, H. S., Brown, J. R., et al., A Quantitative Evaluation of Test Methods for

Brittle Materials, Technical Report AFML-TR70-169.

[3] Robinson, E., Some Problems in the Estimation and

Application of Weibull Statistics, Technical Report UCRL-70555, Sept. 1, 1967.

387
0. 125"

(0. 317 cm) r(2.54

cm R)

r—Groove 3 mils (0.0076 cm) deep x \ 3 mils (0.0076 cm) wide - Equally

\ spaced on each end

T T

—0. 188"

(0. 478 cm)

(5.08 cm)

Macro Tensile Specimen

2 0 RMS

(5.08 cm) Macro Flexural Specimen

_>J L<- 0 . 200" (0. 508 cm)

1 l_i

tSSS) 0.100" (0. 254 cm)

Figure 18. Macrospecimens used for the evaluation of

uniformity and reproducibility.

1.4

1.2 =

1.0

0.8

Io

•HS 0.6

C

0.4

0.2

0.0

Macroflexural Specimens

Strength Ratio versus Volume

1 \ m

= 10"' in?, (6.1 X lO'^cm')

Expected C.O.V.

M=12 Flexural

10 —

(6.1 X 10"«)

10-'

(6. 1 X 10-5)

10"^

(6.1 X 10-")

10-1

(6.1 X 10-3)

10° (6.1 X 10-2)

10' ., (6.1 X 10 ')

(cm')

Figure 19. Strength ratio versus tensile volume for

tensile and flexural (J^-Point) specimens.

388
Discussion

Lange: Refiring can change the properties

of a ceramic. Did you measure grain size both

before and after refiring?

Starrett: We measured grain size before

firing. We have not completed the grain size

measurements after firing. We think the slight

decrease might be due to a slight increase in

grain size, however, the refire temperature was

selected so that grain growth would not occur. (Measurements made at a later date confirmed

that grain growth did not occur during the re- firing cycles.)

Rice: Comments—(1) The fact that you

have a wide grain size distribution and in particular, have large grains scattered throughout

a fine grain matrix would certainly indicate

that there is a volume effect in fracture. This

is because the large grains, regardless of

whether it's a flaw or a microplastic mechanism

of failure, would be the source of failure. This

is one good explanation why you get a volumetric effect. (2) I've used some of your data

as well as quite a bit of my own and analysis

suggests that the strength controlling effect in

alumina is microplasticity and quite probably

twinning. This would go along with the fact

that you don't see much effect of a wide variety

of surface finishes.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing

n/ St^nJ"^^' 1, u . MR« NT A^'' American Ceramic Society, the Office of Naval Research, and the National Bureau of Standards, held at NBS Nov. 2-A, 1970, Gaithersburg, Md. (Issued May 1972).

389

The Effect of Grinding Variables on the Strength and

Surface Finish of Alumina

Rudolf Sedlacek, F. A. Halden, and P. J. Jorgensen ^

Stanford Research Institute

Menlo Park, California 94025

The tensile strength of alumina is used as a criterion for the evaluation of the

effects of variables encountered in diamond grinding of ceramics. The variables studied

include the grit size of diamond, rate of material removal, and sparkout. The difference

in performance of natural and synthetic diamond is shown, and the feasibility of dry

grinding is discussed. The ground surfaces w^ere examined using profilometry, and

transmission electron and scanning electron microscopy. The strengths of test speci- mens having ground or as-fired surfaces are compared.

Key words : Alumina ; diamond grinding ; surface finish ; tensile strength.

1 . Introduction

In recent years, oxide ceramics, which have

certain favorable mechanical and physical properties, have been gaining wider acceptance in

various applications. Some of the properties

contributing to this acceptance are: superior

strength retention at high temperatures, chemical inertness, dimensional stability, hardness,

and dielectric properties. Unfortunately, ceramics are also brittle, which presents a serious deterrent to their use in structural applications.

However, the requirements of modern technology are such that at times the designer can

find no suitable alternative to ceramics. In these

critical applications, the ceramic components

are frequently machined to very close toler- ances. The shaping of ceramics can take place

in various steps of the manufacturing process,

depending on the dimensional accuracy re- quired. For instance, a piece can be either

green-machined and then fired, or bisque-fired,

machined, and then sintered to final density.

However, in the vast majority of cases when

high dimensional precision is required, the

ceramic piece is first sintered to ultimate density and then finished to size. Because of the

hardness of ceramics, the number of machining

methods is limited, and grinding with bonded

diamond constitutes the most widely used industrial method of material removal.

Like many other steps in ceramic manufacture, today's grinding process is the result of a highly pragmatic development and as such does

not have a particularly solid theoretical foundation. As a matter of fact, very little is known

about the physical principles of diamond grinding of ceramics, and most of the existing knowledge is considered proprietary by industry.

One of the better known facts about grinding

is that it controls the quality of the surface of

* Senior Research Ceramist, Associate Director Materials Laboratory, Manager Ceramics Group, respectively.

the ceramic and thereby influences its mechanical properties, which are sensitive to surface

conditions.

The method of material removal described

in the present article was internal cylindrical

grinding. This choice was dictated by the test

method used for strength measurements, i.e.,

the SRI expanded ring test method. The full

description of the test method is available in

the open literature [1]^ and will not be repeated

here. Essentially the test specimen has the

shape of a short thin-walled right cylinder in

which a tangential tensile stress is generated

by application of hydrostatic pressure to the

inner wall. Since this is the locus of maximum

tensile stress, it can be assumed that variations

of grinding parameters in internal grinding

will result in more pronounced changes in

strength. Consequently, the method of grinding

the outside walls and faces of the test specimens was not considered a variable and was

kept constant throughout this study.

2. Experimental Procedure

The material used in the study was Wesgo

Al-995 alumina. This is a nominally 99.5 percent pure alumina body, approximately 97 percent dense, having an average grain size of 18

X 10"*' m. The main reason for choosing Wesgo

Al-995 alumina for this study was that this

material has been studied extensively by vari- ous investigators. It is thus one of the most

thoroughly documented ceramic bodies insofar

as mechanical properties and the reproducibility of the manufacturing process are concerned.

The test material was purchased in the form

of ring-shaped blanks of such size that approximately 12.7 X 10-^ m (0.050 in) had to be re- moved from the internal and external diameters and the faces. The final dimensions of the

test specimens were:

2 Figures in brackets indicate the literature references at the end

of this paper.

391

439-921 O - 72 - 27
Inside diameter: 5.080 x 10-^ m (2.000 in)

Outside diameter: 5.588 x 10-^ m (2.200 in)

Height: 0.762 x 10 = m (0.300 in.)

All final dimensions were kept within ±1.3

X lO-^^ m (0.0005 in) and were checked by

means of air gages and micrometers.

2.1. Grinding Procedure

The machine used in this work was a precision universal cylindrical grinder (Tschudin,

Model 300) equipped with a linear-position

potentiometer for monitoring and reproducing table speeds, and several specialized

holding fixtures were designed for the individual grinding steps. The basic grinding procedure consisted of the following steps:

(a) The blanks were ground on both ends so that their faces were parallel and properly

spaced.

(b) Using a suitable mounting medium, the

faced blanks were arranged in stacks each con- taining seven rings for cylindrical grinding.

The details of the mounting procedure were

rather complex and have been described elsewhere [2, 3]. The cylindrical grinding of the

outside walls of the specimens always preceded

the internal grinding.

The facing and external grinding of all specimens was identifcal in order to isolate and to

better resolve the effect in strength of the variables studied in the internal grinding step. The

wheel used in facing and in external grinding

had a 0.305 m (12 in) diameter, a face width

of 2.54 X 10-2 m (1 in), and was designated

D320-N100M1/16. The following conditions

were used in facing and external grinding:

Table speed: 5.08 x 10"^ m/s (1 ft/min)

Infeed: 6.4 x 10-« m/pass (0.00025 in/pass)

Surface speed: 2.9 x 10 m/s (5,700 sfpm)

Sparkout: 180 s.

2.2. Grinding Variables

The variables in internal grinding whose effects on strength were evaluated were: grit size

of diamond, rate of material removal, type of

diamond (natural or synthetic), wet or dry

grinding, and length of sparkout time. It became apparent that these variables could not

be studied individually and that some compromise in the form of a combination of variables

would be necessary. The diamond grit size offered the largest range of conditions to study

the effect of the surface finish per se on

strength and was varied from 36 to 1200.

All wheels were of the straight type (lAl),

having an outside diameter of 2.54 x 10"= m

(1 in) and a face width of 1.27 or 0.63 x 10"= m

(1/2 in or 1/4 in). All wheels were metal bonded except the 36 grit wheel, which had a vitrified bond, and the 100 grit synthetic diamond

wheel, which was resinoid bonded. For maximum perfection, all wheels were trued against

a silicon carbide wheel mounted on the spindle

of a brake-controlled trueing device, and ro- tated by contact friction. Subsequently, the

wheels were dressed wet with silicon carbide

sticks of proper grade to remove the glaze produced by the trueing operation.

The sparkout time was limited only by consideration of time, whereas other variables

were less flexible. The study of the practicality

of dry grinding, for instance, was limited to

the resinoid-bonded wheel because it was sooii

discovered that metal-bonded wheels, although

their matrix has a much better thermal conductivity than the plastic, wear fast, load up

with grinding debris, and oxidize. The same is

probably true of wheels having a vitrified bond,

but this avenue has not been explored. In terms

of effect on strength, the most difficult variable

to assess in the grinding process was the rate

of material removal. This variable can be re- solved into two main components, i.e., the table

travel speed and the depth of infeed. Both of

these can be studied only over a limited range.

One of the reasons for this is the nature of the

grinding operation employed. Internal cylindrical grinding is undoubtedly the least suited machining operation for rapid stock removal. Unlike surface grinding in which the work piece is

fixed rigidly to the table, in cylindrical grinding the work piece is supported on only one end

and consequently is much less rigid. The grinding wheel in internal grinding is mounted on

the end of an arbor whose length must be somewhat longer than the depth of the cavity being

ground, and the arbor must be smaller in diameter than the grinding wheel itself. Consequently, this arbor is subjected to considerable

bending forces which produce vibrations. However, in the work described the limits in the

rate of stock removal were dictated primarily

by the strength of the bond between individual

rings of the stack.

The makeup of the grinding wheel, i.e., grit

size, diamond concentration, and bond type,

was found to be the most important and limiting factor in controlling the rate of stock re- moval. With the 60, 100, 120, and 320 grit

wheels, the depth of cut presented no problem

over the range investigated (2.5 X 10'^ m to

3.8 X 10-* m - 0.0001 to 0.015 in). In taking

heavier cuts (5 x 10-^ m to 12.7 x 10"* m —

0.002 to 0.005 in) the fine grit wheels (600 and

1,200 grit) tended to slow down. The leading

edges wore rapidly and the metal seemed to

oxidize. The 36 grit vitrified bond wheel was

used only for very light cuts to impart a coarse

finish to stacks previously ground with another

wheel within 2.5 x 10"^ m to 5 x 10"^ m (0.001

to 0.002 in) from the final dimensions. The

392
blocky diamond with large facets, few cutting

edges, and low diamond concentration makes

this type of wheel wholly unsuited for grinding

of dense, hard ceramics. Furthermore, the

vitrified bond lacks the toughness of the metal

bond, and consequently the wheel was broken

in an attempt to make a moderately deep cut,

12.5 X 10-=^ m (0.005 in). The surface speed

most commonly used in internal grinding on

this program was 28.4 m/s (5,600 sfpm). The

highest available speed (34.5 m/s—6,800 sfpm)

was used with the 1,200 grit wheel in the finish

grinding of a few stacks. This speed, however,

exceeds the safety limits of both the wheel and

the spindle. Since no noticeable improvement

of surface finish or strength resulted, this procedure was discontinued.

There are many possible combinations of

variables, and several of them were studied,

i.e., effect on strength of a deep final cut, effect

of a deep cut followed by several shallow cuts

made with the fine grit-size wheel with or without sparkout, etc. These experiments did not

clarify the picture; only heavy infeeds at low

table speeds with coarse grit wheels showed

conclusive evidence of damage.

3. Results and Discussion

In tabulating experimental strength data to

show the effects on strength of individual machining variables, it became apparent that only

a few of the variables employed could be used

as valid criteria for evaluation. The effect of

some variables is so small that it cannot be

detected. The length of sparkout time is a prime

example. Other variables, such as the rate of

material removal, show a considerable effect,

but in such a complicated manner, that it is not

feasible to present the results in a unified logical way. Furthermore, in cases where several

variables were changed simultaneously over a narrow range, it was impossible to decide which

one caused the variations in strength, if any

occurred.

The variable whose effect on strength can be

shown most systematically is the grit size of

the diamond used. Table 1 presents the average

strength values of specimens finish ground with

natural diamonds of the grit size indicated, re- gardless of the previous grinding history. Included are the strength values of specimens

ground by the manufacturer, which can be con- sidered indicative of results obtained in a routine production grinding process.

The relationship between the grit size and

strength is very systematic, although greater

strength differences were anticipated. In particular, lower strength values were expected

from the specimens ground with the two coarsest wheels (36 and 60 grit).

A limited study of the practicality of dry

grinding (without coolant) was made using a resinoid-bonded wheel containing 100 grit size

synthetic diamond. The results are shown in

table 2. This wheel was used without coolant

to remove the last 5.1 x lO"'^ m (0.002 in) in

wall thickness in 2.58 x 10 « m (0.0001 in)

steps and at a table speed of 5.1 x 10"^ m/s (1

ft/min). (The initial grinding was done with

the 320 grit wheel.) Five stacks were ground

in this manner, five more stacks were ground

under identical conditions except that coolant

was used. The results presented in table 2 show

that dry grinding under the experimental conditions employed does not improve strength. As

a matter of fact, dry-ground specimens are considerably weaker than those ground in the con- ventional way, i.e., 21.0 x 10^ N/m^ (30.5 x

lO'' psi) versus 22.2 x 10^ N/m^ (32.2 x 10^

psi).

It may be noted that the average tensile

strength of specimens ground wet with the

same wheel is practically identical to the

strength of the specimens of the group D in

table 1 ground with the 320 natural diamond.

Table 1. Tensile strength of alumina as a function of the grit size of natural diamond in metal-bonded grinding

wheels

Average tensile Standard

Grit Number strength deviation Coeff. of

Group size variation

specimens N/m'XlO' psi X 10' ±N/m=XlO' ±psiXlO' percent

Unknown 12 20.5 29.8 8.1 1.2 4.0 A 36 29 21.3 30.9 11.7 1.7 5.5 B 60 37 21.4 31.1 12.4 1.8 5.8 C 120 48 21.9 31.7 8.2 1.2 3.8 D 320 33 22.3 32.3 7.6 1.1 3.4 E 600 45 22.5 32.6 10.3 1.5 4.6 F 1200 85 22.9 33.2 9.7 1.4 4.2

393
An indication of superior performance of synthetic diamond has also been observed by other

workers [4] using various grades of synthetic

and natural diamond in metal-bonded v^heels.

The lov^er strength of specimens of group A in

table 2 indicates that dry grinding is not a recommendable procedure.

The effect of infeed on strength gave conclusive results only in the case of the heaviest cuts.

For example, a group of specimens were ground

with the 60 grit wheel using a table speed of

2.6 X 10-* m/s (0.05 ft/min) and a depth of

cut 3.8 X 10 " m (0.015 in). The average

strength of these specimens was 18.6 x 10'' N/m^ (27.5 X 10^ psi). This is a significantly

lower value than the strength of specimens of

group B in table 1 which were finish ground

with the 60 grit wheel under less severe infeed

conditions. Similarly specimens of another

stack were ground with the 100 grit resinoidbonded wheel under identical conditions of infeed and table speed. Their average tensile

strength was 20.5 X 10^ N/m^ (29.8 x 10^ psi).

Although this is lower than the average

strength of specimens of group B in table 2,

the difference is less pronounced (7.8%) than

in the case of specimens ground with the metalbonded 60 grit wheel (12.1%). Apparently the

bond of the wheel makes a difference, but the

nature of the difference is not well understood.

It should also be pointed out that in these ex- periments in which a heavy infeed was used,

the resinoid-bonded wheel lost approximately

one-half of the diamond bearing layer, while

the amount of wear of the metal-bonded wheel

was not measurable with the micrometer. These

two examples show clearly that heavy infeeds are detrimental to strength. However, this observation cannot be extended indiscriminately

to the concept of the rate of material removal

which represents the volume of material re- moved per unit time, and is equal to the product of infeed and table speed per unit time. In

the aforementioned cases when a heavy infeed

and slow table travel were used, this quantity

equals 9.1 X 10"" mVs (1.5 x 10"* in Vs. There

were instances in which combinations of infeed

and table speed resulted in considerably higher

rates of material removal without any detrimental effect on strength. For instance, one

stack in group C of table 1 was ground with an

infeed of 12.5 x 10-" m (0.005 in) at a table

speed of 5.1 X 10-« m/s (1 ft/min). In this case the rate of material removal was 6.1 X 10"^

mVs (1.0 X 10-3 inVs). This is 6.7 times higher

than the previous value and yet this stack with

an average strength of 22.5 x 10^ N/m^ (32.7

X 10^ psi) was the strongest of all stacks of

group C. Obviously, infeed and table speed do

not have the same effect on strength, and this

phenomenon probably depends on the rigidity

and vibrational characteristics of both the work

piece and the machine used.

4. Tensile Strength of Alumina

having As-Fired or Ground Surfaces

No matter how gently done, grinding of ce- ramics is undoubtedly a harsh method of material removal. There will always be chipping,

grooving of grains, pullouts, cracking of boundaries, etc. Therefore, it is not surprising that

alternative means of material removal are being sought, for instance, chemical machining.

To our knowledge, no practical solution to this

problem is yet in sight. The question may then

be raised whether or not any postfiring treatment is necessary, provided that the preceding

manufacturing steps are carried out in such a way that no additional dimensional adjustments

are required for the intended application of the

ceramic product. This avenue was briefly explored.

For this study, a different lot of Al-995

alumina was used and the specimens were prepared in three different ways:

Group A—Specimens were green-machined

to a size determined by the experimentally

found shrinkage factor and then fired to final

density.

Group B—Specimens were bisque-fired to

1450 °C, then machined to a size that would

shrink in firing to the desired dimensions.

Group C—Specimens were oversized blanks

which after firing were ground to final dimen^

sions.

The two groups of specimens having the as-, fired surfaces had a number of geometric imperfections caused by the manufacturing process. In the case of group A, the shrinkage

factor was miscalculated so that all specimens

turned out to be undersized on the inside diamTable 2. Tensile strength of alumina ground with and without coolant

Group

Number

of

specimens

Type

of

grinding

Average tensile

strength Standard deviation Coeff. of

variation

N/m ' X 10 ' psi X 10 percent ' ±N/m' X 10° ±psi X 10

A 31 Dry 21.0 30.5 13.8 2.0 6.6 B 34 Wet 22.2 32.2 9.7 1.4 4.3

394
eter. Furthermore, the specimens had the general shape of slightly truncated cones rather

than that of right cylinders. Specimens of

group B did not appear to be tapered but were

somewhat oversized. These specimens had an- other more serious defect, namely, crushed

alumina grit fused firmly onto large areas of

the surface, particularly on one end face. This

grit adhered only loosely to specimens of group

A, and was not found on the blanks of group C.

The wall thickness of all specimens of groups

A and B was relatively uniform, but the outside diameters appeared to vary slightly in a

completely random fashion, indicating a certain

degree of out-of-roundness. Whether this is the

result of improper machining or nonuniform

shrinkage is not known at the present time.

The comparison of tensile strength between the

three groups of specimens is presented in table

3. The difference in strength between the two

types of specimens is obvious, but we cannot

simply state that the ground specimens are stronger (21.4 x 10^ N/m^—30.8 x 10^ psi)

than the as-fired ones (19.9 x 10^ N/m^—28.8

X 10^ psi). The original intention was to make

a comparison of strength based only on the

difference between the surfaces and the means

by which they were obtained. However, the real

situation is complicated by factors stemming

from the manufacturing process which obscured the original objective. In the case of

group A, it is possible to say that the as-fired

surface has an average roughness of 1.3 to 3.8

X 10-*' m (50 to 150 /xin), but the observed

strength and scatter cannot be only the results

of surface conditions. They must also reflect

the dimensional irregularities that upset the

stress distribution in a way not amenable to

analysis.

The observed average strength of specimens

of group B is identical to that of group A, although the overall geometry of these specimens

was considerably better than that of group A,

and therefore a higher strength than in group

A could be expected. Under the circumstances,

we have no choice but to attribute the results

to the surface conditions. This surface can only

be described as a very rough interface between

alumina and air evading any meaningful characterization. However, the macroscopic roughness of the surface should not be confused with

the microscopic surface features controlling the

strength because in that case lower strength

values would have resulted.

The results do not show unambiguously

whether a ground or as-fired surface is preferable to obtain higher strength, but they strongly indicate that in the present state of the art,

grinding is necessary to make up for other

shortcomings of the manufacturing process.

5. Evaluation of Ground Surfaces

In the case of glasses, it is well documented

that in terms of strength, the surface conditions are of critical importance. It is not certain, however, to what degree this applies to

polycrystalline materials, such as alumina. Generally, it is accepted that smoother surfaces are

representative of higher strength. Whether this

is a simple function of density and severity of

surface imperfections, or a more complex

relationship involving some stress-enhanced

chemical reactivity dependent upon surface

conditions, is not known. In either case, an

interpretation and understanding of surface

characteristics is most desirable.

In the work described, two basically different

methods were employed for the identification

and interpretation of surface features, i.e., microscopy and profilometry. Although transmission electron microscopy was also used, most of

the work was done using a scanning electron

microscope. The problems involved are best il- lustrated in figure 1. These are low power scan- ning electron micrographs of Al-995 surfaces

produced by grinding wheels having different

grit sizes of diamond. By using a little imagination, it may be possible to say that the surface

produced by the 36 grit diamond appears to be

somewhat more battered than the others and

that the fewest pullouts were obtained with the

1,200 grit wheel. Beyond that, there is very lit- tle basic difference and it is virtually impossible to decide which grit size produced a partic^

ular finish.

To relate the appearance of the surface to

strength would be almost pure guesswork. From

a practical standpoint, it is nearly impossible to

examine the entire specimen surface before

Table 3. Tensile strength of alumina having as-fired or ground surfaces

Average tensile ^ -p _ strength Standard deviation Coeit. of

Group variation

ZV/m'xlO' psi X 10' ±iV/m'xlO' ±psixlO' percent

A 19.9 28.8 20.0 3.0 10.4

B 19.9 28.8 13.1 1.9 6.4 C 21.2 30.8 6.9 1.0 3.2

395
36 GRIT DIAMOND

SURFACE FINISH:

40-84 /iin. AA

60 GRIT DIAMOND

SURFACE FINISH:

32-84 jUin- AA

120 GRIT DIAMOND

SURFACE FINISH:

40-64 jUin. AA

320 GRIT DIAMOND 600 GRIT DIAMOND 1200 GRIT DIAMOND

SURFACE FINISH: SURFACE FINISH: SURFACE FINISH:

32-48 [Jiin. AA 20-28 Mi"- AA 16-24 jLtin. AA

Figure 1. Scanning electron micrographs of Al-995

Alumina.

testing to find the most damaged spot, e.g., the

highest concentration of pullouts, and to see if

it is the locus of failure. To examine fragments

of broken specimens is a futile exercise in microscopy which may, at the best, lead only to a

subjective interpretation of the studied surfaces.

Profilometric studies were made using a

Clevite Corporation Surfanalyzer 150 System.

The diamond stylus of this instrument has a

2.54 X 10-« m (0.0001 in) tip radius and all

measurements were made using a probe travel

speed of 2.54 x 10-* m/s (0.01 in/s) and a

roughness width cutoff of 7.6 X 10^ m (0.03

in).

Profilometry has certain advantages over

visual evaluation of photographs. It gives a picture of a surface that is more tangible and

easier to interpret, particularly in terms of

depth and spacing of cracks and scratches.

Furthermore, it is the most widely used method

of graphic representation of surface conditions

and it is quite suitable for comparison if equivalent profilometers are used.

The most obvious shortcomings of this method are the fact that the profilometer samples

the surface only in one line and gives no indication of conditions of the surface on either

side of this line. Also, the diameter of the stylus tip limits the depth to which the tip can

penetrate.

Figure 2 shows profile traces of Al-995 alumina surfaces ground by different grits of diamond. Each vertical division represents 5.1 x

10"^ m (20 /xin) and every horizontal division is

2.54 X 10-5 ^

(0.001 in). Subscripts indicate

the arithmetic average roughness (AA) as generated by the profilometer on a different graph

(not shown). It will be noticed that the AA

roughness values are somewhat deceiving. For

instance, the AA roughness value in the last

graph is 5.1 to 7.6 X 10-« m (2-3 /xin). However, there is a 6.6 x 10-" m (260 /lin) deep

crack, whose existence cannot be inferred from

the AA roughness value. The disproportion between total roughness and AA roughness can

be found in all of the graphs, and it is a characteristic of the averaging circuitry of the apparatus. Notwithstanding this shortcoming, we

believe that profilometer traces convey better

the qualitative differences between surfaces

than micrographs. For instance, the AA roughness value for surfaces shown in figure 2 vary

by a factor of 4 to 1 in order of decreasing

diamond grit size. No such sharp distinction

can be obtained from an inspection of the micrographs.

Whether or not these profilometer traces are

typical of the surfaces studied is difficult to decide, because the recordings vary considerably

from specimen to specimen, and also from place

to place on the same piece. Furthermore, it is

doubtful that the appearance of the most typical surface feature can lead to an understanding of the relationship between surface conditions and strength. What is really sought is

some abnormal condition that leads to a de- crease of strength, but so far this phenomenon

has not been identified.

396
Id 120 GRIT ROUGHNESS (AA) 40-*5 p in Jfl 600 GRIT ROUGHNESS lAA) 22-26 (J in.

Figure 2. Profile traces of alumina surfaces ground

with different grit sizes of diamond.

6. Conclusions

From the work described, the following conclusions can be drawn. Grinding of ceramics

is a necessary operation because the ceramics

manufacturing process is presently unable to

produce an as-fired shape having a closely specified geometry. When reasonable care is exercised in grinding, the strength of ceramics is

not markedly affected by changes in grinding

variables over a wide range; however, excessive infeed (deep cut) is detrimental to

strength. Surface features that we presently

can observe and monitor are not a very sure

indicator of strength.

The authors wish to acknowledge the support

of the Naval Air Systems Command who sponsored this work under Contracts NOw-66-

0383-d and NOOO19-67-C-0494.

7. References

[1] Sedlacek, R., and Halden, F. A., Method of tensile

testing of brittle materials, Rev. Sci. Instr. 33,

298-300 March (1962).

[2] Sedlacek, R., and Farley, E. P., Processing of

Ceramics—Surface Finishing Studies, Final

Technical Report under Contract No. NOw

66-0383-d, Stanford Research Institute, May

(1967).

[3] Sedlacek, R., Processing of Ceramic—Surface

Finishing Studies, Final Technical Report under

Contract No. NOOO19-67-C-0494, Stanford Research Institute, November (1968).

[4] Gielisse, P. J., et al.. Ceramic Finishing with

Diamond, Part I: The Workpiece, Ceramic In- dustry 80 February (1968), Part II: The Abrasive and the Bond System, ibid., March (1968),

Part III : External Influences and Cost Analysis,

ibid., April (1968).

Discussion

Gielisse: Mr. Sedlacek, when you compared

strengths of aluminas ground with the manmade 100 grit resinoid wheel versus 320 grit

natural diamond metal bond wheel, did you intend to contrast bond type or did you intend to

contrast the manufactured versus the natural?

There are two possibilities there.

Sedlacek: Well, I really did it simply out of

curiosity. At that moment I had no reason to

anticipate any difference.

Unidentified: I wonder if some of those

spikes as you called them could be due to strain.

My second question is did you or did you not

stress relieve after grinding?

397
Sedlacek: I think that the spikes are due to

pullouts rather than to anything else. We did

not do any stress relieving after grinding.

Lange: In your scanning electron microscopy

work, did you look for and see any surface

cracks? I think you would need higher magnification.

Sedlacek: We have used higher magnifications and we do find cracks. We are always

trying to look for some pertinent feature which

we could correlate with the tendency to break,

but so far it has eluded us. Daniels: Was there any attempt to try to

correlate strength with either the frequency or magnitude of those pullouts?

Sedlacek: No, we have not done that.

398
Edge Effect on the Modulus of Rupture of Ceramic Substrates

Wing C. Lo

Bell Telephone Laboratories, Inc.

555 Union Boulevard

Allentown, Pennsylvania 18103

Because of its simplicity, the modulus of rupture (or flexural strength) test is often used to evaluate the mechanical strength of ceramics. For large flat substrates,

the natural tendency is to test specimens cut from the substrate.

Depending on the method of cutting and the resulting edge condition of the

specimen, the flexural strength can be significantly diff'erent. This is illustrated by

3-point loading data on COi-laser-scribed and diamond-sawed specimens. The modulus

of rupture of a laser-cut sample is about 70 percent of that of a diamond-sawed sample.

Key words: Alumina substrate; edge damage; flexural strength.

To evaluate the mechanical strength of ce- ramic substrate materials, there is a general

tendency to use the bend test (3-point or 4-

point loading) on cut specimens. Although there

is awareness of the effect of the bearing surfaces, loading rate, the span length, size and

surface condition of the test specimen ^ on the

modulus of rupture calculated from bend test

results, that of the edge condition of the specimen is often neglected. Since the damage to

the surface at the cutting edge produced by

different cutting methods can vary with different materials, the resulting modulus of rupture

does not truly represent that property of the

material tested. Of course, the test would be

adequate as a use test for evaluating or monitoring the quality of piece parts from different sources if the method of cutting and preparing

the test specimens is the same as those used for

the piece parts.

This note presents some experimental data

to show that the edge effect on the modulus of

rupture can be highly significant.

Figure 1 shows the modulus of rupture of

two samples of Alsimag 772 - substrates taken

from the same lot. One sample is cut by laser

scribing and the other by diamond sawing. The

laser scribing method uses a pulsed beam from

a CO2 laser to effect a series of hairline cracks

on the substrate which may then be snapped

apart by hand. The CO. laser used is model

40A made by the Coherent Radiation Laboratories with an output wavelength of 10.6/i.m.

The pulse rate is —- 333 pulses/s and the feeding rate is 150 in/min (64 mm/s). The diamond sawing is done with a 0.017 in (0.43 mm)

thick 3 in (7.6 mm) diameter metal bonded

diamond wheel (Norton's designation D220-

NIOOM). It has a diamond imbedded section of

0.125 in (3.2 mm). The wheel rotates at 2,500

RPM. The feeding rate is approximately 1.3

> See for example ASTM F-103.

^ Alsimag 772 is a 99.5 percent alumina body made by the American Lava Corporation, Chattanooga, Tenn., U.S.A.

500

400

200

80

70

.9- 40

30

20

LEGEND: ® SAW-CUT SAMPLE a LASER-CUT SAMPLE

NOTES:

1 - SAMPLE SIZE = 8 2- 3-POINT LOADING FIXTURE USED ; SPAN = 0.25 INCH (6.4 mm)

3 - LOADING RATE = 0.05 INCH /MIN.

(0.021 mm/S) .

J I I I I I L

10 20 30 40 50 60 70 PERCENTAGES

80 90

Figure 1. Distribution of flexural strength of laser-cut

and saw-cut ceramics (Alsimag 772 substrates).

in/min (0.55 mm/s). The size of the test specimen is 0.350 in x 1.300 in (8.89 mm x 33.0mm

with a thickness of approximately 0.025 in

(0.63mm). In testing, the scribed side of the

substrate is placed in tension.

The bend test was performed with a three

point loading fixture using a 0.25 in span and

a loading rate of 0.05 in/min. Eight specimens

from each sample were tested.

The lower value of the laser-scribed sample

(68% of that of the diamond-sawed sample) is

attributed to the damage on the surface at the

cut edge. This is made visible with the aid of a

fluorescent penetrant (see fig. 2.) The damaged

edge of the laser-cut ceramics adsorbs the

penetrant and appears bright when viewed un- der ultraviolet light.

Examination under the microscope using the

replica technique (fig. 3) shows quite clearly

399
CERAMIC EDGES

LASER-CUT

CERAMIC

SAW -CUT

CERAMIC

Figure 2. Zyglo test on laser-cut and saw-cut ceramic

substrates.

the presence of cracks in the amorphous or

recrystallized material in the region where the

laser strikes.

Similar results of edge effect on modulus of

rupture were also reported by P. T. Morzenti.^

Laser-scribed samples of Alsimag 772 and ADS

995 * gave modulus of rupture values about 72

percent and 73 percent respectively of those

cut with a 15/im diamond wheel. However, with

the ERC-105 substrate ^ which has a much

finer grain structure than the other two ce- ramics, the laser-scribed sample did not show

any degradation in strength. This may be ex- plained by the greater energy required to propagate cracks in a fine grain material which has

more grain boundaries than in a coarser grain

material (see figure 4),

These results are not to be construed to the

use of laser-cutting as poor practice but are merely pointing out that when comparing modulus of rupture of ceramics one should be aware

of the edge condition of the test specimens.

Thanks are due to Messrs. J. E. Clark, W. B.

Grupen, F. L. Rowland and P. I. Slick for their

comments and to Mr. P. T. Morzenti for the

use of his data.

' Private communication.

* ADS995 is a 99.5 percent alumina body made by Coors Porcelain Company, Golden Colorado, U.S.A.

^ ERC-105 is a 99.5 percent alumina body developed by Western

Electric Company Engineering Research Center, Princeton, Newr

Jersey, U.S.A.

" The average grain size of the ERC-105 substrates is in the order of 1/im, while that of the other two samples is several microns.

-EDGE

-CERAMIC

LASER-CUT

NOTE

EDGE

CERAMIC

SAW-CUT

THE DARK LINES PERPENDICULAR TO

THE EDGE ARE ARTIFACTS FROM THE

REPLICA.

MAGNIFICATION :

lOO/i

I 1

Figure 3. Top view of edges of laser-cut and saw-cut

ceramic substrates by transcopy replica technique.

CRACK

a = ALSIMAG 772 SUBSTRATE

b = ERC SUBSTRATE

Figure 4. Cracks generated by laser beam on 99.5

percent alumina substrates.

Discussion

Unidentified: The cracks in some of the grain boundary cracks. Is that true?

laser-scribed substrates looked very much like Lo: Yes.

400

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).
VI. PANEL DISCUSSION

J. B. Wachtman, Jr., Moderator

F. Ernsberger, P. J. Gielisse, A. H. Heuer, W. H. Rhodes,

and R. W. Rice, Panel Members

J. B. Wachtman, Jr., National Bureau of

Standards, Gaithersburg, Md.

Panel Moderator:

Our intention here is to put the work presented at this conference into perspective. We

have had a series of excellent invited and contributed papers addressed to various specific

aspects of the sequence of machining, finishing,

characterization, and physical properties. We

now wish to look over this sequence from the

point of view of several factors which are important throughout the sequence and which,

accordingly, provide a basis for a unified point

of view. Such a unified point of view depends

upon an understanding of the processes occur- ring at the surface during machining and the

consequences for characterization and for control of properties. Many of us are interested in

physical properties and the factors controlling

them, but we would like to focus the present

discussion as much as possible on the processes

occurring at the tool-workpiece interface, their

interaction with the machining system, and

their influence on the finished piece. Each of

our panel members has chosen one or more

themes around which to make his remarks.

Each panel member's presentation will be followed by a short discussion period.

It is appropriate to begin with the question

of the mechanism of machining. No one is

better qualified to summarize present views

and suggest areas needing further elucidation

than Professor Peter Gielisse of Rhode Island

University who has already presented a paper

on his single point machining research and who

needs no further introduction.

P. J. Gielisse, University of Rhode Island,

Kingston, R. I.

Panel Member:

I would like to discuss the knowledge we have

on the mechanisms of ceramic finishing processes. It is of course understood that except

for speciality applications, most of this falls

in the area of the mechanical methods of stock

removal or ceramic finishing. In reading over

the various papers and from what I have

heard during this conference I cannot get away

from one particular thought. There remain,

unfortunately, a large number of misconceptions, some of which, and only some of which,

have possibly been done away with as a result

of the actual presentations. We find, for instance, that reference is still made to the celebrated fact that grinding of ceramics is the

same as the grinding of metals and that nothing new needs to be learned. A statement like

that at a conference of this type certainly is

extremely surprising and particularly surprising in the light of many of the presentations

that were made today. A second claim which

could be made is that all one apparently needs

to do to evaluate grinding is to generate some

sort of a grinding ratio or similar parameter

and that physical characteristics and concepts

can be replaced by an apparently more convenient shorthand notation, as seen on many of the

slides on the screen here, which evaluates matters in terms of such (brand name) parameters

as "XDY-MB3." The strength of materials is

in effect claimed to be better expressed as a

function of extrinsic parameters such as grit

size rather than in terms of what is germane

and intrinsic to the process and to the effective

strength of materials. These are just a few

observations and they are given merely by way

of example here. What I am trying to indicate

is what I believe to be the general futility in

our efforts to understand the mechanical re- moval mechanism and the many observed ex- perimental phenomena, in terms of what is

only a part of the total systems analysis,

namely the (important) economic aspects. It

is, however, particularly noteworthy that even

the economic analysis cannot really be fruitfully carried out unless we have an understanding of the full technical aspects. It is therefore

important, I feel, that we continue to develop

the science and technology of the grinding

mechanism so as to generate and develop generally applicable ceramic finishing data. This

information will provide a further understanding of the nature of the influence of the finishing cycle on the finished product, will guarantee

that we have reproducibility in the finished

ceramic part with regard to both the internal as well as external characteristics, and will

generate new methods and approaches in ce- ramic finishing. I think it is not very well understood that the way we grind, the way we re- move stock today, need not necessarily be the

way we remove stock in the future. It is only

when we try to understand the fundamentals

involved in the materials removal process that

401
we will come closer to understanding the possibilities for improvement.

Attempts to further understand these matters should really be concentrated on essentially

four areas. First of all, the recognition and

study of the fundamental differences between

the three basic methods of stock removal:

Bonded Abrasive Machining (BAM), Contained Abrasive Machining (CAM), and Free

Abrasive Machining (FAM). Secondly, the nature of the interaction between the tool and the

workpiece from both a mechanical as well as a

materials point of view. It is very important

that we do not merely look at this from a mechanistic point of view or merely from a materials point of view. These two must come together. Thirdly, the function of each parameter in the tool as well as the workpiece

which contributes positively or negatively to

the stock removal process. Finally, the external

or environmental effects which influence the

system's performance advantageously or disadvantageously as part of the total system, not

merely as an observation of the phenomenon.

In other words, when we observe the phenomenon of the influence of the environment, it

is not sufficient merely to make that particular

observation, but it certainly ought to be tied

to the parameters of the actual finishing system as a whole.

A. H. Heuer, Case Western Reserve University,

Cleveland, Ohio:

It struck me while Bert Westwood was giving his talk yesterday that it would be important to find a "cutting fluid" for alumina.

Probably more alumina is machined by diamond abrasives than any other single ceramic

material and most of this machining is done

wet. If a medium were found in which the

cutting efficiency could be increased by two to

twenty-fold, as Bert showed for some of his

extreme examples, it should have considerable

economic advantage.

K. M. Smith, Applied Physics Laboratory,

Johns Hopkins University, Scaggsville, Md.:

My experience is that high spots on a grinding wheel are particularly damaging when

grinding is parallel to the face of the specimen

as shown in Figure 1. Specimens ground in this

way tend to crack upon refiring. High spots

apparently do less damage when the specimen

itself is also rotated during grinding as shown

in Figure 2. Commenting upon your statement concerning

alumina, I found that morphaline is one of the

safe coolant agents that I could add to water.

Because it is an ammonia compound, it is very

easily volatilized and it does not add any impurities like sodium or potassium that some

coolants would do. So I use morphaline in mistic

Figure 1. Grinding parallel to specimen face.

Figure 2. Grinding with specimen rotating as well as grinding wheel.

mist at about 125 lbs pressure. I must agree

with your first statement. Dr. Gielisse, I don't

think we can compare grinding metals and

grinding ceramics. I think they are two separate materials. They must be handled differently.

P. J. Gielisse, University of Rhode Island,

Kingston, R. I.

What you are stating is exactly the point I was trying to make. Some people want to take

our knowledge of metal grinding (even there

we know very little) and transfer it to the

ceramic area. Now the only similarity I see is

that in both we commonly use, in the majority

of round wheel, and the label of the

402
manufacturers of the machines might or might

not be identical. For the rest, there is absolutely no correlation.

H. A. Perry, Naval Ordnance Laboratory, Sil- ver Spring, Md.:

My horizons have widened at this meeting.

I've been in the habit of thinking of grinding

in terms of mechanical and chemical practices

and also, possibly the actual phenomena used

to explain the Rebinder effect. It comes to

mind, however, that also in the environment

are the occasional electric properties of the

material which are involved in fracture and

similar effects. Perhaps even optical and magnetic or electromagnetic effects are important.

It seems to me that the picture is far more

complicated in my mind than when I came.

P. J. GiELiSSE, University of Rhode Island,

Kingston, R. I.: I think this is quite true. I believe that we

made an observation on triboluminescence and,

again in keeping with the idea of doing the

fundamental studies, we have done studies in

gases; we alluded to that in our paper. We

very definitely see optical effects that are either

the result of gaseous discharges or the result

of triboluminescence. This could make for the

so-called sparks in ceramic grinding. I believe

that from all evidence we have right now, although it is still to an extent meager and

possibly qualitative, it is certainly indicated

that the so-called temperatures of grinding are

considerably lower than most of us have felt

them to be.

A. R. C. Westwood, R.I.A.S., Martin Marietta

Corporation, Baltimore, Md.:

Responding to one of Dr, Gielisse's points, I would like to propose one way in which an en- vironment can influence the ease with which

cracks are initiated in semibrittle solids. Once

initiated, I would expect that the cracks would

propagate and interact to form the chips which

constitute the means of material removal. Consider the stress-strain curves shown in figure 3.

Studies of the room temperature flow and fracture behavior on semibrittle solids, such as MgO monocrystals containing many surface

sources, indicate that only limited work hardening occurs following yielding. Thus, the fracture stress is essentially determined by the

stress at which dislocations first move and multiply, that is, the flow stress. In other words

Fi —' fi. If such a solid is placed in contact

with an environment which facilitates dislocation mobility in the near surface regions,

then the flow stress of the near surface region

will be decreased, e.g., to fo. It follows, of

course, that the fracture stress of this region

Rebinder Effect (Af

)

STRAIN

Figure 3. Schematic illustration of manner in which

the Rebinder effect (the adsorption-induced reduction

in hardness) can facilitate fracture, and hence machinahility, of a semibrittle solid such as MgO,

a

Figure 4. Indenter tracks in freshly cleaved MgO sur- face produced (a) in air—a relatively inactive en- vironment, and (b) surface-active n-hexadecane. The

load on the indenter was 20 g-., and it was moved

from left to right. Note prolific cracking at tracks

produced under the active environment. (After Westwood and Goldheim, J. Am. Ceram. Soc. 53, 142-7

(1970)).

will also be lowered, e.g., to F2. Thus, by reducing the flow stress, the fracture stress is also

reduced, and material removal facilitated.

An example of environment-facilitated

cracking in MgO is shown in figure 4. In this

case, comparative scratch tests were performed

on freshly cleaved {100} MgO surfaces. A 20 g

load was applied to the diamond indenter,

which was traversed at about 2 cm per min.

The environments compared were air (that is,

adsorbed water) and n- hexadecane. Under

403
these conditions, cracks were not formed in

surfaces exposed to air (fig. 2a). Under hexadecane, however, cracking was prolific (fig.

2b), fracture occurring on both {110} planes

intersecting the surface and {100} planes below the surface.

P. J. GlELiSSE, University of Rhode Island,

Kingston, R. I.: Thank you Dr. Westwood. Just one observation. Might I put in a plea that most of our observations in the future be conducted under

conditions that are typical of the grinding

mode. I feel, namely, that in addition to the

phenomena you describe there are unfortunately many processes that simply go unnoticed,

that are the result of high impact-high velocity

type of material deformation processes.

J. A. Mueller, The Carborundum Company,

Niagara Falls, N. Y.:

We've seen photomicrographs of a lot of

samples which appeared to me to be badly

damaged. Some of the parameters of grinding

that were used to grind these samples would

be considered to those skilled in the art as

unsatisfactory. One example was the use of

very heavy in-feeds with slow table speeds

using an aluminum oxide vitrified wheel for

grinding high purity, dense nonmetallics. This

method could do nothing but produce a badly

damaged surface. But my purpose in commenting now isn't to talk about that. My purpose is

simply to point out that there is a wide lack

of communication between those skilled in the

art and the people who are doing this kind of

work. Another thing, I think, that is quite well

taken, concerns the debate about whether or not you can transfer knowledge of metallic

grinding into the field of nonmetallic grinding.

Now I don't know whether you can or cannot,

but I know this—that there are many interesting phenomena that occur in metal grinding

that also occur in ceramic grinding. For ex- ample, I think we saw one just a moment ago

when the type II wheel was compared with

the type I wheel. The gentleman indicated that

the type I wheel grinds on one point. No question about it. We've got documentation in

metal grinding to show that a wheel will grind

on one point and that this point travels around

the wheel face 360 degrees. There is documentation available and there are movies available

that illustrate this phenomena.

Another example stated was grinding a large

ceramic cylinder with a 3 in diameter grinding

wheel running at a slow surface speed.

What happened was no doubt, that the use

of a small wheel, grinding a large part, produced a badly damaged surface. The length of

contact was too great for the wheel to operate

efficiently. Again we have illustrated that, what

happens in metal grinding, happens in nonmetallic grinding.

In metal grinding, when we grind hard

brittle materials we use fine grit wheels. Some

general rules are: hard materials—fine grit

wheels, softer materials—coarse grit wheels,

fine finish—fine grit wheels, coarse finish

coarse grit wheels.

In nonmetallic grinding we have found to

date, that the same rules of grinding apply.

There is a whole area in metallic grinding that

compares to nonmetallic grinding. We are not

certain that we can transfer it bodily. However, I feel that people interested in ceramic

grinding should become acquainted with what

knowledge is available in the metallic field and

then attempt to transfer whatever is applicable. What I am saying is, I think there is a tremendous lack of communication between

the gentlemen here and the gentlemen in the

grinding operations and I think that they

should get together. As a matter of fact, I would like to suggest that possibly when we

grind samples there should be a standard

method for grinding samples. Find the best

way to grind a sample and then use that as a

standard. In this way, better results will be

produced and faster progress will be made.

P. J. GiELiSSE, University of Rhode Island,

Kingston, R. I.: I'd like to make a comment on the first point.

Some of the conditions used in some of the

papers here were not necessarily representative

of those in practice, but this is often a very important and necessary step in broadening our

understanding. I think you often have to look

at a much broader range of parameters than

maybe the ideal ones for industrial practice in

order to come to an understanding of the true

mechanisms. If you can get a better picture of

what goes on by using parameters which really

aren't practical from the industrial standpoint,

you may be able to transfer this knowledge to a better understanding of what will be happening under the conditions that are of practical

interest. So I think there is some communication and appreciation of the point you make.

Many of us may not know as much as we

should about the practice and technology but I think you should recognize also that it is very

often important to use "impractical" conditions in order to generate "practical" knowledge.

A. U. Daniels, Metcut Research, Cincinnati,

Ohio:

Whether or not there is a standard grinding

test, I think it would be extremely appropriate

to work out a standard method for presenting

404
the variables related to grinding or other cutting operations ; that is, a systematic format

for giving feeds, speeds, depths of cut, wheel

types, coolants and so forth. If a standard format for machining data were used in papers

and presentations, one could tell at a glance

what the authors had done in general instead

of having to ask about miscellaneous machining parameters afterward. I think that standardizing data presentations would improve

communications whether or not standard tests are being used.

With regard to transfer of information from

work done on machinability of metals—if I had a lot of wood to cut up and had my choice

of hiring either someone who had never cut

anything up or someone who had spent all of

his life cutting up metal bars, I'd take the man

with at least some cutting experience. In other

words there is much useful machining theory

and experience gained with metals which is

transferable to ceramics. For one thing, metals

aren't all like aluminum or soft steel. Metals

run the full gamut of ductile and brittle properties. In fact, when aerospace alloys, hardened

tool steels, and metal bonded tungsten carbides are considered, there is almost a continuous

spectrum of materials from metallic lead to

aluminum oxide ceramic which must be machined. Therefore, I think ceramic machining

is only one aspect of a larger problem, and I have to ask why Dr. Gielisse would categorically state that everyone should ignore all of

the metals machining background that has

been developed.

P. J. Gielisse, University of Rhode Island,

Kingston, R. I.: First of all, your statement which started

this debate is only partially correct. If the picture really did lie the way you painted it then

I too would say yes. If our objective were

merely to cut ceramics or remove stock, fine, I would also hire the man who has this type of

experience, but that is exactly the point; we

want to move away from this, we want to

understand what happens during this particular kind of cutting. The objective is not

merely to cut, I mean not really with only that

objective in mind. In that case, we could

merely continue turning out grinding ratios

and we would be able to make some sort of

economic evaluation, but it would not give us

the sort of data with which we could accurately go forward.

F. F. Lange, Westinghouse Research and Development, Pittsburgh, Pa.:

Has anyone tried to cut ceramics at higher

temperatures ?

H. C. Miller, Super-Cut, Inc., Chicago, 111.: We have done this with glass with which our

company has had a lot of experience. We frequently get a glass combination that will crack

when you drill it and there just seems to be

nothing you can do. Yet if you will run it in

water as hot as a man can possibly stand, you

can drill beautiful holes without thermally

shocking the glass. There are other ways to go

to higher temperatures for grinding which

have a first order effect.

S. D. Brown, University of Illinois, Urbana,

111.:My son is an ichthyologist of sorts. He likes

tropical fishes and every once in a while he

will come home with a big bottle of some kind

to cut. One time he came home with a carboy.

The best way I found to cut it was to put a hot

wire on it and it just pops right on through.

Now I wonder if you couldn't do some cutting or machining by controlled heat input in low

cross sections.

D. J. Donovan, Data Magnetics, Torrance,

Calif.:

There are some possible industrial applications: changes in the magnetic properties of

ferrites are induced by a hot wire at a given

point. This may have application by locating

the gap in a magnetic head core. Note added in proof:

The normal way to create a gap in a head

core is to make the core in two pieces and then

bond them together. This hot wire method

would not necessarily be a physical gap, but a disturbance in the magnetic path.

J. B. Wachtman, Jr., National Bureau of

Standards, Gaithersburg, Md. Panel Moderator:

We'll take our next two panel members in

sequence and defer discussion until after the

second presentation. We've had many refer- ences to plastic deformation. I believe both of

the next two speakers intend to touch on that

as one aspect of their talk. Our next panel

member is Dr. Fred Ernsberger of PPG Industries. Dr. Ernsberger is well known for his

own work on hardness and other properties of

glass and he has general extensive knowledge

of glass science and technology through the interests of his firm. The presence of a glassy

phase in many, perhaps the majority of technical ceramics, makes it very appropriate to

call on Dr. Ernsberger for his comments on

this aspect of machining.

F. Ernsberger, PPG Industries, Pittsburgh,

Pa.

Panel Member:

Thank you very much Jack. I am glad you

made those introductory remarks because I

405
wouldn't want anyone to think that I am a

specialist in machining of ceramics. My own

area of expertise in glass is more in the area

of strength ; but that does give me some area

of overlap with the material of this conference.

I suppose that's why I was called to come and

be an observer and to give my reactions.

As Jack has hinted, my principal reaction

was to the number of times I have heard the

concept of plastic flow in glasses and ceramics

used very casually. I have some objections to

that which I would like to raise. However,

there are a few other things that I would like

to mention before I get into that.

I feel very comfortable when I hear people

talk about stress-corrosion and delayedfracture effects and so on, because we see those,

of course, in our work with glasses as well.

We use about the same thinking in rationalizing them ; the presence of water, stress corrosion, the Charles-Hillig theory, and so on. I gather that you are not entirely comfortable

with this approach and this rationalization and

I guess we're not either; though perhaps the

chemistry is a little clearer and makes a little

more sense in our case than in yours, so maybe

we're a little more comfortable with it. I'm also comfortable with the concept of

surface flaws and their effect on properties, including strength. Naturally, volume flaws are

also a matter of concern in your polycrystalline

materials. What surprised me a little was Pollock's statement that with his thermally polished filaments of single crystal sapphire he

was able to find internal origins for, I believe

he said, every one of the cases of tensile fracture. He did this by potting is resin or something of the sort. This is something that no

glass person has claimed ever to have done;

either to find internal origins or, indeed, to

find any origin in the fracture of a glass fiber or glass rod specimen under pristine glass conditions. There is too much fracturing and

breakup. Possibly it is because of the higher

modulus of alumina, so that alumina gives a higher stress without so much stored energy.

At any rate, I wish he would come to a glass

conference sometime and tell us how he did

that.

The etching technique to reveal surface flaws

is used, and a particularly pretty example was

the paper we just heard a few minutes ago by

Koepke on magnesium oxide. I think perhaps

some of you are not using this enough. It is

really a very powerful technique. I gather that

it isn't quite as simple or as convenient to

apply in the case of ceramics, particularly

TDolycrystals, as it is in the case of glass. Dilute

hydrofluoric acid is our all-purpose etching

reagent, and it is easy to use at room temperature. It does give in many cases beautiful definitive results which you can hardly get any

other way. You should use it more often, even

if it does mean high temperatures and corrosive gases and things of that kind. I was sorry

to see that in a case where etching could have

been used readily to advantage (I'm speaking

of Dr. Busch's movie), there was no sign that

it had been considered. You remember the first

sequence in that movie where he was dragging

glass over a sphere; behind the point of contact there was what he described as stress

waves or something. Any glass man would immediately guess that those were really cres- cent-shaped cracks which temporarily were

held open by the stress behind the moving ball

and then closed up and became invisible. One

could easily check this hypothesis by simply

etching the sample after the experiment was

done. I hope this will be done sometime with

that sort of experiment. I rather confidently

predict that crescent cracking is what that

effect will turn out to be.

The Rebinder effect was very prominently

mentioned in Westwood's presentation and I have to report that among glass people the Rebinder effect has hardly made any waves at all. Perhaps we are missing something very important; and I, for one, am going to go back

and study this paper and the data very carefully because I can see that at least one man is

convinced that this is a very important effect.

But among glass people, I suspect they mostly

think as I do that the Rebinder effect is just

one of those odd things that the Russians have

come up with, like "polywater." These things are hard to disprove, but in the end will probably turn out to be nothing at all. Now, if I have a little time left, I want to get

to this question of plasticity and plastic flow. I consider it to be a hard thing to defend in most

crystalline oxides and in all silicate glasses. I make one prominent exception: magnesium

oxide. Magnesium oxide obviously shows very

readily definable plastic effects at room temperature. One reason for this is the fact that

it is a highly ionic material like sodium chloride, where you have similar effects ; and another reason is that it has plenty of slip systems, inasmuch as it is face-centered cubic.

Among the more covalent materials (of course,

there is a spectrum with sodium chloride on

one end and diamond on the other, and in this

spectrum I certainly would place alumina

among the more covalent materials) there is

less justification for assuming that slip can

happen at ordinary temperatures. The activation energy for movement of the existing dislocations is simply too great to be theoretically

accountable at room temperature. So when you

postulate this sort of thing, you have to be able

either to explain theoretically how this can be, or else you have to point to a reasonable as- sumption that the temperature at least locally

406
and for a certain period of time was rather

high. That is what I have to assume when I hear discussions of the burnishing of alumina,

of dislocation concentrations near the surface,

and things of that kind.

In glasses, you have what are called "microplastic effects." Until recently, these have been

rather difficult to explain. Under the Vickers

indenter you really can get a permanent identation which is crack free. This bothered me

for a long time, but I published a paper ^ re- cently which rationalizes this, not in terms of

plastic flow, but in terms of densification.

A few of you may find densification a novel

concept so I'll just briefly explain what it is. Glasses having low coordination numbers for

silicon (and for oxygen for that matter) have

a lot of empty space in them. When compressed

hydrostatically, or with various complex

stresses that involve a substantial hydrostatic

stress, they will condense in volume irreversibly. That is, when the stress is removed, you

find that the specimen has actually diminished

in volume. This is not plastic flow in the proper

sense of the word, because plastic flow occurs by shear, it involves changes in atomic partners, and volume is conserved. Densification is

clearly not that kind of a process. It is more of

a displacive than a reconstructive process. I think Stokes put the thing rather clearly this

morning when he talked about the "semibrittle" materials as opposed to the "absolutely

brittle" materials. I would nominate alumina

and our silicate glasses as being among the

"absolutely brittle" materials at ordinary temperatures.

Thank you.

J. B. Wachtman, Jr., National Bureau of

Standards, Gaithersburg, Md. Panel Moderator:

Please hold all the comments that you have

in mind as a result of Dr. Ernsberger's discussion until we have heard from Roy Rice of the

Naval Research Laboratory. Mr. Rice has done

extensive work on ceramic processing including

high temperature pressure forming involving

plastic deformation; he will also discuss the

role of plastic deformation in machining.

R. W. Rice, Naval Research Laboratory, Washington, D. C.

Panel Member:

I would first like to briefly mention what are some of the factors in the occurrence of microplasticity in machining. Slip can occur in some

softer materials (e.g., sodium chloride and

magnesium oxide), and twinning may occur in

some harder materials such as alumina at moderate stresses at room temperature. However,

this does not explain the scope of materials in

' "Role of Densification in Deformation of Glasses Under Point

Loading," J. Am. Ceramic Soc. 51, 545-547 (1968).

which such surface deformation occurs during

conventional machining, nor the extent of such

deformation as reported by many at this conference. The most important factor is probably the high compressive stresses that result

from machining. These stresses will, I think,

considerably enhance the opportunity for slip to occur first by inhibition of cracks due to their

compressive nature, and second, by the high

level of the stresses themselves. The rise in

local temperatures in many machining operations may also measurably enhance the occur- rence of microplasticity.

Next, I would like to look briefly at some of

the possible roles of microplasticity in a typical

machining operation. In a rough machining

operation such as sawing, I would expect that

fracture is most likely the predominant process. Here the role of microplasticity is probably one of aiding the amount, but limiting the

depth of fracture, for example, by multiple

crack nucleation. blunting cracks, and general

energy absorption. In finer machining operations (e.g., most grinding) the propensity for

cracking should diminish (in fact, I have observed this in MgO) and the extent (i.e., depth)

of microplasticity to at least remain constant, if

not increase. In polishing, microplasticity is

probably the main mechanism of developing the

fine surface topography and does not occur as deep as in grinding, as observed by Hockey in

ALO.i and myself in MgO. I think we should

also note that, as has been pointed out previously, too much plasticity under certain machining conditions can be bad, e.g., certain machining conditions applied to soft materials,

such as sodium chloride, would probably be

rather difficult.

I would like to briefly mention, without infringing too much on Bill Rhodes' topic, some

of the material parameters that may affect

microplasticity during machining. First, is

hardness. As I presented in one of my talks,

we generally will expect the depth of microplasticity to decrease with increasing hardness.

Second, is grain size. I have observed the

depths of microplasticity to decrease with decreasing grain size in magnesium oxide, which

is a rational trend. Third, is the effect of porosity. I think that one of the important things

we have to recognize is that the location of

pores (i.e., at grain boundaries, or within

grains) is an important factor. This may be a

key factor in the type and size of chips formed,

and hence, important in both the machining

operation and the topography of the resultant

surface. There is also evidence that pores may

be sources of dislocations; if so, this will also

make the location of pores important (e.g.,

pores within grains should then enhance sur- face work hardening). Impurities are, of

407
course, another important microstructural factor affecting microplasticity.

I would now like to discuss mechanical properties which have clearly been of considerable

interest in the conference. Traditionally, the

view would have been that the effect of machining was to introduce flaws that caused failure,

thus determining strength. Both the growing

evidence for microplastic processes associated

with fracture and the evidence of plastic deformation during machining shown at this

conference indicates that this is a rather shortsighted view. I find mechanical behavior much

more understandable if I consider both microplasticity and flaws where flaws may be much

more prevalent in certain materials, particularly the hard ones ; while microplasticity would

generally be more prevalent in softer materials.

Twinning may also be important, possibly more

at intermediate hardness, but it may not cor- relate as well with hardness as slip. I can see many situations in which microplasticity and

Griffith flaws are competitive as to which may

be the main cause of failure. I can also visualize situations in which these two mechanisms

cooperate to lead to failure, as Bert Westwood

also suggested (e.g., subcritical crack growth

by microplasticity.

)

I think many people tend to take a very skeptical attitude toward microplasticity since the

Griffith concept has been around for a good deal

of time and had a good deal of success in ex- plaining glass behavior. Further, many people

seem to seriously neglect the known or possible

differences between microplasticity in ceramics

and ductility in metals, and hence confuse

them. No one duobts the importance that flaws

can have on mechanical properties. However,

if you look carefully at the evidence for mechanical failure in crystalline materials, I think you will find that the Griffith concept is

no more firmly established than is the microplastic concept. For example, as Bob Stokes

pointed out, there actually has been relatively

little demonstration of Griffith flaws as a re- sult of various machining or other operations

in crystalline materials. In particular, there

has been, to my knowledge, no specific fracture

origins in single crystals demonstrated to be

due to particular (i.e., known preexisting)

flaws which clearly fit the Griffith criteria.

There has been very limited observation of

flaws in polycrystalline materials. One of the

few cases that I can think of is rutile, and

here, flaws were observed only iniarge grain

bodies, did not go clear across the grain, and

the behavior of these bodies was considerably

different than the finer grain materials in

which flaws were not observed.

A basic problem in applying the Griffith

theory where it appears applicable is to explain

the generally pronounced effect of grain size

on strength of materials. It is commonly as- sumed that flaw size is related to grain size,

which is certainly an appealing intuitive

thought, but there is really no firm foundation

that I know of to establish this. On the other

hand, we have seen very good evidence at this

conference that microplasticity does occur during machining of materials such as alumina

and magnesia, and there is evidence in the lit- erature for materials such as silica and tungsten carbide as well. Thus, since nucleation of

cracks by microplasticity is related to grain

size, microplasticity during machining may

offer an explanation for the apparent relation

of the size of at least some flaws to grain size.

There also appears to be an important interaction between the microplastic deformation

from machining and flaws. Hockey has observed that the deformed layer results in sur- face compressive stress. This is quite reasonable, and may be quite important.

Note added in proof:

While it is probably obvious, it is worth emphasizing two important effects of machining

on microplastic mechanism of failure. First,

while microplasticity during machining may

nucleate cracks, it may in other cases inhibit or prevent crack nucleation as noted earlier. Thus,

this can be an important factor in allowing

microplastic mechanisms of failure to operate.

Compressive surface stresses may also aid this.

Second, fresh sources of dislocations (and

probably twins) are often important to microplastic mechanisms of failure in ceramics. This

conference clearly shows machining is a

source of these.

Thus, in summary, while both machining and

material parameters affect microplasticity during machining, it probably always occurs to

some degree. This is important to both the

machining process and the surface character

of the body and thus resultant mechanical

properties. Both flaw and microplastic mechanisms of failure are affected by the microplasticity in machining and each must be evaluated.

S. D. Brown, University of Illinois, Urbana,

111.:While it is altogether possible that very high

temperatures are present locally under the tips

of diamond indenters during indentation or scratching in which permanent deformation of

glass is produced, it is not necessary that this

be the case. Calculations show that even in the

adiabatic approximation, the temperature rise

for 15 percent densification of a typical soda

lime-silica glass need not be more than ca. 10 °C.

408
I believe that the question of whether plastic

flow is involved in the indentation of glass depends somewhat on how one defines plastic

flow. If one is talking about so-called ideal

plastic flow in which there is no volume change,

then of course densification is involved in the

case of glass, the volume changes, and the

plastic flow (if that's what it is) is not ideal.

In the case of glass indentation, the flow occurs under the influences of both hydrostatic pres- sure and shear. I would guess that free volume

is removed from the glass by such a process,

thus leading to densification. The activation

energy of such a process would be expected to

be small on account of the relatively large shear

stresses present during the indentation. These

large shear stresses are absent when the indenter is removed, so the free volume cannot

return at any sensible rate; hence, the densification is permanent until the glass is heated to

a temperature where thermal energy promotes

the annealing reaction. The flow of the glass

under pressure might be termed non-isometric

plastic flow. The reason one does not see similar phenomena in crystals (e.g., ALO., is, I suppose, due to the fact that the crystals do not

contain as much free volume as glasses do.

F. Ernsberger, PPG Industries, Pittsburgh,

Pa.:

It is evident that Sherman Brown has

thought very hard about this from several

angles and I encourage you all to do the same

thing. In fact, that is all I want to do today;

I don't want to try to defend any particular

concept that I have advanced. I just think it is

something that you should think about, particularly before you write for publication. I'm

glad he pointed out that the temperature rise

question hardly enters into the case of indentations. It need not ; you can certainly do them

very slowly. Essentially, I stand pat on what I know about indentations because that is the

only experiment in which I feel reasonably

confident that I know where the temperature is

at all times. In connection with what Roy Rice

has said, let me remind you that I never said

that magnesium oxide does not plastically deform or that sodium chloride does not and as a matter of fact, I don't say that alumina or even

diamond does not deform if the temperature is

high enough at the right time.

R. W. Rice, Naval Research Laboratory, Washington, D. C:

I'd like to ask Fred Ernsberger what his interpretation would be of the curling chip re- moval we saw in the pointed diamond tool from

the South African film. I think it certainly

seems that one's first thought would be to

think in terms of this as plastic flow. If he has

a different interpretation I would be interested.

F. Ernsberger, PPG Industries, Pittsburgh,

Pa.:

I simply fall back on the fact that I cannot

think of any mechanism for plastic flow in

glass at room temperature but I do know that

it flows at elevated temperature by Newtonian

mechanisms ; so that is my answer.

A. H. Heuer, Case-Western Reserve University,

Cleveland, Ohio:

I'd like to make a comment about this question of irreversible densification. My colleague,

Al Cooper, has spent the last three or four

years studying volume viscosity in B2O3 and

has measured (with M. Mizouchi) the kinetics

of the irreversible compaction that occurs un- der hydrostatic pressure. It is a very complicated process. For example, there is no equilibrium volume at an equilibrium pressure; the

volume depends on the pressure history, and

the concept of a "fictive pressure" (equivalent

to a fictive temperature) is not sufficient. To

my knowledge, there is no analogue in crystalline materials, which, when hydrostatically

compressed, will deform elastically unless a phase transformation intervenes, which can

usually not be quenched. I would tend to agree

with Fred that this phenomena of irreversible

compaction is really a different phenomena

than anything occurring in crystals and it probably can occur only in glasses because of their

open structure and lack of periodicity.

A. R. C. Westwood, R.I.A.S., Martin Marietta

Corporation, Baltimore, Md.:

I must rise to Fred Ernsberger's comment

concerning the apparent elusiveness of the Rebinder effect. When this phenomenon was first

reported in the British scientific literature in

1947 (Rebinder, P. A., Nature (London), 159,

866 (1947) ), several Western scientists tried

to reproduce Rebinder's results. They could not,

and consequently some doubt arose concerning

the existence of this effect—doubt which has

persisted into the 60's and, apparently, even

into the 70's. Yet the facts are that this is a

real effect and, in nonmetallic solids, it can

be a large effect. And, as I have said on many

occasions, I believe that it is an effect whose

technological potential make it well worth

studying.

I am convinced that a "hard" look will not

make this phenomenon disappear. The twentyfold increase in drilling rate of soda lime glass

in heptyl alcohol over that observed in water,

demonstrated in our paper, is too large an

effect to just go away. Moreover, the existence

of this Rebinder effect in glass cannot be ex409

439-921 O - 72 - 28
plained on the basis of the two presently popular theories for interpreting environmental

effects in this class of solids, nor can it be interpreted in terms of lubrication or coolant

effects, etc. On the other hand, it does appear

that a consistent explanation can be provided

on the basis of the hypothesis that environment-sensitive flow can occur in glass at room

temperature, and that the changes in hardness

and machinability observed are related to adsorption-induced redistributions of the sodium

and/or hydroxyl ions in the nearsurface layers

of the glass.

Regarding the occurrence of environmentsensitive flow in glasses, let me mention briefly

some experiments in progress at RIAS involving the slow propagation of stable, centerloaded cracks in soda lime glass. We are finding that the energy required to propagate such

cracks is very sensitive to the chemistry of the

test environment. Indeed, our preliminary re- sults indicate that the fracture energy of this

glass in the alkane environments varies in the

same manner as the hardness of the glass.

That is, the harder the glass, the greater the

energy of fracture. Our present interpretation

of this effect is that slow crack propagation in

glass can be facilitated by localized, and en- vironment-sensitive, flow processes in the vicinity of the crack tip, perhaps in an analogous

manner to that in which dislocations emitted

from the tips of cracks in semibrittle solids cause them to grow to critical size. You will appreciate that in this instance we are talking

about glass specimens tested in tension, not in

compression, where compaction may be a con- fusing variable.

Incidentally, regarding the possible mechanisms of flow in glass, Oilman (Oilman, J. J., in "Dislocation Dynamics", McOraw-Hill, New

York, 3-25 (1968) ) has suggested that if one

will allow the possibility of a nonconstant Burgers vector, then it is possible to interpret the

flow properties of amorphous solids in terms of

the presence and movement of dislocations.

J. B. Wachtman, Jr., National Bureau of

Standards, Oaithersburg, Md. Panel Moderator

Please hold your questions for a while so that

we can go on to our next panel member. Dr.

William Rhodes of AVCO. Dr. Rhodes has

extensive experience with special ceramic

processing techniques designed for control of

microstructure and he will discuss the effect of

microstructure on the machining process.

W. H. Rhodes, AVCO Corporation, Lowell,

Mass.Panel Member:

A number of papers have touched on the

effect of microstructure in polycrystalline cer- amics on grinding and polishing behavior, but

it is clear that we've really just begun to enter

on to an era of a coupled analysis of machining

and polishing behavior and physical ceramics.

It is very difficult for me to try to pull a con- sensus of microstructure effects out of this

meeting. We've talked about so many different

materials and microstructural features were

not systematically varied within a given set of

tests. Therefore, the remarks on microstructural features will be quite general and some of

them have already been stated, but I will state

them again for emphasis. I'd like to touch

briefly on grain size dependence, porosity dependence, second phase or chemical effects, and

last, but not least, the problem that we all face

everyday and that is the effect of nonhomogeneous defects in materials.

The lack of systematic studies of the effect

of changing grain size within a given set of

grinding tests without changing any other variables makes it impossible to state unequivocally

what the effect of grain size will be, so I shall

primarily talk about some very general conclusions. Roy Rice has mentioned one already and

that is the effect of dislocation damage on grain

size. With decreasing grain size you can ex- pect a lower level of dislocation damage within

the material under question. This would probably hold true to some degree in terms of

crack damage as well because you can imagine

that the cracks will certainly be blunted,

twisted, and deflected at grain boundaries and

therefore crack damage would probably de- crease with decreasing grain size. There are many fine finishing (polishing) operations

where we have this problem of grain pullouts

and in many cases the surface finish is important. It is almost impossible to obtain a fine

finish without some pullouts and obviously the

depth of the final flaw will be dependent on

grain size as well. In certain applications, such

as the one that was mentioned this morning by

Bob Stokes (the gyroscope application where

very fine tolerances are required for very small

spaces between mating surfaces), the grain size

will be particularly important especially if

there is a touchdown between mating surfaces

and grains are, in fact, pulled out during use. So grain size will continue to be an important

parameter in those applications. Now the tendency for pullout must be related to the chemistry of the system. Many of the commercial

ceramics talked about in this conference are multiphase systems that certainly have what is

normally called a grain boundary phase or a

glass bond phase and we certainly can expect

there to be quite different bond strengths between the matrix phase and the bond phase for

the various materials under question. Of course,

this would be quite different again if we went

to a truly single phase material. I'd like to

410
Figure 5. Hot pressed MgO fabricated with additive.

(250X)

illustrate this point with two slides. In figure

5 we have a material that really hasn't been

talked about too much in terms of fine finishing

and that is polycrystalline magnesium oxide.

This is a material that was made without any

deliberate additive in the hot pressing process.

As you can see, a very good polish was obtained

with very little grain pullout. Figure 6 will

show magnesium oxide of similar grain size

that was made by the lithium fluoride additive

process. You can see a very large degree of

pullout in this particular material. This illustrates that two materials that appear similar

under an optical microscope can behave quite

differently in the final polishing.

Porosity has not been studied with systematic changing of variables. Keeping chemistry

and grain size the same while varying porosity

is a very difficult study to make as a matter of

fact. However, I think we can draw the general conclusion that surface roughness would

increase with increasing porosity. Also, as you

reach very high porosity, you can expect a lot

of chip out around large pores, and the actual

loss in pullout density will go up with decreasing porosity.

For the most part, we have talked about materials as if they were truly homogeneous. I was glad to see the papers today deal with the

Figure 6. Hot -pressed MgO fabricated with LiF additive. {150X)

real world of nonhomogeneous materials. I think this is what we are all striving to eliminate, but a truly homogeneous material within

a one or two cubic inch specimen is the exception rather than the rule in today's market, I'm

afraid. Anyone who is connected with processing is daily trying to eliminate defects. We saw

examples of defects this morning arising from

agglomerates, from impurities, and I can go

through a list of causes that we at AVCO have

come up with. We've found defects from the

above causes and incomplete mixing of a deliberate dopant. You can very often have inhomogeneities arising from this. You can have bits

of a wrapper from your powder container that

would actually get compressed in your material

and during firing leave a void in your final

piece—bristles from brushes that people use to

clean screens—it goes on and on. Actually, towards the end of the week you might have to

worry about dandruff. So it's a never ending

battle, and something that all of us need to be

aware of when we're describing surfaces and

what we can expect as a final surface. We have

to be realistic and acknowledge that there are

these large defects that very well may be the

limiting step in terms of obtaining an optimum

fine surface.

411
In polycrystalline materials such as aluminum oxide, we also have anisotropy effects that

certainly must come into play in the ability to

maintain a surface finish. Anisotropic thermal

expansion is a function of crystal habit that

certainly must leave different retained strains

in a body at different grain sizes under different processing conditions. So this well could be

a variable that might be important in surface

finishing. There are fabrication processes which

will give anisotropic materials a preferred crystallographic orientation and these could have

quite different finishing characteristics than

randomly oriented material of the same grain

size.

I would like to mention several summarizing

guidelines for possible future research on the

effect of microstructure on grinding and polishing. We do need a systematic study of grain

size and porosity in grinding performance.

Some sort of fundamental studies of chemistry

on multiphase bodies and the bond strengths

and how these relate to grinding and polishing

behavior is also warranted. I'll mention again

for emphasis that I think all of us need to eliminate the nonhomogeneous defects that are in

real materials. Bob Stokes mentioned one re- search area that he wanted to see continued,

and I would like to second the appeal for a

systematic study of the depth of crack damage

on surface finishing and polishing.

One area that doesn't really relate to microstructure but was extremely interesting to me

was the fact that you can get quite different

surface finishing effects by varying the type of

diamond grit one uses, and I'd like to see and

encourage our South African friends to con- tinue that body of knowledge on all grades of

commercial material.

L. Berrin, Bell Telephone Laboratories, Allentown, Pa.:

There are two points, mentioned by Bill

Rhodes, that we should certainly consider.

First, for the most part, characterization of the

material that we are investigating is not well

defined. Many of us have heard this mentioned

before, but it needs to be reemphasized until

something is done. Also mentioned was the

need of a systematic study of microstructural

features, e.g., grain size and porosity, as a

function of machining or polishing parameters.

This, of course, is necessary, but in lieu of this

it would at least be desirable to report the grain

size and porosity of the material under investigation at the time we present our research.

Other characteristics of the material that we're

investigating, such as its chemistry or method

of forming, would also be helpful. I think this

criticism of data presentation has come up year

after year, but I think it's something that can- not be overlooked.

J. B. Wachtman, Jr., National Bureau of

Standards, Gaithersburg, Md.:

That takes us back to some of the earlier

comments by members of the audience who are

familiar with the machining of metals and who

have suggested a standard test or a standard

method of presentation of data. Perhaps they

would attempt to answer this question: If I come to you for advice on how to grind a given

material, what information do you want from

me? What do you want to be told about the

material? Do you want to know its hardness,

its porosity, or what are the factors you want

to know before advising on the procedure for

machining a given material?

H. C. Miller, Super-Cut, Inc., Chicago, 111.: Just the things you'd expect. Metals are well

categorized to ASTM and SAE standards. A

description of chemical composition of the

metal would be a necessity; equally important

are hardness measurements and microstructures. There are standard descriptions available for microstructure and also, of course, for

porosity. I haven't heard any mention of, for

instance, use of the standard system for talking

about things in terms of A-1 or A-2 porosity.

Chemical composition can be described as 1040

steel or something of this sort; microstructure

would be in terms of grain size measured by

standard methods, porosity as measured by

standard methods, and hardness by a standard

method. In metals, hardness and tensile

strength are usually very closely related so you

don't always need both but can use one or the

other. At least that much information is

needed before anyone could tell you anything

about or give you a recommendation on metal

removal.

J. B. Wachtman, Jr., National Bureau of

Standards, Gaithersburg, Md.

Is the gross chemical composition sufficient or do you need additional information such as the heat treatment and state of precipitation?

H. C. Miller, Super-Cut, Inc., Chicago, 111.: Yes, usually the heat treatment information

is desirable. Heat treatment is how you arrive

at the condition of the final material and sometimes, that is of additional value. But if you

have any other information for instance, microstructural data that will be helpful also. The

heat treatment is especially important if it's

the type of metal in which phase changes oc- cur, something like titanium. You'd have to

know heat treatment to know whether you had

alpha or beta phases present.

412
J. A. Mueller, Carborundum Company, Niagara Falls, N. Y.:

I would simply add to this that you should

specify your objective; i.e., what you want to

do with the material. Do you want to simply

remove stock or do you want to generate a geometry or do you want to produce a finish or do you want to do all three? State your objective and then we can more logically and effectively select the grinding wheel to do the job.

R. J. Caveny, Diamond Research Laboratory,

Johannesburg, South Africa:

I think even with all these specifications and

our present knowledge, it is going to be very

difficult to get the grinding behavior. I'd like

to quote just one example with tungsten car- bide composites. We ground a certain grade of

tungsten carbide from a certain manufacturer

and we had some pretty good results. We

changed batch and we got completely different

results under the same conditions. We analyzed

both with all the known techniques for cobalt

bonded tungsten carbide and found no differ- ence. We got a piece of it from a different manufacturer. Here again, we found the same

physical properties but again found completely

different grinding behavior. I think it boils

down to the fact that at present we don't

really know enough, and this is what Dr.

Gielisse has said. We must investigate the

fundamentals of grinding to elucidate techniques so that we can do these predictions.

Otherwise we'll just have to do critical tests.

P. J. Gielisse, University of Rhode Island,

Kingston, R. I.: I do agree. Every time the question is asked

as to what you have to know in order to define

the system for grinding, one gets at best some

remarks with regard to materials being ground

upon. One simply forgets the order of magnitude changes that you can cause in the actual

operation by changing parameters which determine the performance of the wheel. There is

sometimes the problem of pullout. The grinding

system is being regarded as something that is

fixed. I mean it is considered to be defined for

us. You observe pullout, what can we do about

it? We should really turn the question around

and ask if the pullout can be prevented ; it can,

indeed, be entirely prevented if you simply

tailor the grinding system to what you are trying to do. Another matter I wanted to comment

on is this attempt to put things on a uniform

basis by using the same grinding conditions.

There can be no such situation as "the same

grinding conditions." For each material, for

each objective, there are different grinding conditions. That's why these bar charts which

compare grinding on the basis of 3000 ft per

min or for another particular boundary condition are essentially useless. We have to move

beyond this.

R. J. Caveny, Diamond Research Laboratory,

Johannesburg, South Africa:

I'd like to make a comment on that. For instance, just consider our results that I quoted

for grinding tungsten carbide—completely different behavior from three different ceramics.

If one was optimizing the grinding conditions

one wouldn't use the same conditions for the

three different ceramics. A standard test is going to be very difficult,

H. C. Miller, Super-Cut, Inc., Chicago, 111.: There is no such thing as a standard test

and that's what I have against attempting to

classify grinding conditions. I know this comes

up because people want to compare things. It

is very nice, particularly for engineers, to draw

curves through points, but the thing is really

that you can only standardize, for instance, with

things such as the mode of grinding; shall we

do it as surface grinding or not, what's the diameter of the wheel, and things like that. But

each material very definitely is going to react

very differently.

J. B. Wachtman, Jr., National Bureau of

Standards, Gaithersburg, Md. Panel Moderator:

Our final panel member is Professor Arthur

Heuer of Case-Western Reserve University

who has done extensive research on the mechanical properties of ceramics including sur- face effects and temperature effects. He will

discuss the effect of heat treatment on ceramic

surfaces and mechanical properties.

A. H. Heuer, Case-Western Reserve University, Cleveland, Ohio—Panel Member:

I must thank Peter Gielisse, because everything I have to say is going to sound mild, gentle and reasonable. I was asked to comment on

the effect of heat treatment, which has been

discussed by a number of conferees. One of the

things that impressed me very much was the

difference between single crystals and polycrystals. I'm thinking particularly of the work

on alumina polycrystals by Sedlacek and by

Kirchner on the one hand, and the work by

Noone and Heuer on single crystals on the

other. The strength increases that were

achieved for polycrystals were, in general, very

small, plus or minus 20 percent. (The only ex- ception is very severe cracks introduced by

thermal shock damage or by very deep and

severe cracks, which are somewhat artificial be- cause these can usually be avoided in practice.)

In contrast, large strength increases have been

obtained in alumina single crystals. The inter413
esting question is what is causing the differ- ence. I think there is one factor which has been

overlooked in the past, but was mentioned very

briefly a few minutes ago by Bill Rhodes, viz.

internal stresses in polycrystals due to thermal

expansion anisotropy. I think that in the case

of an anisotropic material like alumina, these

probably play a more important role in determining the strength of polycrystals than is

generally realized. Part of the difficulty is that

at a zero order approximation, one concludes

that the internal stresses should be independent

of grain size and yet there is very clear evidence that there is something like a Griffith re- lation observed between the strength of polycrystalline alumina and its grain size. As a

corollary question, why does the strength of

polycrystalline alumina never exceed about 100,-

000 psi, even for fine-grained polycrystals? To

understand the differences between single crystals and polycrystals, we have to understand

(on a fundamental level) what are the factors

controlling the strength, how these interact

with the damage introduced during surface

machining, and how they are affected by heat

treatment. (The importance of machining damage in certain cases was illustrated very nicely

in Bob Stokes' paper, where he showed that

the body stresses introduced by machining can

strongly influence electrical and magnetic properties.) Getting back to the eflfect of heat treatment, one of the significant new results was

reported by Rice, Becher, and Schmidt, who

showed how, using a helium-freon etching

"heat treatment", one could dramatically improve the strength of single crystals of arbitrary shapes. As was pointed out, flame polishing of single crystals is rather restricted to

simple shapes. The fact that it had no influence

on the strength of polycrystals (or only a very

limited effect) relates to the point discussed

above.

Another very interesting observation was

given by Sundahl and Berrin, who showed that

impurity segregation to the surface during annealing had a very beneficial effect, in that the

impurities enhanced metalizing. This is clearly a subtle effect, but one that obviously has very

practical applications in many modern electronic uses of ceramic substrates.

There is one further point I want to mention,

with respect to the environment during annealing, particularly the peculiar effect of hydrogen.

Berrin told me in conversation that this is also a factor in his work; there are some funny

things which happen when he heat-treats alumina in hydrogen. Again, this is an area where

we need further understanding, particularly

since furnaces for use above 1500 °C often operate in reducing atmospheres. If there are some poorly understood effects and interactions

occurring in ceramic bodies in hydrogen or vac- uum at elevated temperatcres, continued use

of such environments may be questioned.

Finally, I was impressed by the Bell Labs

paper this morning showing how well surfaces

can be characterized using several new instrumental techniques—Auger electron spectroscopy, low energy electron diffraction, and

reflection electron diffraction—as well as the

better known scanning electron microscopy. I think these techniques can be used very well

to understand changes occurring during heat

treatment. For example, Schmidt and Davey

demonstrated using reflection electron diffraction that the surface layer of mechanically polished sapphire crystals was amorphous. However, this layer could be recrystallized by heat

treatment, and this represents a distinct improvement in our understanding of what happens when machined or polished surfaces are heat treated.

H. P. Kirchner: Carmic Finishing Company

State College, Pa.:

I agree with Dr. Heuer about the importance

of expansion anisotropy, and one thing that

has been clear to me for sometime is that strong

ceramics of phases with anisotropic crystals

could not be made if what is important is the

stress because the stresses are undoubtedly

very high. But what is important is the residual

deformation in an individual grain and if this

is small as it is in a small grain size material

then it is not possible for a crack to open up.

The principal distinction is between the total

deformation of the individual grain (inches per

grain) and the strain (inches per inch).

A. H. Heuer, Case-Western Reserve University,

Cleveland, Ohio:

I'm afraid I don't understand that point. It seems to me in a zero order approximation

that the strain is independent of grain size;

strain, and hence stress, result from change in

length per unit length and this is independent

of grain size. I think it's a problem a number

of people have struggled with and it seems to

be particularly important at this point in time.

R. C. Sundahl, Bell Telephone Laboratories,

Inc., Allentown, Pa.:

Just another point on the effect of anisotropy

I briefly mentioned this morning some measurements of texture. It was found by Y. Nakada

in our laboratory and P. L. Key, F. V. DiMarcello and J. C. Williams at the Murray Hill location of Bell Labs that there is wide variation

in fiber texture in tape case substrates obtained from a variety of suppliers. If the anisotropy of the thermal expansion coefficient were

important one might imagine there would be

414
wide variations in mechanical stren^hs of these

same materials from batch to batch and from

material to material. Another point is that

although there seems to be some correlation

between mechanical properties and grain size,

it is dangerous to try to correlate one parameter solely with another, for example, mechanical propei'ties with grain size, because there are a large number of parameters which have

an influence both on microstructure and mechanical strength. One such parameter is the

chemistry of the ceramic. It is not clear which

of these various parameters dominate and how

many of these observed effects are only indirect

symptoms of the dominant effects.

A. H. Heuer, Case-Western Reserve University,

Cleveland, Ohio:

I agree completely. I think it's a mistake to

think of mechanical strength, say of alumina,

as a function only of grain size. It depends on

several factors and that is the point I was

trying to make.

A. U. Daniels, Metcut Research, Cincinnati,

Ohio:

One comment I can't resist concerns this busi- ness or the use of standard methods of characterization or testing. I think there is fear that

if people use standard methods they will stop

thinking. I believe that's probably the fear con- cerning standard machining procedures ; there's a feeling that maybe engineers do that; they

just look at standard test results and then turn

their minds off and treat the material as being

fully described. I think instead that any kind

of standard methods, if established, could serve

to keep people thinking. Probably we have had

an excellent example here at the conference.

People are using profilometers and at the same

time saying, "Look here, the profilometer says

this, but when you look at the material under

a scanning electron microscope you see that the

materials can be very different with the same

profilometer readings." I think that standard

methods can do two things. They can serve as a bench mark or background from which people

can extend their thinking. Standard tests are always inadequate, they always miss something,

but they do serve as a bench mark or background against which you can begin to make a

sensible comparison of materials. They also do

another thing. There has been talk here about

an information gap between applied and basic

information and those standard test results can

serve to bridge that gap. So that is the reason

I make a strong point that standard tests,

rather than restricting thinking, can serve as a background and a basis for which better description and better communication can take

place.

P. J. GiELlSSE, University of Rhode Island,

Kingston, R. I.: Just a brief comment on that. Being on the

faculty of a college of engineering, I certainly

do not support, of coui'se, the idea that engi- neers stop thinking, and I believe we have to

interpret that quite differently. Secondly, I have

the greatest esteem for the efforts, of particularly your firm, in the area of standardization.

You do mention something, the use of the profilometer and I entirely agree with your comments. People are not using the profilometer

correctly now. They are taking roughness readings, but what can be done with profilometry is

much more than is presently being practiced.

I wish that this would be noticed and hopefully

changed in the future.

A. H. Heuer, Case-Western Reserve University,

Cleveland, Ohio:

I would like to make a rather philosophical

comment, being also on the faculty of a school

of engineering. In reading the ONR-London re- ports in the last couple of years, I've been

struck by their reporting of the establishment

of several departments of tribology in the

United Kingdom. An entire academic discipline

involved with machining and grinding of materials is being established, chairs are being

created, and students are being graduated with

sophisticated training in this particular discipline. To my knowledge, with perhaps the ex- ception of Peter Gielisse's group at the University of Rhode Island, there is very little like

that done in the United States. As an academic,

I can assure you that what is required here is

money—money from the Government, and also

perhaps more importantly, money from industry (the users and producers of ceramics and

ceramic machining wheels). My impression is

that there is only meager support by industry

for the academic disciplines with which they

are most concerned. This communication gap

that has been alluded to is perhaps a result of

insufficient people being formally trained in the

appropriate discipline, and I would hope that

industrial firms could take the lead in effecting a change.

R. W. Rice, Naval Research Laboratory, Washington, D.C.:

I'd like to make two comments tying in with

what both Art Heuer and Bill Rhodes said.

First, concerning porosity, I wonder if it's possible to deposit (e.g., during grinding) sufficient

swarf within existing pores in the material so that if you then anneal the material, you'll get

one result; whereas under other grinding conditions, you may not deposit swarf in the pores,

and hence your annealing results could be very

different. Also, any cleaning process prior to

415
annealing could change the situation. Variables

such as this should be kept in mind in looking

at annealing effects. Second, I would very much

agree with Art Heuer that mechanical properties do not depend entirely on grain size, but

I think we must recognize that grain size is an

extremely pervasive variable. We also have to

recognize that the effects of grain size can be

substantially changed. I think, for example, the

data we showed on grinding direction shows

this rather markedly: if you plot strength ver- sus grain size for specimens ground in different

directions, you'll get rather different curves. I would go along with the view that you have to

be careful in describing strength as a function

of grain size, but I think it has to be recognized

that grain size is a very pervasive variable.

R. F. Firestone, Case Western Reserve University, Cleveland, Ohio:

I'd like to try to put Professor Heuer's re- marks in historic perspective. In 1880, Frederick Winslow Taylor, the Father of Scientific

Management, began to study the art of cutting

steel with single point tools at the Midvale Steel

Company which later became part of the United

States Steel Company of Andrew Carnegie.

Taylor thought the study would take less than

six months to complete.

Four years later he was still making experiments and had hired a graduate engineer to

help with the mathematics. Twenty-six years

later, after having reduced over 40 tons of steel

to swarf, he had developed a formula to determine the optimum cutting conditions and he

had developed the modern high speed tool

steels. The results justified the effort.

The point I'm trying to make is that it took

a great deal of time and money to discover the

best way to machine low carbon steel, a homogeneous material, with a single point tool. Ceramics and the process of grinding are much

more complex. I don't think there is any organization which is willing to put in the time and

money required to solve the problem. This con- ference is certainly a step in the right direction

but I am pessimistic about rapid progress being

made.

P. J. GiELiSSE, University of Rhode Island,

Kingston, R. I.: That's correct. I do think that the present

concept of our university administrations cer- tainly doesn't lend itself to continuing, longterm support of this nature, but I think the

comments you make are in line with what Professor Heuer is saying. What we need is further interaction between the academic world,

that can provide that sort of information, and

the industrial users.

NATIONAL BUREAU OF STANDARDS SPECIAL PUB. 348, The Science of Ceramic Machining and Surface Finishing,

Proceedings of a Symposium Sponsored by the American Ceramic Society, the Office of Naval Research, and the National Bureau

of Standards, held at NBS Nov. 2-4, 1970, Gaithersburg, Md. (Issued May 1972).

416
Author Index

Page

Allgeyer, G. H 53

Becher, P. F 237, 267

Berrin, L 271

Bersch, C 1 Bunnell, C. R 341

Busch, D. M 73

Campbell, W. B 133

Caveney, R. J 99

Colwell, L. V 53

Crowe, J. C 341

Davey, J. E 259

Diness, A. M 1, 309

Douglass, J. W 189

Emery, J. K 293

Ernsberger, F 401, 405

Feick, G 197

Finnie, 1 119

Fujino, S 37

Gielisse, P. J 5, 401

Gruver, R. M 353

Hahn, R. S 59

Halden, F. A 89, 391

Hart, P. E 341

Heuer, A, H 213, 401, 413

Hockey, B. J 333

Horowitz, E 1 Imanaka, 0 37

Jorgensen, P. J 89, 391

Kirchner, H. P 353

Koepke, B. G 317

Lange, F. F 233

Page

Latanision, R. M 141

Lee, D. W 197

Leistner, H. C 113

Levy, P. W 155

Lindsay, R. P 59

Lo, W. C 301, 399

McKinstry, H. A 309

Mineta, S 37

Mueller, J. A 45

Noone, M. J 213

Oh, H. L 119

Oh, K. P. L 119

Pollock, J. T. A 247

Prins, J. F 73

Rhodes, W. H 401, 410

Rice, R. W 1, 171, 193, 237,

267, 365, 401,

407

Schmidt, W. A 259, 267

Sedlacek, R 89, 391

Stanislao, J 5

Starrett, H. S 377

Stokes, R. J 343

Sundahl, R. C 271

Thiel, N. W 99

Vaidyanathan, S 119

Wachtman, Jr., J. B 401

Walker, R. E 353

Westwood, A. R. C 141

White, E. W 309

Wilson, W. A 113

417

Subject Index

A

Abrasion, 73, 120, 355, 356, 360, 362

Absorptivity, 194

Acoustic emission, 341—342

Adsorption, 141-152

AI2O3, see Aluminum oxide

Alumina, see Aluminum oxide

Aluminum oxide, 10, 15, 23, 40, 89-90, 99-112, 233-236, 273,

276, 282-283, 285, 288, 291, 302, 304, 333-339, 353-363,

377-378, 383, 385

diamond polished, 326

polycrystals, 325

sapphire, 237

single crystal, 213-214, 226, 267

Amorphous, 260

Angle

cutting edge, 136

rake, 136

relief, 136

Annealing, 213, 228-230, 354-355

Arc,

coupling, 194

electrode, 194

Auger electron spectroscopy, 279-283, 291

B

Bend testing

at elevated temperatures, 225—226

3-point, 222

4-point, 222-223

Binary maps, 311

Bonded Abrasive Machining (BAM), 14-17 Boron carbide, 326

Brittle-ductile transition, 119, 126-127

Brittle solids, 119

C

CaF2, 142 Camber, 297

Carbides (see also specific material), 368

Ceramics, 45-47, 51, 141, 144, 197, 349-351

characteristics, 7 machining, 200, 204, 207, 343

properties, 24

CESEMI, 309

Characterization, 401-402, 411-412, 414

Charles-Hillig theory, 405

Chatter vibrations, 81-82

Chemical reaction, 12 Chip

accommodation, 54

brittle, 54

continuous, 41 fragment, 40

Computer analysis, 311

Contact length, 62

Coolant,

effectiveness, 56

water soluble, 57 Crack healing, 233-235

Cracking ring, 121, 131

Cracks, 125, 233, 344

Crystallographic texture, 285-288

Cutter, single point, 7, 22, 57, 73

Cutting, 402

diamond, 73

edge angle, 136

environment, 57

fluid, 402

D

Damage

edge, 399

layer, 42

Densification, 406, 408

Depth of field, 309-310

Depth of slip from machining, 371

Diamond abrasive system, 6

Diamond grinding, 6, 73-87, 391-394

description, 14—18

strength, 30

types, 15 Diamond tools, 52, 54-55, 74

grain size effect, 56

performance, 57

single layer wheel, 55

wheel, 99-112

Dislocations, 343-348

in AI.O3, 333-339

in rutile, 374-375

Drilling, 141, 144-147, 193

Dulling, 56

E

Economics, 21 Edge damage, 399

Electric discharge machining, 198-203

ceramics, 200—203

erosion rates, 199

nonconducting ceramics, 202—203

refractory ceramics, 201

spark discharge ceramics, 200-201

Electrochemical machining, 203-204

ceramics, 204-205

Electron beam

air, 194

coupling, 194 vacuum, 194

Electron diffraction low energy, 285, 290

reflection high energy, 285—287

striations, 260-261

Electron microprobe, 278-279

Electronic circuits, hybrid and microwave.

Ellipse fit, 311

Environmental effects, 141, 403, 409

Epitaxial, 259

Erosion, 120, 155

Ion beam, 156

analysis, 123-127

glass, 124

graphite, 124

magnesia, 124

Etch pits, 321

Etching, 406

thermal, 359

F

Fatigue, static, 354

Feasibility study, 309

Ferrite, 324, 349

Figuring, 155

computer controlled, 157

ion beam, 156

Finishing, 113-117, 401-402

theory of, 21-22

Flame polishing, 213, 237, 353

batch, 247

continuous, 247

glass, 242-243

rutile, 242

sapphire, 237-241

419
spinel, 241

strengths, 238, 241

surfaces, 238-241

techniques, 214-219

Flaws, 121, 356, 362

Fluorotrichloromethane, 260

Forsterite, 39

Fracture, 29-32, 146, 317, 407

brittle, 53

delayed, 405

energy, 409

grain boundary, 337

location, 119, 122

mirrors, 267-268, 366

origins, 267, 366

toughness, 130

under spherical indentures, 120

Freon, 260

etch, 262

G

Glass, 54, 141, 148-152, 155, 341-342, 366-367

photosensitive, 206

Glass bubbles, 243

Glass ceramics, 53—54

Glass phase in ceramics, 405, 410

Glassy carbon, 366—367

Glazing. 357, 362

Grain boundary, 345

Grain growth, 236

Grain size, 235

Griffith flaws, 371, 407

Griffith theory, 31

Grinding, 13, 114, 365

centreless, 16

ceramics, 5, 7, 90

costs, 109, 111

cylindrical, 16

direction, 110-111, 365, 372

dynamic aspects, 22

efficiency, 15, 100, 106-107

forces, 23-32, 59, 73

internal, 17 model, 29

modes, 20

noise, 81

principles, 59

rutile, 374

stria, 368, 372

surface, 17

techniques, 402

thermal aspects, 22, 32-34

tool and cutter, 17 H

Hardness, 151, 367, 407, 412

anisotropy, 376

Marsh's theory, 371

Heahng, 356, 357, 360, 362

Helium, 360

Humidity, 353

Hydrogen etch, 261-263

Hydroxide, 260

I

Impact forces, 29

Intermetallic compounds, 201

Internal stress, 232

Infra red methods, 34

Ion beam, 156

bombardment, 155, 333

thinning, 334

Ion microprobe, 284, 290

L

Lapped surfaces, 299

Laser scribing, 399

Lucalox, 312, 341

M

Machinability, 33

Machining, 2, 141, 345, 401-402, 407

abrasive, 14-21, 53, 133, 233

arc, 193

ceramics, 133, 343, 346

chemical, 205

costs, 2

electrical discharge, 198

electrochemical, 203

electron beam, 193

glass ceramics, 53

laser, 193

methods, 2

sonic, 133

sputter, 158

trends, 3

ultrasonic, 17 Magnesia, see Magnesium oxide

Magnesium oxide, 11, 142, 280, 281, 291, 314, 410

crystals, 320

polycrystals, 324

Material removal, 53

Material removal parameters, 60

Mechanical methods of machining, 5 Mechanical and other effects of finishing, 343

Mechanical testing of ceramics, 245

Mechanisms of machining, 401

Melting, 194

MgO, see Magnesium oxide

Microcracks, 235

Micro-flash technique, 37

Microplasticity, 406-408

Microstructure, 95, 383, 410-412

Modulus of resilience, 31 Molybdenum disilicide, 201

N

Nitrides, 368, 369

Nonmechanical methods of machining, 155

Nose radius, 136

Novaculite, 316

O

Optical

fabrication, 189

microscope, 274—275

scatter, 191

surfaces, 155, 189

Optics, 155

Oxides (also see specific material) , 366, 367, 368, 369

P

Panel discussion, 401

Particle size distribution, 129

Phosphoric acid, 260

Photoelectron spectroscopy, 283-284, 290

Plastic deformation, 214, 226, 317, 405, 406, 408

in AI2O3, 333, 337

Polishing, 13, 155, 333-334

chemical, 353, 359

of B4C, 268

of MgALOi, 268

of MgO, 268

of polycrystalline AI2O.1, 267

of rutile, 268

of sapphire, 267

of SiC, 268

of ZrO^, 268

using hot gases, 267

flame, 213, 237, 247, 260, 353

gas, 245

ion beam, 156, 189

ionic, 190

420
Processing variables, 384

Properties

electrical, 349

magnetic, 348

mechanical, 346

optical, 351

Pyroceram, 341

R

Radiation damage, 155, 161 Rebinder effect, 141-142, 406, 409

Residual stress, 11, 334, 337, 348, 349, 362

Roughness, 48, 297

Rutile, 268, 374

S

Sanding direction, 370, 372

Sapphire, (see also Aluminum oxide), 216, 220, 225, 232, 237,

259

filament, 249, 251

Scanning Electron Microscopy (SEM) 102, 272, 275-277, 290,

302, 310, 333, 336

Science of finishing, 21 Secondary electrons, 310

Semiconductor, 259

Shadow effects, 310

Shaping, 45, 155-156, 193

ion beam, 156

SiO., 287, 291

Si,N4, 368, 369

Silicon, 313

crystals, 325

Silicon carbide, 12, 15, 18, 233, 234

Single point cutting, 7, 22, 57, 73

Size effects, 120, 126

Sizing, 45, 53

Sonic machining rates, 137

Sonic motor, 133

Specific energy, 56

Sputtering, 155

atomic process, 157, 171

efficiency, 159

etching, 171, 182, 185

ion beam, 156, 171, 185

rate, 180

r.f., 173, 185 topography, 176, 178, 182, 185

yield, 159

Standards for grinding, 404, 412, 414

Steatite, 40

Stockremoval

processes, 5, 13, 20, 401

rate, 19, 20

Strength, 216, 365, 412, 413

alumina, 93-94, 230, 377, 383, 394, 395

abraded, 338, 356

sapphire, 226, 227, 269

brittle (statistical theory), 121, 126

ceramics, 11, 12, 24, 241

effect of surface perfection and internal defects, 214, 216, 223

flexural, 353, 355, 356, 357, 360, 362, 378, 383, 384, 385

produced by chemical polishing, 268

produced by flame polishing, 222, 238

sputtered surfaces, 181, 185

tensile, 93-94, 226, 227, 353, 378, 384, 385, 394, 395

Stress concentration, 357, 359

corrosion, 405

localized, 362

residual, 334, 338, 348, 349

yield, 370

Stress intensity factor, 130

Stylus instrument, 272-274

Sub-surface

damage, 317, 333, 334-337

grinding, 334

Substrates, 113-114, 271-272, 293

alumina, 272, 301

Surface

analysis, 8, 9

annealing, 228

characterization, 309

chemistry, 272, 278-284, 288-289

condition.. 317

contamination, 215, 224

crystallography, 272, 285-289

damage, 9-10, 99, 101, 233, 302, 317, 341

defects, 272, 274, 275, 302

degradation, 216

dependent properties, 2 energy and power, 55, 142

features of flame polished rods, 222

finish, 2, 66, 111, 377

grinding, 99-109, 317, 333, 336

integrity, 68

perfection, 215

replica, 301

roughness, 309

strength, 215

stress concentrations, 371

testing, 9

texture, 300, 303-306

topography, 2, 272-278, 289

treatment, 377, 383

Surface and subsurface analysis and characterization, 271

T

Techniques, non conventional, 17 Techniques and mechanisms of removal and shaping, 5 Techniques and mechanisms of surface finishing, 213

Template cutting, 138

Thermal expansion anisotropy, 414

Thermal processes, 22

Thermal shock, 193-194, 233, 358-360

Thick fihn, 299

Thin film, 300

adhesion, 271, 288

circuits, 271-274, 289

Tool shapes, 135-136

Tools, single point 356

Transmission electron microscopy, 277—278, 290, 333-334, 337

Tribology, 415

Tungsten carbide, 204

Twinning

rutile, 242, 244-245, 268

sapphire, 239, 244^245, 267, 333, 337, 372

U

Ultrasonic

cutting, 120

machining, 119

V

Vaporization, 106, 107 W

Waviness, 296

Wear mechanisms, 105 Weibull

distribution, 122, 378

parameters, 123, 126

statistics, 378-380, 285

Wheel

grinding, 45—52

structure, 54—55

wear, 63—69 wear parameters, 60-63

Y

Yield stress, 370

Z

Zirconium carbide, 201, 204

Zirconium diboride, 201, 204

ZnO, 234, 235

1

421

FORM NBS-1 14A (1-71)

BIBLIOGRAPHIC DATA

SHEET

1. PIIRl rCATION OR RFPORT NO

NBS SP-348

2. Gov ' t Ac cess ioo No.3. Recipient's Accession No.

4. TITLE AND SUBTITLE

Science of Ceramic Machining and Surface Finishing

5. Publication Date May 1972

6. Performing Organization Code

7. AUTHOR(S)

S. J. Schneider, Jr. and R. W. Rice

8. Performing Organization

9. PERFORMING ORGANIZATION NAME AND ADDRESS

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, D.C. 20234

10. Project/Task/ Work Unit No.

11. Contract/Grant No.

12. Sponsoring Organization Name and Address

The American Ceramic Society, 4055 No. High Street, Columbus,

Ohio 43214

Office of Naval Research, Arlington, Virginia 22203

National Bureau of Standards, Dept. of Commerce,

13. Type of Report & Period

Covered

Final

1 A on ^orinp Appnfir ( rvri <=

15. SUPPLEMENTARY NOTES Washington, u. 20234

16. ABSTRACT (A 200-word or less factual summary of most significant information. If document includes a significant

bibliography or literature survey, mention it here.)

This volume presents the proceedings of the Symposium on the Science of Ceramic

Machining and Surface Finishing held at the National Bureau of Standards in

Gaithersburg, Maryland on November 2-4, 1970. The symposium was jointly sponsored by

the Baltimore-Washington Section of the American Ceramic Society, the Office of Naval

Research, and the National Bureau of Standards, The purpose of the conference was to

survey the developing science of ceramic machining and to stimulate further progress

by discussion of current problems and research programs. In addition to a panel

discussion, 37 review and original research papers were presented with attention

focused on four main areas: (1) Techniques and Mechanisms of Removal and Shaping

(2) Surface Treatment (3) Analysis and Characterization of Machining Effects and

(4) Mechanical and Other Effects of Finishing. An edited version of the floor

discussion following each paper is given.

17. KEY WORDS (Alphabetical order, separated by semicolons)

Ceramics; ceramic machining; mechanical effects of machining; removal and

shaping of ceramics; surface treatment.

18. AVAILABILITY STATEMENT

HcI UNLIMITED.

I I FOR OFFICIAL DISTRIBUTION. DO NOT RELEASE

TO NTIS.

19. SECURITY CLASS

(THIS REPORT)

UNCL ASSIFIED

20. SECURITY CLASS

(THIS PAGE)

UNCLASSIFIED

21. NO. OF PAGES

431

22. Price

$5.25

USCOMM-DC 66244-P71

U.S. GOVERNMENT PRINTING OFFICE: 1972 0—439-921