DATE DUE

NUV o

/^ o ^ 983

; MAY 1

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1990

GAVLORO PR (NT CO IK U.SA.



A111D3 07E1S3

NATL INST OF STANDARDS & TECH R.I.C

A1 11030721 53

Symposium on the Sol/The science of cera V QC100 .U57 N0.562, 1979 C.2 NBS-PUB-C 19

(A o NBS SPECIAL PUBLICATION 562

U.S. DEPARTMENT OF COMMERCE / National Bureau of Star.da •-01 Admin. Bldg.

OCT 1 1981

-1 S j. 0 8 3

QC

100

U57

0.562

1979

:.2

The Science

of Ceramic Machining

and Surface Finishing II
NATIONAL BUREAU OF STANDARDS

The National Bureau of Standards' was established by an act of Congress on March 3, 1901.

The Bureau's overall goal is to strengthen and advance the Nation's science and technology

and facilitate their effective application for public benefit. To this end, the Bureau conducts

research and provides: (1) a basis for the Nation's physical measurement system, (2) scientific and technological services for industry and government, (3) a technical basis for equity in trade, and (4) technical services to promote public safety. The Bureau's technical work is per- formed by the National Measurement Laboratory, the National Engineering Laboratory, and

the Institute for Computer Sciences and Technology.

THE NATIONAL MEASUREMENT LABORATORY provides the national system of

physical and chemical and materials measurement; coordinates the system with measurement

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standards, and data on the properties of materials needed by industry, commerce, educational

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provides engineering measurement traceability services; develops test methods and proposes

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The Institute consists of the following centers:

Programming Science and Technology — Computer Systems Engineering.

'Headquarters and Laboratories at Gaithersburg, MD, unless otherwise noted; mailing address Washington, DC 20234.

2 Some divisions within the center are located at Boulder, CO 80303.
OCX-- v»c.

The Science of Ceramic Machining

and Surface Finishing II

1 (SXMCO

U6H

Proceedings of a Symposium

held at the National Bureau of Standards,

Gaithersburg, Maryland,

November 13-15, 1978

Edited by

B. J. Hockey

Center for Materials Science

National Measurement Laboratory

National Bureau of Standards

Washington, D.C. 20390

and

R. W. Rice

Naval Research Laboratory

Washington, D.C. 20390

Sponsored in part by

Air Force Office of Scientific Research

American Ceramic Society

Office of Naval Research

U.S. DEPARTMENT OF COMMERCE, Juanita M. Kreps, Secretary

Luther H. Hodges, Jr., Under Secretary

Jordan J. Baruch, Assistant Secretary for Science and Technology

NATIONAL BUREAU OF STANDARDS, Ernest Ambler, Director

Issued October 1979
Library of Congress Catalog Card Number 79-600149

National Bureau of Standards Special Publication 562

Nat. Bur. Stand. (U.S.), Spec. Publ. 562, 524 pages (Oct. 1979)

CODEN: XNBSAV

U.S. GOVERNMENT PRINTING OFFICE

WASHINGTON: 1979

For sale by the Superintendent of Documents, U.S. Government Printing Office, Washington, D.C. 20402

Stock No. 003-003-02122-9 Price $9.

(Add 25 percent additional for other than U.S. mailing)
FOREWORD

Machining and surface finishing operations play a vital role

in the production and reliable use of ceramics. It is essential,

therefore, that the full scope of this technologically important

subject be periodically reviewed and evaluated. One of the most

effective ways of accomplishing this is through meetings, such as this Second Symposium on Ceramic Machining and Surface Finishing.

These meetings provide the necessary forum for the direct exchange

and discussion of new data, recently developed techniques, and

advances in our understanding of underlying principles.

This Symposium, in particular, has focused on both the

scientific and technological advances in ceramic machining and

related topics. The interaction that has occurred at this meeting

between science and technology will provide many direct benefits

and, hopefully, will act as a stimulus for further progress.

Through publication of these proceedings, the information presented

at this symposium, together with any derived benefits, can be

shared with the scientific and industrial community at large.

The National Bureau of Standards is pleased to be responsible

for this publication, and to have joined with the Air Force Office

of Scientific Research, the Office of Naval Research, and the

American Ceramic Society in sponsoring the symposium.

John D. Hoffman

Director

National Measurement Laboratory

National Bureau of Standards

iii
PREFACE

The first Symposium on Ceramic Machining and Surface Finishing

was held eight years ago at the National Bureau of Standards. That

meeting covered, for the first time, various topics on the scientific

and technological efforts that represent the full scope of ceramic

machining and surface finishing. It was, thus, largely responsible

for the development of what is now known as the Science of Ceramic

Machining and Surface Finishing. The resulting publication 1 continues

to be a useful reference on the subject.

The present volume records the proceedings of the Second

Symposium on Ceramic Machining and Surface Finishing held at the

National Bureau of Standards, Gaithersburg , Maryland on November

13-15, 1978. Attendance at this meeting, augmented by attendees

to the concurrently held Fall meeting of the Basic Science Division

of the American Ceramic Society, numbered over two hundred and

included representatives from seven countries including the USA.

The primary purpose of this meeting was the same as that of

the previous meeting, namely: "To survey the developing science

of ceramic machining and to stimulate further progress by discussion

of current problems and research."

In keeping with this theme, the meeting provided not only a timely review of subsequent research on the mechanisms and techniques of machining, surface characterization, and the effects of

finishing on material properties, but also incorporated data on

recently developed materials and information gained from related

research efforts. As an example, new insights into the mechanisms

of material removal during multipoint, fixed abrasive machining

have been gained from recent indentation and solid particle impact

studies. Significant advances in the technology of machining and

surface finishing have also been achieved, in particular, through

the application of single-point diamond machining to ceramics.

Current results for this technique are doubly encouraging in that

they suggest that still further advances are possible. Similarly,

advances have taken place in the rapidly developing area of nondestructive evaluation, and the initial application of techniques

such as acoustic surface wave emission shows considerable promise

in the rapid evaluation of machining- induced flaws. In addition,

the proceedings of this meeting record continuing efforts to better

define the nature of machined or polished surfaces, the effects of

gaseous and liquid environments, and quite importantly, the effects

of machining on the various properties and reliability of ceramics.

Also presented are a number of specific examples which illustrate

the technological progress made in machining and finishing ceramics

for important commercial applications, such as turbine components,

laser windows, and electronic substrates. These examples clearly

reflect the benefits derived from the interaction between technology

and science, and, hopefully, will provide the justification for

further research on machining.

The Science of Ceramic Machining and Surface Finishing, S. J.

Schneider, Jr. and R. W. Rice, eds . , National Bureau of Standards

Special Publication 348, 1972.

iv
Looking to the future, it is hoped that several avenues of

research and development are stimulated by this latest conference

and its proceedings. Although significant progress has been made

in understanding the general aspects of traditional abrasive

machining together with the flaws that are produced, possible new

developments might include: 1) further improvements in machining

technology that lead to substantial reduction in the cost of

finished products (most specifically, in ceramic turbine components);

2) further improvements in the detection and characterization of

flaws (e.g., cracks) and other damage produced by machining; and,

3) a more detailed understanding of the correlation between the

nature and distribution of machining flaws and the properties

(mechanical, optical, dielectric, etc.) of ceramics. With regard

to the latter, mechanical reliability for example, may be improved

by not only limiting the maximum flaw size, but also by deliberately

increasing the density of subcritical flaws. Through multiple flaw

interactions, the effectiveness of any single flaw may well be

limited and resistance to further surface damage improved. In

addition it is felt that the potential suggested by current results

on single point diamond and laser, or laser assisted machining —

as well as other non- conventional techniques — should certainly

encourage further research.

Finally, a few comments on the organization of these proceedings are in order. For the most part, we have attempted to follow

the outline of the symposium program; however, some changes were

deemed necessary to better reflect the content of the papers.

Also, as is often the case, many papers could be placed equally

well in different sections and, hence, section headings should

be regarded as general guidelines. Discussions of the papers have

been included, and in every case were transcribed from handwritten

notes obtained during the meeting. In most cases, these comments

were returned to those involved to verify their accuracy.

Roy W. Rice

Bernard J. Hockey

v
ACKNOWLEDGMENT

The editors wish to acknowledge several organizations and

individuals for their aid in conducting the Symposium and in the

publication of these proceedings. In particular, financial

support was received from the Air Force Office of Scientific

Research, the Office of Naval Research, and the National Bureau

of Standards. The additional sponsorship and organizational

assistance of the American Ceramic Society is also acknowledged.

Similarly, we wish to acknowledge Dr. A. M. Diness, Office of

Naval Research; Major Will Simmons, Air Force Office of Scientific

Research; Mr. Charles Bersch, Naval Air Systems Command; Dr. John

B. Wachtman, Jr., National Bureau of Standards; and, Dr. Benjamin

Wilcox, National Science Foundation, for their roles in developing

the Symposium program. Further, it should be noted that this

symposium was held in conjunction with the Fall meeting of the

Basic Science Division of the American Ceramic Society. Accordingly, we wish to thank Dr. Nancy Tighe, National Bureau of Standards, who as program chairperson for the American Ceramic Society,

assisted in coordinating the scheduling of the two meetings.

Finally we wish to acknowledge the much needed assistance

provided during the symposium and in the preparation of these

proceedings by Mrs. Pellechehia and Mrs. Craig of the Naval

Research Laboratory and Mrs. Salpino and Mrs. Stang of the National

Bureau of Standards. Special thanks are due to Rosemary Maddock of

the National Bureau of Standards for her extensive efforts and

guidance in preparing this volume for publication.

vi
ABSTRACT

This volume presents the proceedings of the Second Symposium

on the Science of Ceramic Machining and Surface Finishing held at

the National Bureau of Standards in Gaithersburg , Maryland, on

November 13-15, 1978. The symposium was jointly sponsored by the

Air Force Office of Scientific Research, the American Ceramic

Society, the Office of Naval Research, and the National Bureau of

Standards. The purpose of the symposium was to review recent

progress by discussion of current problems and research efforts.

The proceedings contain forty of forty-three research papers,

together with floor discussions, presented at the Symposium.

Specific topics covered include: 1) mechanisms of abrasive

machining and surface finishing: 2) technology of machining and

surface finishing; 3) characterization of machined surfaces; and

4) effects of machining on the properties of ceramics. Also

included is an edited discussion on hardness and its relation to

machining

Key words: Abrasion; ceramics; ceramic machining; effects of

machining on ceramic properties; hardness; surface characterizations; surface finishing.

vii

CONTENTS

Page

Foreword . . . iii

Preface iv

Acknowledgment vi

Abstract vii

Session I: Mechanisms of Abrasive Machining and

Surface Finishing

Abrasive Wear in Ceramics: An Assessment 1 A. G. Evans

Nucleation of Median and Lateral Cracks Around

Vickers Indentations in Soda-Lime Glass 15

J. T. Hagan, M. V. Swain, and J. E. Field

Fragmentation and Damage Penetration During

Abrasive Machining of Ceramics 23

H. P. Kirchner, R. M. Gruver, and D. M. Richard

Single-Point Scratches as a Basis for Understanding

Grinding and Lapping 43

A. Broese van Groenou, N. Maan, and J. B. D. Veldkamp

Experiments on High-Speed Grinding of Ferrites

and Glasses 61

A. Broese van Groenou and R. Brehm

Effect of Workpiece Properties on Grinding Forces

in Polycrystalline Ceramics 75

B. G. Koepke and R. J. Stokes

Glass Grinding Mechanisms and Influence on the

Finished Surface and Strength 93

S. Malkin and M. Huerta

The Fine Grinding of Glass and Ceramic Material

Using Conventional Grinding Wheels 107

E. D. Doyle and S. K. Dean

Some Observations on Environmental Effects

when Diamond Drilling 115

G. A. Cooper

The Effect of Chemical Environments on the

Fracture of Ceramic Surfaces 139

R. E. Cuthrell

Some Results on the Wear of a Bronze - Bonded

Grinding Wheel 147

A. Broese van Groenou

New Concepts on Surface Finishing and Its

Application to Ceramics — Recent Progress in

Ultra-Fine Finishing in Japan 157

0. Imanaka and M. Okutomi

ix
Mechanism and Application of the "Mechanochemical

Polishing Method Using Soft Powder" 171

N. Yasunaga, N. Tarumi , A. Obara, and 0. Imanaka

Extended Discussion: Hardness and Its Relation

to Machining 185

R. W. Rice and A. G. Evans

Session II: Technology of Machining and Surface Finishing

Optimization of Multistage Grinding Operations:

The Choice of Consecutive Removal Rates 191

A. Broese van Groenou

Grinding Stone, Concrete and Ceramics with

Coated Abrasives 201

R. C. Lokken and E. J. Duwell

Turbine Component Machining Development 209

D. W. Richerson and M. W. Robare

TM Machinability Studies on Macor Glass-Ceramic 221

D. G. Grossman and D. L. Taylor

Effects of Various Polishing Media and Techniques on

the surface Finish and Behavior of Laser Glasses 231

R. L. Landingham, A. W. Casey, and R. 0. Lindahl

The Fabrication of Alumina Tooth Roots by Diamond

Contour Grinding 247

G. L. Messing, L. G. McCoy, and C. R. Hassler

Simultaneous Finishing of Both Surfaces

of Hemispherical Domes 255

R. Gentilman, F. Cooke, and P. Miles

Abrasionless Machining Methods for Ceramics 261

R. F. Firestone

Shaping Silicon Compound Ceramics with a Continuous Wave Carbon Dioxide Laser 283

S. M. Copley, M. Bass, and R. G. Wallace

Optical and Surface Physical Characteristics

of Diamond-Machined Infrared Windows 293

D. L. Decker, D. J. Grandjean, and J. M. Bennett

Electrochemical Grinding and Surface Damage for

Cemented Carbides , , 305

S. Malkin and R. Levinger

A New Ultra-Fine Polishing Technique for Sapphire 317

T. Sugita, S. Kawakami , and 0. Imanaka

Fine Polishing of Sapphire Crystals 325

M. Ikeda, A. Yamada, Y. Kokaj i , S .. Kimura

x
Session III: Characterization of Machined Surfaces

Surface Characterization of Optical Materials 337

H. E. Bennett

The Nature of Strength Controlling Machining

Flaws in Ceramics 351

R. W. Rice and J. J. Mecholsky, Jr.

Acoustic Surface Wave Measurements of Surface

Cracks in Ceramics 379

B. T. Khuri-Yakub, A. G. Evans, and G. S. Kino

Characterization of Surface Damage in Ceramics

Using Selected Area Electron Channeling 395

J. Lankford and D. L. Davidson

Y-Deflection Modulated Secondary Electron Images

in the Evaluation of Ceramic Surface Finish 407

L. P. Cook, E. N. Farabaugh, and C. D. Olson

Latent Scratches on Optical Glasses that

Appear After Cleaning 417

T. Izumitani, E. Miyade, S. Adachi , and S. Harada

Session IV: Effects of Machining and Surface Finishing

on Ceramic Properties

Machining Flaws and the Strength Grain Size

Behavior of Ceramics 429

R. W. Rice

Surface Machining Effects on the Strength

Distributions of Aluminas 455

T. E. Easier, T. A. Countermine, R. E. Tressler,

and R. C. Bradt

Effect of Surface Finish on the Strength of Hot

Pressed Silicon Nitride 463

C. A. Andersson and R. J. Bratton

The Effect of Surface Finishing on the Strength of

Commercial Hot Pressed Silicon Nitride . . 477

C. Cm. Wu and K. R. McKinney

Surface Condition and Strength of Silicon CarbideSilicon Materials 483

T. J. Whalen and L. L. Terner

The Effect of Grinding and Polishing on the Strength

of Annealed Soda-Lime Silica Glass 489

R. Bowles

Surface Flaw Effects on Crack Propagation from

Subsonic Water Drop Impact 501

Y. M. Ito and M. Rosenblatt

xi
Surface Characteristics of Diamond-Machined Ceramic

Surfaces Evaluated by High Fluence Diagnostics . . M. J. Soileau, J. 0. Porteus, and D. L. Decker

513

Symposium Attendees 519

Author Index 529

Subject Index 530

xii
Session I

Mechanisms of Abrasive Machin

and Surface Finishing

Session Chairmen

A. Diness, Office of Naval Research

B. Wilcox, National Science Foundation

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

ABRASIVE WEAR IN CERAMICS:

AN ASSESSMENT

A. G. Evans

University of California

Berkeley, CA 94720

The mechanisms of material removal during abrasive wear have been examined,

and preliminary analyses presented. Primary emphasis has been placed on the

lateral fracture mechanism, in an attempt to elucidate both its realm of importance and the concomitant material removal rates. Thermal stresses induced by

plastic penetration have also been shown to be a potential source of material

removal, especially at low levels of the normal force.

Key Words: Abrasive wear; ceramics; fracture, plastic penetration; thermal stresses

1. Introduction

The abrasive wear of ceramics has been the subject of extensive empirical investigation [I].

1 For example, material removal rates have been measured on various materials

as functions of the normal force, P

n

, the horizontal (frictional) force, P , the wheel

speed, etc. However, the detailed mechanisms of material removal have only been cursorily

explored. Yet, some understanding of mechanisms is an essential prerequisite to optimization (especially for complex processes such as abrasive wear). The intent of the present

paper is to examine possible mechanisms of abrasive material removal, as a prelude to

suggesting research studies that should elucidate details of the various removal processes.

Several investigators [2-4] have recently recognized the important role of lateral

fracture in abrasive material removal. This mechanism of wear is explored in some detail,

to determine both its realm of importance and the parameters that might influence the removal rate. Other potential mechanisms of material removal are then examined to assess

their significance as alternative or superposed influences.

2. Lateral Fracture Mechanisms

2.1 Regime of influence

Studies of indentation conducted under quasi-static [4], sliding [2,3] and impact

conditions have indicated a threshold for lateral crack formation. A definition of this

threshold is of great significance, because it defines the lower limit of pertinence of

the lateral fracture mechanism of abrasive wear. A transition to another mechanism may

also obtain under very severe grinding conditions [2]. However, this regime could be just

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

1
another manifestation of lateral cracking (perhaps involving a multiplicity of lateral

cracks as observed during projectile impact [5]); consequently, this "transition" is not

afforded specific consideration in the present analysis.

During "steady-state" grinding conditions, a ceramic will contain a distribution of

microcracks, formed as a consequence of the stresses produced by the penetration of the

grinding particles. These pre-existent microcracks are considered to be the sources of

lateral fracture. The pre-existent microcracks will develop into lateral cracks when the

grinding stresses attain the requisite level. The tensile stresses that induce lateral

cracks develop primarily during load removal [6,7] (i.e. residual stresses), and exhibit

a maximum either within the plastic zone (in work hardening materials) or at the elastic/

plastic interface (in non-hardening materials). It has recently been demonstrated [8] that

elastic/plastic indentation fields can be approximated in the immediate vicinity of the

plastic zone by the elastic/plastic solution for a spherical cavity, 2 and remote from the

indentation by the elastic solution for a half-space: the intermediate stresses being obtained by interpolation. This procedure also provides estimates of the residual tensile

p stress, azz

, obtained by subtracting the elastic unloading stresses from the elastic/plastic stresses at peak load. In particular, the residual stresses in the elastic zone (the

zone of primary interest for fracture analysis [4]), are obtained at zero load by matching

the elastic stress at the indenter interface, a!^, to the equivalent elastic/plastic stress

at peak penetration, a, (i.e. to the hardness H), viz;

5

e = /3H\/z_yr a*u ] (1) zz V2 / \AJ/ u

2 + a

2 z2

where

2u = r

2 + z

2 - a

2 + /(r2 + z

2 - a

2

) + 4a 2 z

2

and then subtracting this stress from the elastic/plastic stress, o^;

(4---*X£?r)h(StS)]

(2)

(z > 0)

By noting that, ay

~ H/3 [4,6,10], and that, r

p

~ a [4,5], the residual stress at zero

load becomes;

R

(3)

°zz

If _ 5

e V (—\ I 322

- a [ + 3

/z 2 - r

2 \]j

\

ZZ ° zz/ \2 / |/u"(u 2 + a

2 z2

) " 9(z 2 + r

2 )3/2[ \ z 2 + r 2/]j

2 . With the obvious exception of a stresses near the surface, which are poorly predicted

by the spherical cavity solution.
The peak axial value of this stress,

zz H/2

occurs at the elastic/plastic interface, and is appreciably larger than the peak stress at

-R maximum penetration [8] (~ H/12). Note, however, that a decreases rapidly as the relaR tive plastic zone size increases, e.g.3 a = 0.27H at r = 2a. zz p

SURFACE

PLASTIC i ZONE l °zz

Fig. 1

DISTANCE FROM AXIS (r/a)

The spatial dependence of the residual tensile stress along a plane through

the intersection of the plastic zone with the penetration axis: the

stresses were estimated from an analytic elastic/plastic solution for a spherical cavity and an elastic solution for a half space.

The spacial variation of the stress along a plane through the location of maximum

tension is plotted in Fig. 1. The stresses can be approximated by

zz

R

J zz

= o(l - r/2a) (r < 2a)

(r > 2a)

(5)

where a is the peak stress. A stress-intensity factor solution for an axisymmetric crack

located in such a linearly varying stress field has recently been derived by Lawn and

Evans [11]. This solution demonstrates the existence of an absolute minimum P* in the normal force required for crack activation. For the stress field represented by eq. (5), P*

is given by

= 14^ e K VH3 (6)
where K

£

is the toughness of the material, 3 is a parameter related to the shape of the

indenting particle [11] (e.g. 3 = 2/tt for a Vickers indenter) and g = H/a. In a material

that contains a wide size distribution of pre-existent microcracks, there will be a relatively high probability of lateral crack nucleation at P*. However, more generally, P*

will need to be exceeded: by an amount related to the size range of microcracks in the

vicinity of the indentation. The parameter K

c

4 /H 3 in the expression for P* is considered

to afford a reasonably accurate relative measure of the influence of the material variables

on the fracture threshold, as indicated by recent observations on radial fracture [12].

However, the magnitudes of the coefficients - being strongly dependent on the value of the

peak stress - could be appreciably in error and should only be regarded as order of magnitude estimates.

2.2 Material removal rates

The amount of material removed by the passage of each abrasive particle that exceeds

the fracture threshold force P* is determined (in part) by the extent of the lateral fracture. Most observations of lateral fracture indicate that the fracture extent is related

to the extent of the prior radial fracture (4,5]; but the generality of this result has

not been substantiated. Nevertheless, the present analysis of material removal rates will assume a relationship between lateral crack extension and the more extensively characterized radial crack extension, C. The average depth of the lateral cracks, < h >, also influences the material removal (Fig. 2a). The only available analysis of lateral crack

depths indicates that < h > is proportional to the plastic zone radius r . This relationship will also be assumed in the present analysis.

The maximum volume of material that could be removed by the passage of an abrasive

particle is (Fig. 2);

V. - 2 < h. > C

i £

i (7)

where z. is the distance of motion. This expression will be the basis for estimating variations in the material removal rate.

The plastic zone radius r

p is related to the plastic contact radius a, under both

dynamic [5] and quasi-static [4] conditions. We will use the direct proportionality,

r « a, so that,

P

< h > a i^a (8)

where i|i is a constant * 1. However, the radial crack extension C differs appreciably

under quasi-static and dynamic conditions [5]: being more extensive (for the equivalent

contact area) under dynamic conditions. For typical abrasive wear situations, both quasistatic and dynamic conditions can obtain. A preliminary analysis that embraces both

quasi-static and dynamic influences is attempted herein. The analysis considers the exten3 A tenuous rationale for this observation is that the presence of the radial cracks en- hances the stress intensity factor at lateral cracks toward the limit of their extension,

by relaxing the circumferential constraint and permitting additional z-displacements.

4
sion of the lateral cracks normal to the motion of the abrading particle (Fig. 2a), since

this distance defines the width of the potential material removal zone. The situation to

be modelled is defined in Fig. 2b; it consists of an abrasive particle subjected to normal

force P , moving with a prescribed lateral velocity v p

. n

* If the velocity v is very low, the consequences of the lateral motion can be approximated by a quasi-static lateral force P = yP (where p is a friction coefficient). Although this condition is rarely encountered, it is instructive to examine the influence of

the lateral force, vis-a-vis the normal force. The normal force dictates the penetration

according to the normal hardness relation [11];

H = P

n

/(3Tra2 (9)

(b)

Fig. 2 a) A schematic indicating the nature of the lateral cracking that leads to

material removal

b) A schematic indicating typical conditions that apply to an abrading

particle during abrasive wear.

Both the lateral and normal forces are expected to influence the stress field. By analogy

with the equivalent elastic problem [13], it is anticipated that the lateral force will

tend to rotate the stress field in the direction of P^, in the sense that the tension is

enhanced behind the contact zone and suppressed ahead of the contact. But, there should

be no appreciable influence on the stresses contained in the plane normal to P . The latter have the primary effect on the lateral crack extension normal to the motion of the

abrading particle. It is thus concluded that under quasi-static conditions, the width of

the lateral fracture zone should be relatively insensitive to the magnitude of the lateral

5
force. The zone width may, therefore, be approximately characterized by a relation similar

to that obtained for radial cracks under conditions of normal penetration; viz., for a Vicker's indenter [14],

C/a = F

1 [(Kc /H/a), (H/E)] (10a)

where is a function, which for c/a > 2 reduces to;

>29(-M

\H/a/

2/3

F

1 « 0.29[ -^J (10b)

giving;

C

= 0.29 /_^\

2/3

00c)

(ua)

2 /3 [

KJ

Combining equations (7), (8) and (10c) yields the quasi-static material removal rate;

[ 0.58i|<

] VjTA i LU0)

7 /

6 J(kc

2 /3 hV2

)

1

More generally, the lateral motion of the abrading particle will lead to dynamic effects. These dynamic effects arise both from the prescribed motion v^ of the particle, and

from local instabilities. The instabilities occur following a fracturing event and result

in an abrupt increase in the local particle velocity. They are manifested as the lateral

force oscillations typically observed in studies of the motion of abrading particles. The

impact events that result from the instabilities are likely to have impulse components

both normal and parallel to the surface. The normal component will affect the penetration,

but both components will influence the dynamic stress field. We will assume that, as in

the quasi-static case, the lateral component of velocity does not generate stresses that

significantly influence lateral crack extension in the direction of present interest (Fig.

2a). The effect of the normal particle velocity v can then be assessed from experimental

results obtained for projectile impact [5];

C * X [VV

n 2/K

c ] V3

0*)

where r

p is the radius of the projectile and A is a constant. The normal component of the

velocity acquired by the abrading particle depends on the magnitude of the normal force P

n and the height z that the particle rises above the surface during the instability (Fig. 2b);

\

2 = 2P

n z/m (13)

6
Where m is the mass of the grinding wheel. The height z will be influenced by the detailed

character of the instability, the imposed lateral velocity v^ and the friction coefficient

y. The influence of normal velocity on the extent of the fracture will be assumed in the

present analysis to be equivalent to that found for projectile impact. This assumption is

based on the premise that the stresses which influence the fracture relate to the condi5 tions that prevail soon after initial contact: as dictated by the impedence, the initial

velocity mismatch and the dynamic hardness. Clearly, experimental studies of the velocity

dependence of C are needed to assess the utility of this assumption. The crack extension

obtained from eq. (12) is then;

C

=

p

n z (v ry )

K

2/3

F(r . m) (14)

where F is an undetermined function. It is interesting to note that the dependence on the

normal force is identical to that obtained for quasi-static indentation (eq. 11). Now,

assuming that eq. (9) still affords a reasonable estimate of the contact radius, equations

(7), (9), and (14) can be combined to give;

2 F(r

p ,m) (P ).

y /6

v n i

L H 1/2

2/3

A.- (15)

This result is quite similar to the quasi-static result except that influences of the

lateral velocity and the friction coefficient emerge, through the instability height z.

No information is presently available concerning the magnitude of the instability; this is

clearly an important parameter for future study.

100

E

JL

O

CVJ X

»-Q

o

<

o

50 -

£ 30

20

10

1

1 1 1

£T ^-PREDICTED

1

/ SLOPE, Pn213 -

1 1 1 0.1 0.2 0.3 0.5

NORMAL FORCE, Pn

(N)

1.0

Fig. 3 A plot of the lateral crack

width as a function of the normal

force [3], indicating the P^/ 3 relation.

7
Fig. 4 A plot of the material removed

as a function of the applied force [15].

50 100 200 300

DEPTH REMOVED (/tin)

Fig. 5 A plot of the inverse of the

material removal against the material

dependent parameter K

2/

3

H

1 ^ 2 predicted by the lateral fracture analysis.

8
A comparison of the predictions of the present simplified analysis with available

results should provide a basis for assessing the merits of this approach to abrasive material removal. The dependence of the lateral crack width on the normal force has previously been examined by Veldkamp et al [3] in MgZn ferrite, and shown to approximately

2/3

conform with the predicted P

n

' relation (Fig. 3). The dependence of the material removal

on the normal force has been explored by Hartley and Wilshaw [15] (Fig. 4). A non-linear

force dependence with an exponent ranging between 1.07 and 1.3 was found, similar to the

7/6 predicted dependence of P ' . The important dependence on the material properties K

c and

H can be ascertained from the study of Rice and Speronello [16]. A comparison of their

results with the predicted material dependent quantity, H

1 /

2 K

c

2//3

, reveals a consistent

trend (Fig. 5). However, variations of the instability height with the material type have

not been incorporated in the comparison, and these could account for some of the residual

trends in the experimental results.

The reasonably good correlation of the available abrasive wear results with the lateral

cracking model encourages further study. These studies would include the characterization

of lateral cracking under conditions that simulate abrasive wear, and estimations of the

zone of material removal vis-a-vis the zone of lateral fracture.

3. Other Mechanisms

3.1 Thermal stresses

It is widely recognized that heat is generated during a grinding operation. This heat

derives primarily from the plastic work expended by the penetration of the abrading particles. An upper bound (adiabatic) estimate of the thermal stresses that develop as a consequence of the heat generation is derived in this section. The potential influence of the

thermal stresses upon material removal are then assessed.

The plastic work W

p expended during indentation can be estimated from the macroscopic

force-penetration curve, (Fig. 6) as;

P

W

p

= /Pd6

p (16)

o

where 6 p is the plastic penetration. The plastic penetration is directly related to the

contact diameter 2a through the profile of the penetrator. For a penetrator of fixed profile (Fig. 6);

<5 p

= a cote (17)

The penetration pressure (i.e. the hardness, H) for a fixed profile indenter is independent

of the penetration. Hence, for an axisymmetric penetrator, combining equations (9), (16)

and (17) and integrating gives;

It is interesting to note the

m

inverse dependence on the hardness. This arises because the

plastic penetration decreases more rapidly with the increase in hardness than the

9
SURFACE PENETRATOR

plastic y ZONE '

Fig. 6 A schematic indicating the elastic

&i and plastic 6 penetrations by a fixed

profile indenter.

PLASTIC

WORK = /PdSp

concomitant increase in the penetration force.

The work expended during a lateral traverse of an abrading article depends very sen- sitively on the details of the traverse. In the presence of instabilities in the lateral

motion, individual impacts will occur at various sites along the traverse. The plastic

work at each impact Site will be given approximately by eq. (18). Under adiabatic

conditions, if all of the plastic work is converted into heat, the temperature T generated

within the plastic zone at each impact site is;

[(wskpVp]" 1 (19)

where c is the specific heat and p is the density. Setting r = a and combining with eq.

r r (9) then gives;

( 3 3 cote t ( H \

V

4

) VC J

(20)

Note that the temperature rise is independent of the applied force. Heat flow during penetration will clearly reduce this temperature, as dictated by the differential equation;

dfT

dr2 (21)

where k is the thermal conductivity.

10
The local increase in temperature will generate thermal stresses. By assuming the

plastic zone to be a spherical enclave in an infinite body (Fig. 7), the thermal stresses

are analogous to the stresses that develop due to thermal expansion mismatch [17], viz,

3 2EaT

rr

3(1 - v)

3(1

(?)

/ V 3 (r > r

p

) (22)

a rr

~ °ee 2EaT/3(l - v) (r < r

p

)

where r is the distance from the center of the plastic zone. An upper bound value for the

tensile stress generated by the local plasticity, obtained from eqs. (20) and (22) is thus;

e cote

4(1 - ; 0

Fig. 7 A plastic enclave in an elestic

matrix indicating the stresses developed

by heat generation in the plastic zone.

PLASTIC

ENCLAVE

/ ctHE \

V

pC

p /

(23)

Inserting some typical values for E, p, c and a into eq. (23) yields values of a na in the _2-i P ee range ~ 10 to 10 H. These stresses approach the stresses that result from plastic

penetration (eq. 4). Thermal stresses should thus be regarded as a possible source of material removal, particularly in materials with large values of thermal expansion coefficients, elastic modulus and hardness, and low values of specific heat. Since the thermal

^An elastic solution is not strictly valid, because the thermal shear stresses, which are

a maximum at r

p

, superimpose on the penetration stresses and extend the plastic zone.

11
stresses are independent of the normal force, the thermal influences should be particularly

effective at low force levels. The thermal stresses could either enhance the lateral

fracture just above the threshold force P*, or perhaps, provide an independent mechanism

of material removal below P*.

3.2 Elastic stress pulses

Impact events can, in certain situations, lead to the generation of Rayleigh waves

[18], The Rayleigh wave can produce large amplitude, short duration tensile pulses,

which introduce circumferential cracks that link together to cause material removal [19].

A similar mechanism could apply in abrasive wear, under analogous circumstances, e.g., the wear of very hard materials at low force levels. An important feature of this mechanism

is that crack development is impeded by large values of the elastic wave velocity.

4. Summary

Some mechanisms of material removal during abrasive wear have been assessed. A preliminary lateral fracture model has been developed. The model indicates that a minimum

force is required to initiate lateral fracture. The threshold force is determined primarily by the fracture toughness and the hardness.

The material removal that occurs by lateral fracture above the threshold has been

estimated to depend on the normal force, the hardness and toughness as well as on the

the lateral velocity and friction coefficient (through local instabilities caused by the

fracture process). A comparison of the predictions that emerge from the preliminary model

with available data indicates encouraging correlations. Further development of lateral

fracture models thus appear to be a promising course for future studies of abrasive wear.

An upper bound estimate of the thermal stresses that result from plastic grooving

indicates that these stresses can be appreciable. The stresses are, to a first approximation, independent of the normal force, suggesting that the thermal stresses are likely

to be most important at low force levels, i.e., when the lateral fracture mechanism is

least effective. The thermal stresses are enhanced by large values of the thermal expansion coefficient, elastic modulus and hardness and small values of the specific heat.

Thermal effects are most likely to emerge in materials with these properties. Thermal stresses can either enhance the extent of lateral fracture just above the threshold, or provide

an alternate material removal mechanism.

Finally, one other mechanism has been suggested, by analogy with impact problems. It

entails the growth of cracks by elastic tensile stress pulses generated by deformable particles. However, no specific estimates of the realm of importance of this mechanism have

yet been obtained.

ACKNOWLEDGEMENT

This work was supported in part by the Office of Naval Research and in part by the

Division of Materials Sciences, Office of Basic Energy Sciences, U.S. Department of Energy.

12
REFERENCES

I] Gielisse, P. J. and Stanisloa, J., The science of ceramic machining and surface

finishing, NBS Special Technical Publication, [348], 5, (1972).

2] Swain, M. V., Fracture mechanics of ceramics (Ed. Bradt, R. C, Hasselman, D.P.H.

and Lange, F. F.), Plenum, NY, 3, 257, (1978).

3] Veldkamp, J.D.B., Hattu,N. and Snyders, V.A.C., Fracture mechanics of ceramics,

3, 273, (1978).

4] Evans, A. G. and Wilshaw, T. R. , Acta Met., [24], 939, (1976).

5] Evans, A. G. , Gulden, M. E. and Rosenblatt, M. , Proc. Royal Soc, [A361], 343, (1978)

6] Lawn, B. R. and Wilshaw, T. R., Journal of Mater. Sci., [10], 1049, (1975).

7] Swain, M. V., and Hagan, J. T., Jnl . Phys., [9], D, 2201, (1976).

8] Evans, A. G. , ASTM Special Technical Publication, in press

9] Perrott, CM. , Wear, [45], 293, (1977).

10] Tabor, D. , Hardness of Metals, Oxford: Clarendon Press, (1951).

II] Lawn, B. R. and Evans, A. G., Jnl. Mater. Sci. [12], 2195, (1977).

12] Lankford, J., to be published.

13] Hamilton, G. M. , and Goodman, L. E., Jnl. Appl . Mech., [33], 371, (1966).

14] Evans, A. G. and Charles, E. A., Jnl. Amer. Ceram. Soc, [59], 371 , (1976).

15] Wilshaw, T. R. and Hartley, N.E.W., (Proc. Third European Symposium on Comminution),

Chemia GMBH, Weinheim/Bergestrasse, (1972).

16] Rice, R. W. and Speronello, B. K. , Jnl. Amer. Ceram. Soc, [59], 330, (1976).

17] Seising, J., Jnl. Amer. Ceram. Soc, [44], 419, (1961).

18] Evans, A. G. , Treatise on materials science and engineering, (Ed. Preece, CM.),

Academic Press, in press.

19] Adler, W. F. , Jnl. Mater. Sci., [12], 1253, (1977).

Discussion

CHOUDRY

1) The superposition of the compressive stress (which occurs during loading) and the tensile stress (which occurs during unloading) could be questioned since these two are operational at different times.

2) The calculation of temperatures through the integration of the 'Adiabat

again to be looked from the time point of view, i.e. during compression

o & e (strain) increase giving the adiabat; which you integrated; however during the tensile relaxation (which is a part of your treatment)

13
the adiabat would reverse itself (in the absence of internal friction etc.)

completely and the heat dissipation (area between the two curves) would be

Zero. In any case i.e. even in the presence of internal friction etc, the

adiabat integral of your paper would grossly overestimate the temp.

EVANS

1) In elastic/plastic problems, the residual stresses are obtained by

subtracting the elastic stresses associated with unloading from the elastic/

plastic stress at peak load (See Hill's book on plasticity). This is standard procedure.

2) The thermal stresses calculated in this paper refer to the heat generation in the plastic zone during penetration. This heat results directly

from the plastic work (the irreversible component of the penetration) expended during penetration. The unloading, as noted above, is elastic and

therefore does not contribute to the heat generation, not can it absorb

the heat generated during penetration. The heat generation does influence

the local plastic flow properties, through the temperature dependence of

the contributive equations. The appropriate hardness is thus the 'adiabatic*

hardness which may be lower than the isothermal hardness.

M. SRINIVASAN

1. What are the mechanisms by which the lateral cracks curve toward the

surface after propagating a distance c?

2„ In such a situation will Mode I critical stress intensity factor alone be

sufficient?

EVANS

iy

detion

The residual stress calculations indicate principal tensile stress

trajectories that initiate within the plastic zone and progress gradually

toward the surface. The conventional wisdom would then anticipate crack

velopment along these trajectories. However, some crack surface interac

should be expected in this case, causing the cracks to approach the surface

before the principal tensile stress trajectory reaches the surface. Crack

propagation is exclusively Mode I.

S. W. FREIMAN

Based on your analysis of the formation of lateral cracks during machining could you also not calculate the sizes of the radial cracks formed?

Shouldn't you then be able to calculate the strength of the workpiece after

machining?

EVANS

The radial crack sizes can indeed be characterized according to a similar

scheme. The prediction of the strength after machining would also require

that the residual stress distribution be known. This could be estimated

using the methods described in the paper and then used to predict the strength.

14
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice/ editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

NUCLEAT I ON OF MEDIAN AND LATERAL CRACKS AROUND

VICKERS INDENTATIONS IN SODA-LIME GLASS

J.T. Hagan, M.V. Swain and J.E. Field

Physics and Chemistry of Solids

Cavendish Laboratory

Cambridge /England

The median, radial and lateral cracks that form around Vickers indentations in

soda-lime glass have been studied. Cross sections of the plastic indentations have

revealed a system of intersecting flow lines in the deformed zone directly beneath

the indenter; these flow lines lie close to the elastic shear trajectories beneath

a pointed indenter and also around a radially expanding cylinder. It is suggested

that the nucleation of both the median (radial) and lateral cracks arise from the

inhomogeneous nature of the subsurface deformation. The nucleation of the bulk

median crack occurs at the intersection point of two limiting flow lines and is

therefore similar to crack nucleation from a pile of dislocations at an obstacle,

or the interaction of dislocations on two intersecting slip planes. The radial

cracks are nucleated from the interaction of similar flow lines that develop on

the sample surface and around the indentation. The lateral cracks appear to form

as shear cracks along the flow lines and are propagated by the unloading residual

stresses. The interaction of the flow lines may also be responsible for the

powdering or crushing of material that occurs at high indenter loads.

Key Words: Fracture; plastic indentations; crack nucleation; median, radial and

lateral cracks; shear flow lines; residual stresses.

1 . Introduction

When blunt or sharp indenters are loaded slowly or impacted on brittle solids, the

extensive cracking - media n, (radial) and lateral cracks - that forms around the contact can lead to erosion, wear and strength degradation. These cracks are also basic to such processes like abrasion, machining and drilling since all these processes may be regarded as

the summation of damage at all contact sites. The relevance of these cracks to a better

understanding of the above processes demands a proper study of the origins of the various

cracks that develop around blunt (elastic) and sharp (plastic) indentations.

The case of the radial, median and lateral cracks around blunt contacts during impact

has been treated by Chaudhri and Walley [1] and it will not be considered here; they,

however, suggest that the radial and lateral cracks, that develop during elastic impact,

arise from purely frictional effects between the newly created fracture surfaces as they

are brought together during overloading.

Of more interest, however, are the median, radial and lateral cracks that form around

plastic indentations; this is because they form at loads as low as 0.5N. This paper is,

therefore, concerned primarily with the origins of the median, radial and lateral cracks

that develop around plastic pyramidal indentations.

These cracks are illustrated in figure 1. The median (radial) cracks, pp, form across

the major diagonals of the plastic indentation; the lateral cracks are subsurface cracks

that give rise to the halo, h, around the plastic indentations. Within the plastic zone itself and extending from the edge of the indentation is an annular ring of cracks, cz; these are clearer in figure lb which shows a Vickers indentation in toughened soda-lime

glass. The cross-sectional views are illustrated schematically in figure 2. The median

crack is a semi-^circular crack, npn, with the deformed zone (the hatched region) at its

origin; pp is the section of the other orthogonal crack. The radial cracks, rc, are

15
Figure 1. Vickers indentations in annealed and toughened soda-lime glasses respectively.

shallow surface cracks which emanate from the edge of the deformed zone: the different

situations which give rise to radial rather than median cracks are discussed later . The

lateral cracks, lc, develop from the bottom of the deformed zone and extend in saucerlike fashion to the specimen surface [fig. 2] . (See Lawn and Swain [2] , Swain and

Hagan [3] and Hagan and Swain [4] for a full description of these cracks).

Figure 2. Schematic of median, lateral and radial cracks, mc, lc, rc, respectively around

sectioned Vickers indentations. The oval-hatched zone marks the extent of the deformed

The propagation of these cracks, especially of the median cracks, is fairly well understood and there are standard fracture mechanics expressions to describe the propagation

(Lawn and Fuller [5] , Evans and Charles [6] ). The more important problem of the initiation of these cracks has not been treated. The initiation process is often complicated

by the fact that the material may be crystalline or amorphous and material properties like

grain size, d, hardness, H, and the critical stress intensity factor Kic would be expected

to influence the nucleation process. Lawn and Evans [7] have recently suggested a model

for the nucleation of the median crack in glasses. Their model relies on the existence

of subsurface fortuitous flaws of the right critical dimensions in the vicinity of the

elastic-plastic boundary where the tensile stresses are highest. In this paper we present

a new model for the origins of the median and lateral cracks around plastic indendations

in soda-lime glass. Details of this model have been published by Hagan and Swain [4] .

Since the median and lateral cracks are associated with the plastic deformation under

the indentation, the study of the origins of these cracks requires examination of cross- sectional views of the plastic indentations. The indentations were sectioned by making the

indentation on and near the tip of a pre-existing crack. Similar sectioning technique has

been employed by Peter C8 ] and Mulhearn [9] .

a

zone

2. Experiments

16
3. Results

A summary of the ooservations are illustrated with figures 3 and 4; figures 3 and 4

show the surface and sub-surface damage around and under Vickers indentations in soda-lime

glass, at indenter loads of 2N and 50N respectively. Figures 3a, and b illustrate surface

traces of the radial or median cracks around the plastic indentation on the specimen

surface. The halo, h, around the plastic indentation in figure 3b is reflected light from

the sub-surface lateral cracks. The sub-surface view consists of a deformed zone, dz, directly beneath the indenter, and the lateral cracks, lc, which emanate from the bottom of

the deformed zone. At such low indenter loads, there is no evidence of the trace of the

median crack in the bulk. The traces on the specimen surface are therefore shallow radial

cracks confined to narrow regions of the specimen surface and are propagated by unloading

residual stresses. In most cases the radial cracks could be watched propagating after the

load has been taken off completely.

At the higher indenter load of 50N (fig. 4) the surface damage consists of traces of

the characteristic radial/median and lateral cracks. There are also some shallow cracks,

cz, in figure 4a, that occur within the surface of the plastic indentation itself. Similar

observations have been reported by Mliller[ 1 0

]f or Vickers indentations in phenolic resin.

The sub-surface deformation consists of a deformed zone and traces of both the median and

lateral cracks, mc and lc respectively, as in figure 4b. The deformed zone is made up of

a series of intersecting inhomogeneous flow lines which are characteristic of materials

undergoing radial flow (Peter, [8] and Nadai [11] ). These flow lines are clearer in

Figure 3 (a), (b) show 2N Vickers indentations on a pre-existing crack, ss. The subsurface views (c) and (d) show no evidence of median cracks and are dominated by the

lateral cracks.

17
Figure 4 (a), (b) and (c) are surface and sub-surface views of a 50N Vickers indentation.

Sub-surface deformed zone, &z_, consists of network of flow lines aa,bb and cc as in (c)

;

(d) , (e) and (f) show sub-surface damage of another 50N indentation.

figures 4d-f which are sub-surface deformation of another 50N load indentation. Some of the

flow lines, arrowed in figure 4d, appear to degenerate into lateral cracks outside the

plastic contact; they also coincide with the short shallow cracks, ss, within the surface

of the plastic indentation. The dominant lateral cracks, lc, again emanate from the bottom

of the deformed zone (figure 4b). The trace of the median cracks appears at the bottom of

the deformed zone at t and n the intersection points of two of the flow lines in figure 4.

At higher indenter loads, extensive crushing of the materialwithin the deformed zone occurs and the flow lines become unresolvable. Such entanglement of shear flow lines and

the subsequent break up of material within the deformed zone may constitute the early

stages of the crushing of the material within the contact zone.

18
4. Discussion

From the sub-surface deformation, it is obvious that the radial, median and lateral

cracks are nucleated as a result of the inhomogeneous nature of the deformation beneath the

indenter

.

It appears that the lateral cracks are nucleated as a result of the special nature of

shear deformation in glassy materials; shear deformation in glassy materials in the absence

of any long range order is non-constructive, and the overstraining of atomic bonds leads to

bond breakage and the degeneration into shear cracks. These cracks, however, do not propagate under load because of the compressive stresses that exists across the prospective

crack path. The driving force for the propagation of these lateral cracks is provided by

the unloading residual stresses that develop around the indentation as the load is taken

off.

The nucleation of the median cracks at the points of intersection or blocking points

on a flow line is similar to crack nucleation from the interaction of dislocations on two

intersecting flow lines or the pile up of edge dislocations against an obstacle (see Zener

[12], Stroh [13]). The nucleation and propagation of the median crack is therefore by the

wedging action of the deformed zone beneath the indenter and not by the indenter itself

as is commonly believed and on which most analyses are based; this is, however, only true

for wide angle (blunt) wedge for which the deformation is radially from the indenter faces.

The shallow radial cracks confined to the specimen surface are probably nucleated

by the interaction of similarly occurring flow lines that develop on the surface of the

specimen (Peter [8] ). These radial cracks may also form to accommodate displacement

of material in two orthogonal directions from the indenter faces.

The shallow surface cracks, cz (figure 3a) within the plastic indentation may be

traces of the sub-surface flow lines on the surface. They are more probably formed by

the large component of elastic stresses which are associated with wide angle indentations

in highly elastic solids (Hirst and Howes, [15] ). These cracks are therefore similar

to ring cracks associated with Hertzian elastic contacts except that they do reflect the

symmetry of the indenter (Hagan and Swain [4] , Hagan[16] ).

5. Conclusions

It has been established that the cracks that develop around plastic indentations in

solids arise from the inhomogeneous nature of the sub-surface deformation. This has been

confirmed by recent studies of sub-surface deformation of indentations in LiF and NaCl.

It is also obvious that the median and radial cracks would be mainly responsible for the

strength degradation of materials. The propagation of the lateral cracks from the bottom

of the deformed zone in a saucer-like manner to the specimen will remove material with it

as it propagates; these lateral cracks will therefore enhance the erosion of the brittle

materials. For efficient grinding and polishing, however, the dominant median crack

should be avoided by using low polishing loads but high enough for sub-surface lateral

cracks to extend to the surface. The interaction of the shallow radial cracks with

the sub-surface lateral cracks would be responsible for material removal without introducing

the strength degrading median cracks.

The authors are grateful to Dr. M.M. Chaudhri for useful discussions and comments on

the paper. This work was supported by the Ministry of Defence (Procurement Executive)

and the Science Research Council.

19
References

[I] Chaudhri, M.M. and Walley, S.M. , Damage to glass surfaces by the impact of small glass

and steel spheres, Phil. Mag. Vol. 37, 2, 153-65 (1978)

[2] Lawn, B.R. and Swain, M. V. Microfracture beneath point indentations in brittle solids,

J. Mater. Sci. , JJO , 113-22, (1978)

[3] Swain, M.V. and Hagan, J.T., Indentation plasticity and ensuing fracture of glass,

J. Phys. D: Appl. Phys. 9_, 2201-14, (1976)

[4] Hagan, J.T. and Swain, M.V., The origin of median and lateral cracks around plastic

indents in brittle materials, J. Phys. D: Appl. Phys. JJ_, 2091-202, (1978)

[5] Lawn, B.R. and Fuller, E.R., Equilibrium penny-like cracks in indentations, J. Mater.

Sci. 10 [1], 113-22 (1975)

[6] Evans, A.G. and Charles, E.A., Fracture toughness determination by indentation,

J. Am. Ceram. Soc. 59 C 7-8] 371-2, (1976)

[7] Lawn, B.R. and Evans, A.G., A model for crack initiation in elastic/plastic

indentation fields, J. Mater. Sci., \2_, 2195-9, (1977)

[8] Peter, K.W., Densification and flow properties of glass in indentation experiments,

J. Non. Cryst. Solids, 5, 103-15 (1970)

[9] Mulhearn, T.O. , The deformation of metal by Vickers-type pyramidal indenters,

J. Mech. and Phys. of Solids, 7_, 85-96, (1959)

[10] Muller, K. , Determination of Vickers hardness under load - a method considering

viscoelastic behaviour of plastics. The science of hardness testing and its

research applications, J.H. Wesbrook and H. Conrad eds, pp. 291-9, (ASM, Metals

Park, Ohio, 19 73)

[II] Nadai, A., Plasticity (McGraw Hill - N.Y. , 1931)

[12] Zener, C, Fracturing of metals pp. 3-31 (Cleveland, Ohio; Am. Soc. Metal 1948)

[13] Stroh, A.N., A theory of the fracture of metals, Advances in Physics, _6, 418-65

( 1957)

[14] Hirst, W. and Howse, M.G.J.W., The indentation of materials by wedges, Proc. Roy.

Soc. A311, 429-44, (1969)

[15] Hagan, J.T., unpublished work, (1978)

Discussion

GIELISSE

Would, under continued loading, the lateral cracks continue to go downwards, as initially divested or come up to the surface? Have you followed

crack extension under continued loading experimentally in time?

HAGAN

We have not followed the lateral crack extension during continual loading; this has, however, been done for the median crack though. The main

20
driving force for the propagation of the lateral cracks is provided by the

unloading residual stresses. Continued loading will, however, provide an

overloaded system and it is possible for the flow lines to break through the

specimen surface outside the contact area. Note that the flow lines are not

just directed downwards, but they follow logarithmic spiral trajectories and

there is a tendency for them to turn towards the surface.

HOCKEY

Did you notice the formation of cone cracks beneath sharp indenters in

glass ?

HAGAN

No, we did not observe any cone cracks around Vickers indentations in

soda -lime glass,, However, the shallow cracks, CA , in Fig. 16 probably arise from the large elastic radial tensile stresses which develop on the

surface of the specimen for Vickers indentations in highly elastic solids.

These cracks are, therefore, similar to the collar on the 'skirt' of cone

cracks. In fact, for Vickers indentations in fused silica glass, the

dominant cracks are the cone cracks.

21

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

FRAGMENTATION AND DAMAGE PENETRATION DURING ABRASIVE MACHINING OF CERAMICS

H. P. Kirchner, R. M. Gruver, and D. M. Richard

Ceramic Finishing Company

State College, PA 16801

Mechanisms of fragmentation and damage penetration during abrasive

machining of ceramics were investigated using indentation fracture mechanics. Specimens of glass and hot pressed silicon nitride were damaged by

controlled scratching and single point diamond grinding. Vertical and

horizontal forces acting on the material were measured for various diamond

shapes, set depths and wheel speeds. The damage was characterized by optical

and scanning electron microscopy. Cracks, similar to those formed by static

contact and particle impact, were observed. The vertical forces were correlated with the extent of damage yielding results consistent with those

expected based on indentation fracture mechanics.

Initially, the penetration of damage was analyzed in terms of point

contact loading. This analysis resulted in improved understanding of the

effect of diamond shape on the variations of the forces and damage penetration which were manifested by variations in the dimensionless contact

constant x- Subsequently, the results were analyzed in terms of line

contact loading which allowed the results for hot pressed silicon nitride

and various diamond shapes to be represented by a single value of xKey Words: Abrasive machining; grinding; ceramics; glass; silicon nitride;

impact; contact stress; stress intensity factor; fracture mechanics; fragmentation; damage penetration; strength.

1 . Introduction

Empirically developed processes are used for abrasive machining 1 of ceramics despite

several efforts to establish a more scientific basis for these processes [1-7]. It is

highly desirable to develop methods to control fragmentation, to minimize penetration of

subsurface damage and to increase the efficiency of material removal. To aid in accom- plishing these objectives, it is necessary to have improved understanding of the

fracture mechanisms. Based on extensive studies of mechanisms of localized impact damage

[8-20], there has been a growing conviction that the methods of indentation fracture

mechanics can be applied to aid in understanding of ceramic machining. This approach

is the basis of the present investigation.

The various types of damage occurring during localized impact are sketched in

figure 1. These include Hertzian cone cracks, lateral vent cracks, and radial cracks

[8,14]. In addition, crushing is frequently observed. Hertzian cone cracks are formed

by static contacts or low velocity impacts with blunt objects such as spheres or punches.

Impact velocities during abrasive machining are in this low velocity range (10-100 ms"l).

Even though there are many similarities between the damage resulting from "static" contact loading and that from impact loading, there are also significant differences. One

such difference is that, in impact loading, there is less subcritical crack growth so that

fractures originate at higher stresses. In the case of Hertzian cone cracks, the ring

Abrasive machining includes grit blasting, free abrasive machining and other processes,

as well as grinding.

23
Ring Crack

Contact Zone

Cone

Hertzian Cone Crack Lateral Vent Crack Radial Cracks

Fig. 1. Types of cracks formed by localized impact.

cracks formed at the surface may branch at regular intervals which are determined by the

fracture stress and the crack branching stress intensity factor of the target material

[16]. The wedge-shaped fragments formed by this crack branching mechanism may subsequently be subject to material removal.

Impacts by either sharp objects or blunt objects at relatively high velocities cause

radial and lateral vent crack formation. Radial cracks may be perpendicular to the

surface, radiating from the impact site as shown in the top right hand view in figure 1,

but also may occur at other angles to the surface. Radial cracks arise because, as indentation increases, the tensile stresses tangential to the contact area increase and

the radial stresses decrease, causing a shift from ring cracks to radial crack formation.

Lateral vent cracks are saucer shaped cracks that form mainly parallel to the surface

but that tend to come to the surface at the periphery. These cracks form, as the load is

released, by reaction of the elastically deformed material against the permanently

deformed material in the contact zone. Because these cracks form after the radial cracks,

they are frequently separated into segments extending between the radial cracks. When the

cracks extend to the surface at the periphery, these segments are frequently removed as chips or fragments. Because of this tendency toward material removal by lateral vent crack

formation, Lawn and Wilshaw [8] and others have suggested that this mechanism is responsible for material removal during abrasive machining processes such as grit blasting or

grinding.

The depth of penetration of surface damage and its subsequent subcritical crack

growth are the principal factors controlling the strength and integrity of ceramics

intended for use in structural applications. This surface damage may be induced when

the part is machined during the original manufacturing process or subsequently during

use. Clearly, for optimum performance, the ceramics should leave the manufacturing process

with minimum surface damage. Penetration of surface damage and the effect of surface

damage on the remaining strength were investigated previously for scratched and indented

specimens [21-24]. Since that time, the new information from impact investigations has

become available making possible new approaches to understanding abrasive machining

processes.

The depth of damage induced by contact stresses can be estimated by the following

equation [25]

24
IC x

c 3/2 (1)

in which Ktq is the critical stress intensity factor for the particular material, P is

the contact force or load, C is the crack depth, and x is a dimensionless contact

"constant" that incorporates the effects of uncertain geometrical factors, interface

friction, etc. x can be determined in special cases from the theory of hardness indentations but is usually determined by experiment. Appropriate subscripts can be used

with x and C to designate the symbols related to various types of cracks such as Hertzian cone cracks, lateral vent cracks and radial cracks. The analysis assumes that

the crack propagation in the far field is independent of the conditions (indenter shape,

surface roughness) in the contact zone. However, experience has shown that many important cases fall in a less severe damage range.

One objective of this research was to investigate the applicability of eq. (1) to

damage caused by abrasive machining. This included characterization of the various

types of damage and determination of appropriate values of x for these types of damage and

abrasivemachining conditions. The effects of variables such as abrasive grain shape,

set depth, and wheel speed (impact velocity) were investigated.

The remaining strengths of ceramics with subsurface damage were estimated using

KK

= Y a

p C

1/2

(2)

in which ap is the fracture stress and Y is a dimensionless constant determined by crack

shape and proximity to the surface. Values of Y have been determined by fracture mechanics

techniques for a wide variety of cases [26,27]. Equations (1) and (2) can be combined and

solved for the remaining strength (a

F

) by eliminating the crack length (C) yielding

= IC /o

°F 1/3 1/3

v Y X P

It has been common to use eq. (3) to describe the effect of surface damage on* the strength

of ceramics without verifying that the cracks do grow to the predicted dimensions. In the

present investigation, the surface damage was introduced by scratching and single point

grinding and care was taken to characterize the cracks to assure that the crack dimensions

were properly accounted for by the above equations.

2. Procedures

2.1. Materials

Several criteria were used to select materials for this investigation. Perhaps the

most important criterion was the ease with which cracks and other damage could be observed.

This criterion was the principal reason for the selection of soda-lime-silica glass 2 and

hot pressed (H.P.) silicon nitride 3

. Other factors considered were the potential importance of the materials as structural ceramics, availability and cost.

2.2. Techniques for material removal

Two techniques were used to remove material from ceramic surfaces; scratching and

single point diamond machining. For scratching tests, a microtome which consists of a

2Float glass, Centre Glass, Inc., State College, PA, thickness 6 mm.

3 NC-132, Norton Company, Worcester, MA, grain size 1-2 pm.

25
massive carriage sliding very smoothly on metal ways was modified to provide a means to

press a loaded diamond point against a ceramic plate specimen mounted on the carriage while

the carriage was pulled past the diamond point at approximately constant speed by means of

a weight acting on a cord running through a pulley. The diamond points, which were 75°

coned wheel dressers were loaded with various weights to form scratches of various depths

[23].

The single point diamond machining was done using an apparatus located at the Department of Chemical Engineering, University of Rhode Island. The equipment consists of a high speed milling machine 4 with pressure cylinders arranged to drive the table under the

tool at a constant speed of 0.017 ms

- ^. A dynamometer, mounted on the table was used to measure the vertical and horizontal forces transmitted from the tool to the specimen.

The forces were converted to electrical signals by Kistler quartz load cells (Model

504A) and displayed on a Tektronix Type 564 storage oscilloscope. In most cases, the

traces were photographed to provide a permanent record of the forces experienced by the

specimen (Fig. 2).

Fig. 2.

H.P. Si

flat,

Load traces for single point grinding of

N., top trace is vertical force (50 ym

2!d urn set depth, wheel speed 8000 rpm).

The forces were calibrated using a pendulum falling from known heights. This apparatus

was used to construct calibration curves giving the actual vertical and horizontal forces

in terms of the nominal forces determined taking into account the transducer range, vertical

scale setting of the oscilloscope, etc. Usually, the vertical forces were larger than the

horizontal forces. The results were analyzed" in terms of the vertical forces because this

is the principal factor controlling penetration of surface damage.

The diamond points used in these experiments were single crystals ground to form 120° cones. Normally, the tip radii were about 40 ym. In some cases flats were ground on the

tips to form truncated cones. These flats were 50 or 100 ym in diameter. The diamond

points were fastened to the periphery of a two and one-half inch diameter wheel which

rotated at speeds ranging from 4,000 to 12,000 rpm.

A typical test was performed as follows:

1. The height of the diamond was adjusted so that it would traverse the specimen at

the desired depth beneath the surface.

2. The wheel speed was adjusted to the desired value using a stroboscope (General

Radio 1531-AB).

3. The horizontal trace on the oscilloscope was adjusted so that the load was recorded just as the diamond point entered the specimen and recording stopped when the

specimen had traversed the entire distance under the diamond point. The vertical traces

represent the load calibrated in arbitrary units.

^Precise Milling Machine, Model II.

26
4. The trace on the oscilloscope was photographed.

5. The diamond point was examined for possible damage using a magnifier.

All of the present tests were performed using down grinding; that is, the table and

the diamond point were moving in the same direction during contact.

2.3. Characterization

Standard nomenclature for describing the features formed by the diamond points was previously adopted [23]. The term "scratch" refers to the entire result of scratching

or single point grinding and includes both surface and subsurface effects. The "track"

is the path formed by the diamond point in the surface. The individual striations in

the track are called "grooves."

The damage was characterized by optical and scanning electron microscopy (SEM).

Various methods were used to expose the subsurface damage. In some cases, the specimens

were beveled through the track (1-2° bevel). The beveling was done by holding the specimens at the desired beveling angle on a lap rotating from the edge of the bevel back

toward the track. This direction of rotation was used to avoid extending the subsurface

damage. The effect of beveling is to expose the subsurface damage along the beveled

surface. In other cases the specimens were notched from the surface opposite that con- taining the track. Then, the specimens were fractured back from the notches toward the

tracks, thereby avoiding propagation of the machining damage. The swarf and the changes

in the diamond points were also characterized.

The remaining strengths of the damaged specimens were measured in flexure by three

point loading on a one inch span at a loading rate of 20 Ns~l.

3. Results and Discussion

3.1. Scratching experiments

Glass plates, approximately 10.7 x 4.3 x 0.6 cm, were scratched using various loads.

In each case the specimen was scratched lengthwise along the center of the large face.

Then, the plates were cut crosswise to form bars and a large fraction of the bars were

fractured to measure the remaining strength after scratching. Tnus, the scratches were

parallel to the fracture surfaces. In almost all cases the fractures originated at the

scratches. The depths of damage were measured by viewing the damage through the sides of

glass bars using a microscope with a scale in the eyepiece. At a low load (9N) a very

smooth track was formed. At higher loads the tensile stresses behind the diamond point

generated characteristic U-shaped cracks. At the highest load (38N) lateral cracks were

formed on both sides of the track.

The loads (P) were plotted vs. the depth of damage raised to the three halves power

as suggested by eq. (1), in figure 3. Taking the slope as Kjr/x with Kic = 0.8 MNm-3/2

[28] yielded x

= 0.053. This estimate is higher than a previous value obtained by impact

tests [17]. The data are too sparse to prove that the depth of damage varies with load

as predicted by eq. (1) but the results are consistent with this possibility.

The variation of the average fracture stresses with the inverse square root of the

depth of damage was plotted based on eq. (2) and compared with theoretical curves for semicircular and edge notch flaws (fig. 4). The slope indicated by the data is close to that

for edge notch flaws, perhaps indicating that the scratches act as edge notches. The

average remaining strengths are plotted vs. the load on the diamond point in figure 5. Six to 10 results were averaged for each load. Results for a total of two specimens in

which the fractures did not originate at the scratches and three specimens that yielded

unusually high values were omitted from the averages. The fracture stress decreases

with increasing load on the diamond point as expected. The average remaining strengths

are plotted vs. P-l/3, as suggested by eq. (3), in fiqure 6. Although the range of the

data is limited, the results are consistent with a straight line through the origin.

27
40-

i 30

O

O

S 20

>

10

10 20

,3/2 _

30

m

3/2

x I07 40

iin

co

<u

o

o

o

c-l/2_ m

-l/2

Fig. 3. Vertical load vs. c

3//2 for qlass

scratched by a diamond point.

Fig. 4. Average fracture stress vs. C-V2

for glass scratched various loads,

~ 50

40

c

o> £ 30

o>c

'co

e 20 a> <r

a>a>| 10 ><

Q

Brackets Indicate Standard Deviations

_L

10 20 30 40

Vertical Load - N

50

M

I

40

C

<u

CO

enco

E

a> lT

a>a>ok_a>><

30

.E 20

10

0.

/

4

J 1 L

0 0.10 0.20 0.30 0.40 Q50

p-l/3_ N

-i/3

Fig. 5. Average remaining strength vs, vertical load for scratched

glass plates.

Fig. 6. Averaqe remaininq strength vs P-l/3 for scratched glass

plates

.

H.P. silicon nitride plates (4 x 14 x 0.3 cm) were qround and lapped. Then, the

plates were carefully scratched lengthwise by a diamond point under various loads and

cut crosswise to form bars. The depths of scratch damaqe were measured. As shown in

figure 7, the tracks do not contain deep grooves. The smoothness along the track may

indicate microplastic flow under the diamond point. A crack projects downward from the

track. Examination of a 1° bevel through a track shows that there has been considerable

loosening of the individual grains in the material under the track. This loosening led

to frequent pull out of the grains as the bevel was lapped.

The variation of the crack depth (C) with load on the diamond point is illustrated

in figure 8. The load is approximately linear with C3/2 as expected based on eq. (1)

althouqh it appears that a somewhat better fit would result from using a lower exponent.

Taking K IC = 4.7 MNm" 3 / 2 [29], the value of x was calculated from the slope of the curve in figure 8, yieldinq a value of 0.20.

28
Fig. 7. Track and subsurface damage in Fig 8. Vertical load vs. (Crack Depth)3/2

H.P. silicon nitride (vertical for H.P. Si 3N4 scratched with a load 38 N). diamond point.

3.2. Single point grinding experiments

3.2.1. Glass

Glass was machined by single point grinding in two series of experiments. In the

first series, the set deptn (the depth to which the diamond is set to penetrate the

glass) was varied from 25-100 ym, using "sharp" diamonds. Examination of the results

gave the impression that the observed damage was too severe to be realistic in terms of

practical applications. Therefore, in the second series smaller set depths, 13 and 25 ym, were used. Also, the sharpness of the diamond points was varied by using both sharp

diamonds and diamonds in which the tips were blunted to form flats 50 or 100 ym in

diameter.

The impact causes fragmentation of glass in the track and crack formation under and

alongside the track. The fragments have a bimodal size distribution. Large numbers of

very fine particles (1-5 ym) are produced in the grooves by crushing. Frequently, these

particles are packed against the sides of the grooves as the diamond point passes.

Smaller numbers of large fragments (40-400 ym) form by chippinq of material at lateral

cracks alongside the track. The shapes of these cracks and the resulting fragments are similar to those of lateral vent cracks observed to result from impact and "static"

indentation. Although investigators of machining damage have referred to these cracks

by several other names, they will be referred to as lateral or lateral vent cracks in

this investigation. The damage under the tracks was studied. Blunt diamonds form wider

tracks than those formed by sharp points. Figure 9 shows the damage under a track formed

by a 100 ym flat diamond at 25 ym set depth and 6000 rpm. The wide track formed by this

diamond and the wide separation of the radial cracks under the track are characteristic

of the damage due to blunt diamonds. Voids caused by removal of chips formed by lateral

cracks are shown on both sides of the track.

In some cases the forces at the start of the scratch were greater than those farther

along in the scratch. Apparently, the initial impact was sufficient to induce vibrations

in the wheel, shaft and mounting which resulted in these higher forces which were damped

29
surface damage in glass ground

Hat HiamnnH with 9^ nm colFig. 9. SuDSUrTdte uaiiidye in yidb:> yr ui with 100 ym flat diamond with 25 ym set

depth, 6000 rpm.

out as the grinding continued. The forces increased with increasing set depth, as expected. When the specimens were not quite flat, or were mounted not quite parallel to

the table travel, the depth of cut varied during the traverse with corresponding variation

in the forces. One might expect that, at constant table speed, increasing the wheel speed

would decrease the volume of material swept out by the diamond point on each rotation and

thus decrease the forces but this decrease in forces was not observed consistently with

glass. The forces increase with decreasing sharpness of the diamond point as expected;

that is, the forces increased in the series sharp point -> 50 ym diameter flat -* 100 ym

diameter flat.

In order to learn more about the variation of force with groove depth 5 and to explain

some of the observations in the previous section, the groove depths were measured at

intervals along some of the tracks yielding interesting results. The increase in groove

depth with vertical load is illustrated in figure 10. The variations of the groove depths

for tracks made at 4000 and 12000 rpm are shown in figure 11. Groove depths are, with few

exceptions, much greater than the set depths and the differences between groove depths and

set depths are greater at 12000 rpm than at 4000 rpm. This latter observation indicates

why the forces do not decrease consistently with increasing wheel speed. Evidently, the

extra force necessary to remove the added depth at higher wheel speed is offset to some degree by the decrease in the horizontal distance swept out so that whether there is an increase or decrease in the force in a particular case depends on the balance of these

opposing effects.

Fig. 10. Vertical load vs. groove depth

for single point grinding of

glass (set depth 50 ym, 4000

rpm, track no. 14)

.

5 It is important to remember that the groove

track and does not include the subsurface d

Fig. 11. Groove depth vs. distance from

start of track for single point

grinding of glass (set depth

25 ym, tracks 3 and 4)

.

is measured only to the bottom of the

30
For single point grinding with sharp diamond points the lateral crack lengths increase

with increasing set depth and wheel speed and, therefore, with increasing force. Because

the lengths of lateral cracks vary substantially even within short distances along a particular track, it is difficult to find an accurate way to represent the above variation. Therefore, it was decided to compare the lengths of the largest lateral cracks

observed for various grinding conditions. The results of this comparison are given in

figure 12. The increases in lateral vent crack length with set deoth and wheel speed,

shown in the figure, occur as a result of the increases in forces that occur with

increases in these variables.

Fig. 12. Largest lateral crack length vs. set

depth for single point grinding of glass (sharp

diamond points, 4000 and 12000 rpm).

'O 25 50 75 100

Set Depth -pm

At 6000 rpm and 13 ym set depth, the largest lateral crack length decreased with

decreasing sharpness of the diamond point as follows: sharp point, 420 ym; 50 ym flat,

290 ym; 100 ym flat, 150 ym. This suppression of lateral crack propagation with decreasing

sharpness of the diamond points was unexpected.

The penetration of damage was measured in glass subjected to single point grinding

with sharp diamond points at various wheel speeds and relatively large set depths. The

variation of the maximum crack depth raised to the three halves power with the vertical

40-

/

/

/

/i/o

/

/

h /

/

/.

• / 3/p

/ . Fig. 13. Vertical force vs. (crack depth) ' for single point grinding of glass with sharp

diamond points at various set depths and wheel

/<y' speeds.

50 100 150 200 250

m x „3/2__3/2 „ |q7

31
force is plotted in figure 13 which shows that the depth of damage increases with vertical

force as expected and that the relationship is approximately linear with C,3/2. The value

of x was calculated and found to be 0.10 which is about twice the value determined from

the scratching tests. This difference is attributed mainly to the fact that the diamond

points used for single point grinding had smaller tip radii than the points used for

scratching.

In a later experiment at smaller set depths the variations in forces and depths of

damage for sharp and 50 ym flat diamonds were investigated with the results shown in

figure 14. The forces were higher for the 50 ym flat diamond but the damage penetration

averaged slightly less than for the sharp point. The values of x were estimated and

found to be 0.14 for the sharp point and 0.038 for the 50 ym flat.

• Sharp Diamond —-°- — 50 jjm Flat

/

Fig. 14. Vertical load vs. (crack depth) 3/2 fo

single point grinding of glass with sharp and

50 urn flat diamond points at various set depths

and wheel speeds.

-ma,S£ x 10'

3.2.2. H.P. silicon nitride

There are significant differences in fragmentation and crack formation of glass and

H.P. Si 3N4. The higher Kjq of H.P. Si 3N4 tends to restrict fragmentation and crack formation compared with that in glass but this is partially offset by the greater forces

required to drive a diamond point through the H.P. Si 3N4 . The differences in crack formation are illustrated by the pictures of the tracks in figure 15. With sharp diamonds,

the tracks in glass and H.P. Si 3N4 are similar in appearance but the scale of the lateral

cracking in H.P. Si 3N4 is much smaller than that in glass. With blunt diamonds, at small

set depths and low wheel speeds the tracks in H.P. Si

- 3N4 are very smooth and there is

little or no chipping (lateral crack formation) alongside the track. As the scale of

the damage becomes smaller the response of the material tends to become less brittle,

as expected.

The total volume of silicon nitride fragments formed is much smaller than that of

glass. The fragments have a bimodal size distribution which consists of fine powder

formed by crushing and larger fragments formed by chipping. The shapes of the fragments

formed by chipping are similar to those observed in glass.

32
A. H.P. Si 3N4 , sharp diamond point, topwheel speed 8000 rpm, bottom-wheel speed

4000 rpm.

4000 rpm

B. Glass, sharp diamond point, top-wheel

speed 8000 rpm, bottom-wheel speed

4000 rpm.

D. Glass, 50 ym flat diamond, top-wheel

speed 8000 rpm, bottom-wheel speed

4000 rpm.

Figure 15. Comparisons of tracks formed in H.P. Si 3N4 and glass (set depth 13 ym).

The groove depths in the H.P. Si 3N4 specimens varied along their lengths as they

did in glass. At the initial impact site the groove depths were always much greater

than the set depths. As the midpoint of the track was approached, the groove depth

became almost equal to the set depth. In this respect, this case is different from glass

in which the groove depth was always substantially greater than the set depth. The groove

depths increased with increasing vertical force as shown in figure 16. Comparisons for

sharp, 50 ym flat and 100 ym flat diamond points show decreasing slope with decreasing

sharpness. The spacings between the curves are roughly equal to the differences in the

set depths.

Gielisse, Kim and Choudry [30] have shown that grinding forces decrease with increasing

wheel speed, pass through a minimum, and then increase at high wheel speeds. Lateral crack

formation responds to these variations in forces as shown in fiaure 17 for sharp diamond

points. As expected, the maximum lateral crack lengths decrease with increasing wheel

speed, pass through a minimum and then increase.

33
Set Depth 25pm Fig. 16. Vertical force vs. depth of groove for

single point grinding of H.P. Si^N, using 50 ym

flat diamond, wheel speed 8000 rpm?

j i L J_ JL 20 40 60 80 100

Depth of Groove -pm

800-

100-

Fig. 17. Maximun lateral crack length vs. wheel speed for single point grinding of

H.P. SiJi. (sharp diamond point, set depth

50 ym).

^

2000 4000 6000 8000 10000 12000

Wheel Speed-rpm

The maximum lateral crack lengths (C. ) for diamonds of varying sharpness were determined

for various grinding conditions. Then, the vertical force at each such lateral crack was estimated from the oscilloscope trace and plotted vs. the maximum crack length to the three

halves power as shown in figure 18. The principal result is that these lateral crack

lengths fall in the same range for the diamond points of varying sharpness despite the

large variations in load from one type of point to another. This result indicates that

the variations in x for the various geometries roughly offset the effect of the variations

in force for the three degrees of sharpness.

Substantial differences in the morphologies of subsurface cracks were observed in

H.P. Si 3N4 ground with points of varying sharpness as shown in figure 19. The pictures

show cross-sections formed by fracturing the specimens back toward the tracks. The shapes

of the grooves accurately reflect the shapes of the diamonds. The depths of penetration

increase with increasing sharpness and wheel speed. For sharp points, the most penetrating

34
damage usually consisted of a radial crack perpendicular to the surface and directly under

the groove. For flat diamonds, cracks usually radiated from the corners of the grooves.

In many cases these cracks were the deepest damage.

250

200

0>o

i? 150

DO

<u > 100

• Sharp diamond —o— 50pm flat

x 100 ym flat

/

/

/

/

/

Fig. 18. Vertical force vs. (max. lateral crack

length) 3 /

2 for single point grinding of H.P.

Si^N^ with diamond points of varying sharpness.

20 40 60 80 100

c? /2 -m3/2x I07 120

Fig. 19. Variations in subsurface damage with diamond shape and wheel speed; H.P.

Si 3N4 , set depth 13 ym.

In addition to radial and lateral cracks, two other types of cracks were observed in

specimens ground with flat diamonds. Surfaces of fractures parallel to the track contain

a succession of small conical (Hertzian cone crack) surfaces, one of which is shown in

figure 20. The cone crack continues under the track from the exposed conical surface.

The second type of crack was usually observed in fracture surfaces perpendicular to the

tracks and has a bulbous appearance (fig. 21). No explanation for the formation of these

bulbous cracks is available.

35
Fig. 20. Evidence for Hertzian cone crack

damage in H.P. Si 3N4 ground with

TOO ym flat diamond (25 ym set

depth, 8000 rpm).

Fig. 21. Bulbous crack formed under

tracks in H.P. Si 3N4 (50 ym

flat diamond, 25 ym set depth,

4000 rpm).

Composite photographs of the beveled portions of tracks formed in H.P. Si 3N4 by

diamond points of varying sharpness were presented previously [31]. The increasing track

width with decreasing sharpness of the diamonds was clearly shown. The groove depth

decreased with decreasing sharpness of the diamonds. This decrease was observed by noting

the depths of the smooth portions of the track that are formed by shearing of material at

the bottom of the track. The results were 26 ym for the sharp point, 23 ym for the 50 ym

flat and 12 ym for the 100 ym flat. Arc-like cracks were observed on the beveled surface

under the track. These cracks are believed to be associated with the conical or bulbous

cracks noted previously. The horizontal intervals between these arc-like cracks were estimated and were found, on the average, to correspond to the distance traveled by the

table between revolutions of the diamond. Thus, the material responds to the individual

impacts as originally hypothesized.

The vertical force vs. the depth of damage to the three halves power was plotted with

the results shown in figure 22. This figure shows strikingly that x in eq. (1) varies

with the sharpness of the diamond point. It is clear from this figure that, in practical

applications, in which diamond points or abrasive grains become less sharp during use, the

tendency toward damage penetration decreases despite the increased load on the diamond

point as it wears. Busch and Prins [32] have shown that, in multipoint grinding of a hard ceramic, initially sharp diamonds are rapidly dulled. In view of the reduction in

damage penetration, further investigation of the use of dull diamonds for grinding

ceramics, is justified. The values of x were estimated for the three sets of data with

the following results: sharp diamond, x = 0.46; 50 ym flat, x = 0.13; 100 vim flat,

x = 0.065. These data show that x depends strongly on sharpness.

It is not uncommon for diamond points to break, especially when used at large set

depths and low wheel speeds. Therefore, one might question whether or not the diamond

points maintained their integrity in these tests. Perhaps the most important tests were

those involving the points of varying sharpness, so one set of these points, exposed to

lengths of cut ranging from 56-73.5 cm at set depths of 13 and 25 ym was characterized.

The points were examined under the optical microscope with reflecting light perpendicular

to the points. The reflecting areas were photographed and the dimensions of the flats were measured on the photographs. The 50 and 100 ym diameter flats remained close to

their original dimensions. A firmly adhering layer of silicon nitride that obscures

much of the surface of each of the diamonds was observed by SEM. As indicated by the

36
—• Sharp

o 50 pm flat

* lOOpmflat

3/2

Fig. 22. Vertical force vs. (crack depth) ' for H.P. Si^N* ground with diamond points of

varying sharpness (combined data 4000 and 8000

rpm, 13 and 25 ym set depth).

J 1 I L 20 40 60 80 100

c3/2_ m

3/2 x|07

intense white light emitted from the contact zone, high temperatures are generated by the

single point grinding [33]. The silicon nitride contains a low melting magnesia and silica

rich intergranular phase that would be expected to melt at these temperatures. The

presence of this viscous material may account for the adhesion of the silicon nitride to

the diamonds. Despite the fact that the surfaces are obscured, it is clear that the

diamonds have retained their original shapes. Therefore, it can be concluded that, at

least for moderate set depths, the diamonds maintain their integrity in machining H.P.

Si3N4-

4. General Discussion

For diamond scratches in glass, the depth of damage increased with increasing load

as expected. Although the results are sparse so that a definite relationship cannot be

specified, the variation of the depth of damage (C) with load (P) is consistent with

the C3/2 dependence expected based on eq. (1). The remaining strength increases with

C-l/2 as expected based on eq. (2) and the slope of the curve was close to that expected

if the scratch acts as a notch. Looking at the overall phenomenon, the remaining

strength increases with P~ 1/3 as expected based on eq. (3).

For diamond scratches in H.P. Si 3N4 , the depth of damage also increases with increasing

load, as expected. The results were plotted assuming the Cj/2 dependence but it appears

that a better fit could be obtained with a lower exponent.

In single point grinding of glass, the damage increased with increasing load on the

diamond point. Groove depths were measured and were found to increase strongly with

wheel speed. Therefore, the forces do not depend simply on the volume swept out by the

diamond point on each revolution. Instead, the force depends in a complex way on the

volume swept out by the diamond point which decreases with increasing wheel speed and the

extent of damage outside this swept out volume which increases with increasing wheel

speed. These opposing effects offset each other to some degree so that the variation

of the force depends on the balance of these effects in a particular case. At low wheel

37
speeds the volume swept out is the dominant effect so that the force decreases with

increasing wheel speed. At higher wheel speeds the extent of damage outside this swept

out volume becomes dominant so the force passes through a minimum and increases at high

wheel speeds.

The extent of damage reflects the factors discussed above. The lateral crack lengths

increase with increasing set depth and wheel speed and, therefore, with increasing force.

The depth of penetration of damage under the track increased with force on the diamond

point, as expected. The variation is approximately P « C,3/2 as in eq. (1). The forces

increase with decreasing sharpness of the diamond point but the damage penetration does

not increase correspondingly. The effect of this is that the values of x decrease to

compensate for the lack of the expected increase in C3/2.

In single point grinding of H.P. Si'3N4, the groove depth also increased with the

vertical force as it did for glass. However, near the midpoint of the track, after the

vibrations due to the initial impact have decayed, the set depth and groove depth were

very nearly the same in most cases.

The extent of the damage increased with the vertical force, as expected, for both

the lengths of lateral cracks and penetration of damage under the track. For individual

diamond points the variation in the extent of damage was consistent with P « c3/2 based

on eq. (1). The vertical forces increased with decreasing sharpness of the diamond and,

again, this increased force was not accompanied by a correspondinq increase in 0^/2. The

damage under the tracks formed by the 50 and 100 ym flats differed somewhat from that due

to sharp diamonds in that the radial cracks frequently emanated from the corners formed by

the edges of the flats. In addition, there was evidence of two other types of cracks;

Hertzian cone cracks and what we have called bulbous cracks. The frequencies of occurrence

of these cracks are consistent with formation of the cracks in response to the individual

impacts of the diamond points.

The values of x for penetration of damage under the track are qiven in table I for

various materials and test conditions. The strong dependence of x for single point

grinding on the sharpness of the diamond is evident for both glass and H.P. Si'3N4. The

differences observed in values of x from scratching and single point grinding tests are

also consistent with the differences in the diamond point radii used for the two methods.

Table I. Estimated values of x for various materials and test conditions.

Glass H.P. Si

3 N 4

Scratching Test

Sharp (^1 50-300 ym tip radius) 0.053 0.20

Single Point Grinding

Sharp M0 ym tip radius) 0.10,0.14 0.46

50 ym Flat 0.038 0.13

100 ym Flat — 0.065

The values of x for H.P. Si 3N4 are consistently about four times the comparable values

of x for glass. This consistent difference indicates that some factor that is important

for abrasive machining was not taken into account in the derivation of eq. (1). The

principal difference in the response of glass and H.P. Si 3N4 to abrasive machining seems

to be the more extensive crushing of the glass. It seems reasonable that, if the diamond

point acts through a layer of crushed material, it will act as though it is less sharp

than would otherwise be the case. The lower values of x observed for qlass are consistent

with this explanation. Decreased penetration of damage in the presence of crushing was observed previously in a study of damage due to impacts by spheres [14]. This earlier

observation is also consistent with the proposed explanation.

38
The consistent decrease in x with decreasing sharpness of the diamond point suggested

that it might be possible to find an explanation for the variations in x- The principal

damage to be accounted for is the radial cracks which are extended in increments parallel

to the track with each rotation of the diamond. Physically, this situation is somewhat

different from that for which eq. (1) was derived. Lawn and Marshall [34], derived

an equation for an idealized straiqht-crack model for indentation induced failure. It was assumed that once a straight crack was developed, it is driven by essentially localized, mouth-loading forces representing effective wedging components of the applied load.

The result for straight cracks is

P

L P K

IC

= X

L ^172

=

X

L j^TTT (4)

where P = P|_ L is the load on an indenter of length L and xj_ "is again a dimensionless

contact constant which incorporates uncertain geometrical and frictional factors. In our case the contact lengths of the flat diamonds were taken as the diameters of the

flats. There was fractographic evidence of the contact width of the sharp diamond points

in the tracks formed by single point grinding. This width is determined by the combined

effect of elastic and microplastic deformation under the diamond point. The contact was assumed to be symmetrical so that the length of contact could be taken equal to the width.

The widths of the smooth contact regions in a small number of specimens were examined.

On the basis of these measurements a contact length of 20 ym was selected. The results

of these calculations are plotted in figure 23 which shows that the assumption of line

contact loading as described above allows the data to be represented roughly by a single

line with a value of xl = 0.030.

iE

- 4

I" •

c

=>

k.<UQ. T3q 2 o

o

> I

Sharp Point 50 pm flat 100 pm flat

.o3f- y Fig. 23. Vertical load per unit contact / length vs. (depth of damage 1 /

2 for single

, point grinding of H.P. Si 3 N 4 with diamonds

° of varying sharpness

o

J I I L 0.005 0.010 0.015 0.020 0.025

c

l/2 -ml/2

5. Conclusions and Recommendations

The types and extent of damage observed in this investigation depended strongly on the diamond shape, set depth and wheel speed. Sharp diamonds induced radial cracks

perpendicular to the surface and lateral cracks. Flat diamonds induced radial cracks

emanating from the edges of the contact zone and lateral cracks which usually were shorter

than those induced by sharp diamonds. In addition, Hertzian cone cracks and so-called

bulbous cracks were observed in several cases in tracks formed by flat diamonds. As the

grinding conditions became less severe and the contact zones were smaller, there was

39
evidence that the response of the materials became less brittle.

The penetration of the damage was analyzed initially in terms of eq. (1) and it was found that the values of x decreased with decreasing sharpness of the diamond point.

Subsequently, it was assumed that the mechanism of damage penetration involved line contact

loading which was represented by eq. (4). This assumption allowed the results for H.P.

Si^N^ to be represented by a single value of xThe penetrations of damage in H.P. Si 3N4 and glass were compared. The values of x for glass were approximately a factor of four smaller than those of H.P. Si^^ indicating

that the Kjc does not account for all of the differences between the materials. Therefore,

the mechanism of damage penetration in glass must be somewhat different from that in H.P.

Si 3N4. One obvious difference is that glass crushes much more extensively than H.P. Si'3N4.

Therefore, the diamonds may act to wedge open the cracks through a layer of crushed glass

making each diamond act somewhat less sharp than otherwise expected. The differences in

x between glass and H.P. Si^^ are consistent with such an explanation.

It is highly desirable to investigate less severe grinding conditions, multipoint

grinding and other materials. Use of less severe grinding conditions may yield a better

description of the role of microplastic processes in abrasive machining. Further advances

should make it possible to predict the penetration of surface damage for each set of

abrasive machining conditions. This information would be very useful for prediction of

the strengths of machined ceramics and for the development of machining processes in which

each successive step is designed to remove the damage of the preceding step so that

machining time required using fine-grained abrasives can be minimized.

Acknowledgment

The authors are pleased to acknowledge the advice and assistance of Professor P. J. Gielisse and Mr. Ray McLaughlin of the University of Rhode Island and the sponsorship of

the National Science Foundation.

References

[1] The Science of Ceramic Machining and Surface Finishinq, S. J. Schneider and R. W. Rice, eds. NBS Special Publication 348 (1973).

"

[2] Gielisse, P. J. and Stanislao, J., Mechanical methods of ceramic finishing, Reference

1 , pp. 5-35.

[3] Oh, H. L., Oh, K. P. L., Vaidyanathan , S., and Finnie, I., On the shaping of brittle

solids by erosion and ultrasonic cutting, Reference 1, pp. 119-132.

[4] Gielisse, P. J., Kim, T. L., Goyette, L. F. and Nagarkar, R. V., Fundamental and

practical evaluation of the ceramic finishing process, U. of Rhode Island Final

Technical Report, Contract N00019-72-C-0202 (1974).

[5] Mecholsky, J. J., Jr., Freiman, S. W., and Rice, R. W., Effect of grinding on flaw

geometry and fracture of glass, J. Amer. Ceram. Soc. 60 [3-4] 114-117 (1977).

[6] Rice, R. W. and Speronello, B. K., Effect of microstructure on rate of machining

of ceramics, J. Amer. Ceram. Soc. 59 [7-8] 330-333 (1976).

[7] Rice, R. W. , Machining of ceramics, in Ceramics for High Performance Applications . J.J. Burke, A. E. Gorum and R. N. Katz, eds., pp. 287-343 (Brook Hill Publishing

Co., Chestnutt Hill, Mass., 1974).

40
Lawn, B. and Wilshaw, R. , Indentation fracture: principles and applications, J. Mater

Sci. 10 1049-1081 (1975).

Wiederhorn, S. M. and Lawn, B., Strength degradation of glass resulting from impact

with spheres, J. Amer. Ceram. Soc. 60_"[9-10] 451-458 (1977).

Evans, A. G., Strength degradation by projectile impacts, J. Amer. Ceram. Soc.

56 [8] 405-409 (1973).

Evans, A. G. and Wilshaw, T. R. , Quasi-static solid particle damage in brittle

solids, I. Observations, analysis and implications, Acta Met. 24, 939-956 (1976).

Evans, A. G. and Wilshaw, T. R., Dynamic solid particle damage in brittle materials:

an appraisal, J. Mater. Sci. 12, 97-116 (1977).

Adler, W. F. and Hooker, S. V., Impact damage in single crystals, Effects Technology

Report No. TR 77-49, Contract N00014-76-C-0744 (1977).

Kirchner, H. P. and Gruver, R. M. , Localized impact damage in glass, Mater. Sci. and

Eng. 28, 153-160 (1977).

Kirchner, H. P. and Gruver, R. M., Localized impact damage in stressed members, Mater

Sci. and Eng. 28, 249-255 (1977).

Kirchner, H. P. .and Gruver, R. M. , Branching of Hertz Cracks, from Fracture 1977 , Vol. 3, D. M. R. Taplin, ed. pp. 959-964, (University of Waterloo Press, Waterloo,

Ontario, 1977).

Kirchner, H. P., Analysis of localized impact damage in biaxially stressed glass

plates, J. Amer. Ceram. Soc. 61 [3-4], 161-163 (1978).

Kirchner, H. P. and Gruver, R. M., The effect of localized damage on energy losses

during impact, Mater. Sci. and Eng. 33_, 101-106 (1978).

Kirchner, H. P. and Gruver, R. M. , Localized impact damage in stressed Al 2O3 and

Si3N 4

, Mater. Sci. and Eng. 34, 25-31 (1978).

Kirchner, H. P. and Gruver, R. M. , Localized impact damage in a viscous medium (glass

from Fracture Mechanics of Ceramics , Vol. 3, R. C. Bradt, D. P. H. Hasselman, and

F. F. Lange, eds., pp. 365-377 (Plenum, New York, 1978).

Kirchner, H. P., Gruver, R. M. , and Walker, R. E., Strength Effects Resulting from

Simple Surface Treatments, Reference 1, pp. 353-363.

Gruver, R. M. and Kirchner, H. P. Effect of surface damage on the strength of Al 2O3

ceramics with compressive surface stresses, J. Amer. Ceram. Soc. 56 [1] 21-24 (1973).

Gruver, R. M. and Kirchner, H. P., Effect of Environment on penetration of surface

damage and remaining strength of Al 2O3 , J. Amer. Ceram. Soc. 5_7 [5] 220-223 (1974).

Marshall, D. B., Lawn, B. R. , Kirchner, H. P., and Gruver, R. M. , Contact-induced

strength degradation of thermally treated alumina, J. Amer. Ceram. Soc. 6]_ [5-6]

271-272 (1978).

Lawn, B. R. and Fuller, E. R., Equilibrium penny-like cracks in indentation fracture,

J. Mater. Sci. 10, 2016-2024 (1975).

Tada, H., Paris, P., and Irwin, G., The Stress Analysis of Cracks Handbook , (Del

Research Corporation, Hellertown, PA 1973)

Smith, F. W. , Emery, A. F., and Kobayashi, A. S., Stress intensity factors for

semi-circular cracks, Part 2, semi-infinite solid, J. Appl . Mech. 34, Series E,

453-959 (1967).

41
[28] Wiederhorn, S. M. , Subcritical crack growth in ceramics, from Fracture Mechanics

of Ceramics , Vol. 2, R. C. Bradt, D. P. H. Hasselman and F. F. Lange, eds. (pp.

613-646, Plenum Press, New York, 1974).

[29] Petrovic, J. J., Jacobson, L. A., Talty, P. K. , and Vasudevan, A. K. , Controlled

surface flaws in hot pressed Si*3N4, J. Amer. Ceram. Soc. 58 [3-4] 113-116 (1975).

[30] Gielisse, P. J., Kim, T. J., and Choudry, A., Dynamic thermal aspects of

ceramic processing, University of Rhode Island Final Technical Report, Contract

N00019-71-C-0244 (1971),

[31] Kirchner, H. P., Gruver, R. M., Richard, D. M., Fragment formation and damage

penetration during machining of ceramics, Ceramic Finishing Company Technical

Report, Contract NSF-C-APR77-19818 (1973).

[32] Busch, D. M. and Prins, J. F., A basic study of diamond grinding of alumina,

Reference 1, pp. 73-87.

[33] Gielisse, P. J., Kim, T. J., Choudry, A., Short, J. F., and Turker, E. J., Force

and wear analysis in ceramic processing, University of Rhode Island Final Technical

Report, Contract N00019-72-C-0202 (1972).

[34] Lawn, B. R. and Marshall, D. B., Contact fracture resistance of physically and

chemically tempered glass plates: theoretical model, Phys. Chem. Glasses 18 [1]

7-18 (1977).

42
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SINGLE-POINT SCRATCHES AS A BASIS FOR UNDERSTANDING GRINDING

AND LAPPING

A. Broese van Groenou, N. Maan and J.B.D. Veldkamp

Philips Research Laboratories

5600 MD Eindhoven, The Netherlands

Considering grinding as a mult i- scratch process we have used single-point, lowspeed, scratch experiments with pyramidal diamonds in an attempt to understand

such quantities as the force ratio and the specific energy in terms of the shape

of the diamond and of the properties of the workpiece, here ferrites, glasses

and various oxides. The force ratio is found to be dominated by the geometry of

the diamond, while the interface friction is small. The specific energy for

groove formation is determined by the Vickers hardness. A depth dependence ("size

effect") occurs with the harder materials. Moreover, scratch experiments yield

information on the resistance of cracking on a microscopic scale and therefore

give some idea of the relative strength of the grain boundaries with respect

to the grains

A liquid environment of e.g. n-alcohols does not have a pronounced influence on the force ratio and the specific energy for grooving, in contrast to water. Experiments made with a loose particle between a binder material (brass, pitch, steel)

and an oxide or a metal shows that the particle motion depends on its shape and

on the hardness of the two materials. Rolling occurs where the hardness values

are about equal. For a spherical particle a high load favours rolling, deforming

both materials to a moderate extent. Sharp particles behave in an irregular manner and scratch both materials. Finally, scratching is a useful means of testing the

homogeneity of materials.

Keywords: Ceramics; chipping; crack formation; fracture toughness; friction;

hardness; liquid softening; oxides; scratching; size effect; specific energy.

1. Introduction

Grinding, lapping and abrasive polishing of brittle materials like ceramics and glass

are often necessary but complicated processes. One essential element of the grinding system

is the interaction between the abrasive particles and the workpiece: here the material re- moval takes place and here the residual deformation is formed. This can be seen in an experiment in which a slightly tilted workpiece is ground: where the wheel first touches the

sample many scratches are seen, which at greater depth influence each other (fig. 1).

The basis for understanding grinding is therefore the scratching action of sharp points, and

a model may be built on the phenomena caused by a single point. In our study such scratches were made at a low speed (less than 1mm/s). Apart from making experiments somewhat easier,

this speed range is expected to give simpler results than when scratching at speeds used in

grinding, e.g. 30 m/s, where heat generation may in any case cause complications. The

low-speed experiments provide a basis for the high-speed problems, which must be tackled

separatedly. In this paper attention is focussed on the behaviour of various materials - glass, crystals, polycrystals - under the action of a sharp pyramidal point. Experimental

details are given in the next section.

Two modes of deformation are of interest in scratching: the plastic deformation underneath and the brittle fracture around the point, as can be seen in the figure. For

both phenomena it appears that the scratching speed is important, at the low speeds used

here. For the plastic deformation it should be kept in mind that a sharp point causes a considerable strain rate, of the order of £ = vs /b. This is in the range from 0.1 to 103 sec

- '' for a scratching speed vs of 1— 1 03 ym/s and a scratch width b of several microns.

43
Fig. 1. Photograph of scratches on polycrystalline Mng # ^jZnQ _ 2^e 2 .08®k mad- e ^>J a number

of diamonds on an electroplated wheel , running at 75 m/s . Since the sample was slightlytilted, the start of the scratches could be made visible.

' -\

A slow-scratch apparatus therefore allows the investigation of plastic behaviour over a reasonable range of £, although the deformation is limited to small volumes. This speed

range is also of interest for the second mode of deformation, crack propagation, because

for low stresses slow crack growth has been observed in glass and AI2O3 [ 1-k]

Returning now to high-speed grinding, the connection with scratching is found in that

both processes remove material and both need a certain mechanical action, expressible in

force and energy. The characteristic quantities are usually taken to be the force ratio,

i.e. the ratio of the force component parallel and normal to the surface, and the energy

necessary for removing a unit volume (for that reason it is called the specific energy)

[ 5-7] • Low-speed scratching may clarify the precise meaning of these quantities; for ex- ample, what is the relation between force ratio and the Coulomb friction in the interface

between particle and workpiece? Does the type of abrasive particle influence the process?

And is the specific energy typical of a material or of a grinding system? For the last

question it is important to realize that the specific energy depends closely on the mode

of deformation: removing material by plastic deformation requires a much higher energy

than by brittle fracture (section 3). This is one reason for the interest shown in the

types of cracking that accompany a scratch; another reason is that the presence of microscopic cracks weakens a product. Cracking during scratching in glass was studied by Peter

[ 8] , who noted that the "median" crack extended further not only at a higher load but also

at a lower speed. For other types of cracks, to be classified in section k, load-speed

diagrams can be given [9,10] . Two types of cracks have recently been investigated theoretically: fracture near an indentation [ 11,12] and around a scratch [9 1 • It appears

that the propagation of cracks on a microscale allows a description with reasonable accuracy, in spite of the unknown stress field around a pyramidal indenter.

Scratching may also be useful in the study of the possible influence of water and

other liquids on material behaviour, leading to better insight into their role in the crea- tion of surface damage. It is well known from the work of Westbrook [ 13] and Westwood [ 1V|

that the hardness of a number of materials (e.g. LiF, MgO, A12 0, glass) varies with time

in a suitable liquid environment. A scratching experiment allows much shorter observation

times (Westwood used 1000 sec.) and moreover it is possible to verify a lubricating action

at the contact between particle and workpiece from the value of the force ratio (section 5)«

The action of water on the propagation of slow cracks in glass is well known [ 1] ; in the

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

44
same speed range scratches can be made [ 8] and slow crack growth observed. Although such

low speeds are not directly relevant to grinding, they may influence the residual damage

in a ground surface. Also, for slower processes such as lapping, liquids may play a role.

Lapping is a process where material is removed by loose particles. How "loose" these

particles really are is also a subject that can be investigated in a single-point scratching apparatus. A variety of lappan materials can thus be investigated (section 6).

Leaving aside the other applications of a scratching apparatus, we discuss the relevance

of scratching data to grinding in section 7-

Scratches were made by moving a polished sample under a vertically loaded diamond.

During the motion the horizontal force on the diamond was recorded. Afterwards the depth

and width of the scratch were measured under a microscope. (Details have been given in

ref. 15.) The apparatus, shown in fig. 2 and 3, contains a balance to enable scratching to

be carried out at loads as low as 0.01 N. The maximum load is about 3N. The sample-holder

is driven by a motor at speeds in the range from O.k to kOO um/s. The diamonds were ground

into square pyramids of about k mm size, with apex angles of 2 9 = 100°, 120° and 136°.

The radius of curvature of the tips was bettern than 0.3 um as can be checked in a darkfield microscope. A tungsten carbide indenter was ground to the same shape (20 = 120°),

but could only be made with a radius of curvature of about 1 um, owing to its microstructure. During scratching in air the point and its reflection in the surface were observed

with the aid of a microscope connected to a TV monitor.

Scratches were made in a liquid environment by surrounding the diamond with a droplet.

The same apparatus could be used for observing loose particles by replacing the diamond

holder by a flat plate.

The standard three-point bend test was used in order to determine the critical stress

intensity factor, Kj c

. The double torsion test was used for determining the crack speed

dependence of Kj.

A range of polycrystalline ceramic materials was used (various ferrites, ZnO, AI2O3)

as well as crystals of MgO, ZnO and lead-hexaferrite . A number of glasses were also investigated. All samples were highly polished with diamond abrasives. For the etching of the

MgO crystals Stokes ' recipe was followed [ 16] .

Fig. 2. Scratching apparatus Fig. 3. Schematic drawing of the scratching

apparatus. Al, A3: upper arms, A2 : lower arm;

S1, S2: leaf springs; V: counter weight and

damper; C, F: coarse and fine vertical adjustments; L: load; D: diamond; BE: bending

element used for force measurements; S: sample;

Mo: motor; E: x-y and rotation stages; Mi:

micromanipulator.

2. Experimental

45
3. Grooves

3.1. Force ratio

Examples of scratches on a MnZn ferrite are shown in figure k, for various loads,

at a speed of 1+00 um/s and an apex angle of 136°.

Fig. k: Influence of load on the presence of crack types in crystalline materials. View

from above of scratches made on Mn Zn ferrite. Scratching speed: hOO um/s. Loads: 0.1 (a),

0.7 (b) and 3 N (c). Incident light, bright field.

Grooves are found at all loads, whereas cracks become visible at higher loads. During

scratching the horizontal force is recorded (fig. 5). At low loads a much smoother curve

is found than at higher loads. The average value of F can be used to calculate the force

ratio

f=F

t

/Fn

. (1)

f values have been determined for three pyramids and six materials [ 15] . The results allow

a simple presentation in a f - cot 0 plot, where 29 is the apex angle of the pyramidal

particle (fig. 6). Independent of load or material, f is equal

t(s)

Fig. 5. Recording of horizontal force component, Ft , for a load of 0 . 1 N (bottom) and

0.5 N (top) on MnZn ferrite.

Fig. 6. Force ratio f = F^./Fn

as a function

of cot 9 for various loads on the diamond. Al

materials are included (from ref. [ 15] ).

to cot 9 in a good approximation. The same result has been obtained for a number of metals

[ 17] •

The significance of this result can be seen from the directions of the force components

on the contact plane between diamond and workpiece (fig. 7). The vertical load Fn and the

horizontal component F+ combine to give the total force F-tot ( we consider scratching by a leading plane, so there is no resultant component in the third direction). The components

of F-^Q-f. along the contact plane F^ and perpendicular to it, F^ can be determined for a known value of 9. The experimental results show that F^/F^, i.e. the friction coefficient

on the contact plane, is much smaller than the force ratio, F^/Fn

. Denoting F// /Fj^ by

46
y = tan a, we can write f as

The ]i values are 0.1 or smaller. For large 6, i.e. blunt particles, the u value is more

important since cot 8 is small. In the limit of a sphere, at very shallow penetration,

f = u. For deeper sphere-scratches a correction is needed [ 1 T » 1 8] .

Fig. 7. Diamond geometry showing the contact

area AB. The load, Fn , and the measured tangential force, F^, are combined to form the

total force F+ 0

^.. The components perpendicular (Fj_ ) and (F^ ) to the contact area can easily be deduced. The force ratio f equals

tan The friction coefficient is y = tan a

Fig. 8. Force ratio versus load for two

pyramid materials, diamond and tungsten

carbide, for scratches on MgO crystals.

The geometry of the scratching point is not only important for a diamond pyramid. For

a tungsten carbide point with 28 = 120° similar results are found, e.g. on MgO (fig. 8).

Below a certain load the force ratio depends on the load, i.e. on depth of penetration.

This effect corresponds to the larger tip radius of the tungsten carbide point, which has

a varying 9 value and therefore an increasing f until the depth is larger than the radius.

With our diamond tips such an f-load variation has frequently been an indication of wear

on the tip. After regrinding the load dependence had disappeared. Experiments with roundedoff cones, as used in gramophones, are not suitable for a simple interpretation for this

reason. The influence of rounding-off has also been noticed by Prins [ 6] .

The horizontal force, as given in figure 5> shows an increasing number of fluctuations

with increasing load. Part of these are due to irregularities in the workpiece material,

such as pores or grain boundaries. The load dependence is due to the occurrence of various

cracks in and around the grooves, which will be discussed below. First the type of deformation will be treated and a quantitative characterization given.

3.2. Deformation

During scratching the diamond moves through the material (cf. fig. 7) and pushes

the debris up in front, as can be seen on the microscope picture. In exceptional cases the

debris is still fastened to the groove [7] (fig. 9), in most cases it consists of short

pieces of material and lies around the groove. The groove itself is a highly deformed area. From its depth d and width b, as determined under a microscope after scratching, the groove

cross-section is found to be congruent with the diamond pyramid [ 15] ; for grooves on some metals, however, the depth/width ratio was too low [ 17] • A slight pile-up of material along

the groove has been found in the softer ceramics.

47
Fig. 9. Glass ciarls fastened to groove. Fig. 10. Dislocation pattern on cleaved MgO,

perpendicular to scratch (28 = 13(P;Fn

= 0.2N

etched; horizontal and vertical cracks under

the groove

.

Fig. 11. Scratch on Gd

3 Ga

5 0 12 crystal (F = 0.05N). Scratch length is 100 ym. On the

left a microscope picture of the scratch, on the right an X-ray topograph of the same area.

The high pressure on the contact plane leads to dislocation generation and movement. Dislocations are found deep under the surface of crystalline samples; an example is given in

figure 10 for MgO. This result was obtained after cleaving perpendicular to the scratch

and etching; the dislocations along [ 110] extend down to 100 ym for a load on the diamond

of 0.1 N. Even further from the scratch one can detect the stress field corresponding

to the plastic deformation near the groove by means of X-ray topography (fig. 11). In

polycrystalline materials the deformation on cleaved surfaces has never been found at

large depths, presumably because the grain boundaries formed a barrier.

3.3. Scratching hardness, specific energy

More quantitative statements can be made about grooving from the measured depth and

width and the force components, by introducing the scratching hardness,

kF

b

48
and the specific energy,

e = 2F

bd~

In the "ploughing" model [ 15, 19, 20] these quantities are related to the contact stress

a, derived from the total force and the contact area (AB in fig. 7), and the friction

coefficient u by

e = o{ 1-u tan 0

)

H

g = a( 1+y cot 6).

(5)

(6)

200

6

100

50

20

i/) JO

F„ =01

28

' 136" V

« 120"

- • 100" /

X /

- •/'

100 200 500 1000 2000 5000

»- Hardness Hv (kg/mm2

)

Fig. 12. Specific energy values for a range of

material^ as a function of their Vickers hardness

(1 kg/mm = 10 N/m ). Load: 0.1 N; three pyramid

have been used [ 15].

In our experiments the u and 6 values are such that the second terms are only minor cor- rections. This is confirmed by the experimental result that, within our accuracy, e and

H

s are equal. The results for a number of ceramic materials are given for a low load,

F_ = 0.1 N, in fig. 12 as a function of the Vickers hardness, measured on these materials

A good correlation is found, the slope being about 2.5. This indicates that the formation

of a scratching groove is closely related to "the plastic yield that occurs in indentation.

3.U. Size effect

The specific energy values do not depend on the shape of the pyramid, but on load, or which is more interesting, on depth. Such a size effect is well known in grinding metals

[21] and occurs also in grinding ceramics [22-24 • Experimental data for low-speed scratching are given in figure 13a for some ceramics and in 13b for several ferrites

200

2 5 10 d (fjrn)

WO

50

20

01 02

x SrFeuOig

* (Mr,,Zn)Fe2 0t

» (Ni,Zn)Fe2 0i

Fig. 13. Specific energy, e, versus depth of scratch, d, for all apex angles. For clarity

of presentation the white ceramics are shown in (a) separate from the black ceramics (b).

The values labelled steel (T)(a) have been taken from grinding results of Brecker and Shaw

([21] ). The black squares for AI2O3 refer to measured d values, the open squares have d

calculated from b.

49
il

)

It is remarkable that the decrease of e with d is greater for harder materials. The porous

ZnO sample (Sy. = 2.109 N/m2

, p = 8%) shows a small size effect; this is not due to porosity,

because for a crystal with H = 10.1 09 N/m2 we found a small size effect too. v

The influence of the scratching speed has been determined for a MnZn ferrite, MgO and

a number of glasses. The results show a clear increase of scratching hardness and specific

energy with speed in the range between 0.1 and 1+00 Urn/ s (fig. For metals a similar

increase has been found [ 17]

20

- 15

£ 10

- 5

-

Indentation » Scratching •

o Glass

MnZn ferrite

Fig. 14. Scratching hardness as a function of

scratching speed (v ) for glass and MnZn ferrite.

_l I I

1 1 L_ 10"' 10"' 1 10' 102 Scratching speed vs Ipm/s)

k. Cracking

Glass

Cracks of various kinds occur along and underneath a scratch. Examples have already

been given for a MnZn ferrite (fig. k) and for MgO (fig. 10). Before returning to these

materials, some attention will be given to the glass-like materials, which are simpler

(no anisotropy, no grain boundaries). Surface cracks (fig. 15) and sub-surface cracks

(fig. 16) in a NaCa glass allow the following inventory:

•

1

I

2 7

Fig. 15: Influence of load on the presence

of crack types in glass. Views from above

of scratches made on SF58 glass, scratching

speed kOO um/s. Loads: 0.1(a), 0.5(b) and

3 N(c). Incident light, bright field.

Fig. 16: Cross-section of a scratch in glass

The median crack is clearly shown. The open

space is bounded by lateral and sub-surface

cracks. Load 0.5 N scratching speed ^00um/s.

Incident light, bright field.

1. On the surface cracks run from the groove at an angle of about 30 . These cracks, which

are lateral with respect to the scratching direction, start in front of the scratching

pyramid and move away from the groove while the pyramid continues its motion. The "lateral"

cracks show some resemblance to the horseshoe shapes found around scratches made by spheres

[25] and to the cracks [ 11,12] underneath a static pyramidal indenter. In a film on scratching made by Busch [ 26] the speed of propagation of the lateral cracks is about the same

as that of the scratching particle. Quite often there is a correlation in the length of

neighbouring lateral cracks: the second one is shorter than the first and the third smaller

again (fig. 15b).

50
2. Underneath the surface a vertical crack starts out just underneath the top of the pyramid

[9] (fig- 16). This "median" crack occurs especially when the pyramid is used in the leading

edge position [ 9] • A minimum load seems necessary before the extent of the stress field is

large enough to nucleate this type of crack. In Peter's extensive work ([8] , see also Dick

[ 27] ) the extent of cracking is found to depend on the load and on the speed of scratching.

3. On the surface, especially at higher loads, and at lower speeds two lateral cracks are

seen to combine to form chips (fig. 15c). The second crack thereby takes a curved path and

ends up about perpendicularly to the first crack [ 9] • For a theoretical treatment the com- plex stress field seems forbidding. The chips are particularly interesting for the mechanism of material removal. For one scratch the chips often remain connected to the surface

and a second scratch seems necessary to harvest the loose material. At the end of section

k.2. we shall return to this question. In the microscope the chips are seen to be formed

behind the scratching point.

k. Subsurface cracks , propagating parallel to the surface (fig. 16), are easily seen on the

glass surface from their optical effects (fig. 15b). These cracks are formed behind the

scratching particle (as observed in the film by Busch [26] ) and are similar to the cracks

that start on the unloading of a hardness indenter [ 12] .

The cracks described start near the scratching particle, i.e. in the complex stress

field of a pyramid. The chips and subsurface cracks are formed behind the point where the

residual stress due to the groove adds up to the field from the point itself. To some extent

similar complexities occur for the median and the lateral cracks, as long as these are

short. To add to the difficulties of an analytical treatment, the cracks are formed in a material that already contains a number of cracks. This may be responsible for the 1-2-3

effect in the lengths of subsequent lateral cracks (fig. 15b). It may also be relevant

for the formation of a groove, i.e. the specific energy may be lower in a cracked material,

because the effective hardness will be somewhat lower due to a lower effective stiffness.

A similar effect is seen in the widening of a groove when it reaches the end of a sample.

The analytical treatment of the extent of cracking for a given load on an indenting

particle was first given by Roesler [ 28] for the cone crack and by Lawn [ 11] and Swain [ 12]

for the median and penny-shaped crack. From the known stress field it is possible to deduce

when the stress intensity factor Kj reaches the critical value for crack propagation Kjc

. At a large distance from the loaded point the derivation is simple: for a load Fn the

stress at a distance r is O = F

n

/irr^. If r is also the crack length, the Kj value is

KT = oVr, which immediately leads to

r = (Fn

/7TKI

) 2/3

(7)

For small cracks, however, the presence of the groove cannot be disregarded. Since the

details of the stress field are unknown, one can use the experimental results on lateral

cracking to find a correction term. Veldkamp [ 9] found a satisfactory description of the

experiments with the relation

= a K-j.r(r+6b (8)

where a is a constant and b is the width of the groove. The Kj values deduced from (7) and

(8) agree reasonably well with values found from the double torsion test [9] for the propagation of slow cracks. The effects are illustrated in figure 17, where the decrease of r with increasing speed is shown.

t2r 40

30

20

Fig. 17. Crack length taken on both side of a scratch, (2r) as a function of scratching speed

(v ) for MgO.

04 40 400

51
1Cr

3 M

(m/s)

10-'

10"

10"'

. double torsion scratching

05

6 3 Kj(10 N.rrfT)

0.5

3 10

K,(10 N.m"7) —

05

6 3 K.1 10 Nm"7)

Fig. 18. Speed dependence of KT as determined from scratching experiments on glass (left),

MnZn ferrite (middle) and a MgO crystal (right). Kj values from a double torsion experiment

have been added for the glass [ 9] •

By using eq_.(8) the results can be presented as a vs -K diagram, as so often is used in

stress-corrosion cracking of glass [ 1-h] (fig. 18). The appearance of a given type of crack

is also a function of scratching speed, and a diagram can be made that combines the effects

of speed and load (fig. 19). Grooves are observed at all loads and speeds, but the various

cracks are found in a limited region only.

polycrystalline materials

g»c Urn)

g*l Urn)

F n

glass

Fig. 19 Regions in a load (Fn )-speed (v ) diagram where the various types of cracks occur,

for a polycrystalline MnZn ferrite (lefl) and a glass (right). Grooves (g), lateral cracks

(l), median cracks (m), subsurface cracks (s) and chips (c) are found in these materials.

h.2. Crystalline materials

Anisotropic materials often show the same types of crack as glass but the direction

of cracking may be quite different for materials with pronounced cleavage planes. An

example is the surface of a scratched MgO crystal (fig. 20) where the planes of easy

cleavage show up clearly. In the cross-section shown in figure 10, horizontal and vertical

cracks were visible; the horizontal crack extends over the length of the scratch [29] . The case of cracking makes it difficult to determine the groove dimensions. The extent

of cracking in MgO shows a speed dependence (fig. 17)

•

Fig. 20. View from above of scratch made on

(010) plane in the [100] direction of MgO

single crystal, load 0.1 N, scratching speed

40 ym/s. Incident light.

52
Median cracks have also teen observed in a MnZn ferrite, scratched by a pyramid with

a leading edge (fig. 21). The surface cracks are of the lateral type, accompanied by chips

(figs. 3 and 22). Here too, speed effects have been found [9,10] and the types of cracks

can be given in a load-speed diagram, as was done for glass (fig. 19), but now the subsurface

cracks are missing.

Fig. 21. Cross-section of a scratch in Mn Zn

ferrite. Clearly shown is the median crack.

Load 1 N, speed kOO um/s. Incident light,

bright field.

Fig. 22. Scratch made on MnZn ferrite. The

two types of chip are marked with I and II.

Load 0.2 N, scratching speed k um/s. Scanning electron microscope.

In a polycrystalline material it can easily be shown by scratching where cracking takes

place, along the grain boundaries as in fig. 23 or through the grains (fig. 2k). The

example in fig. 2k has exceptionally strong grain boundaries; one chip is seen to contain

even three boundaries. As in glass, the large chips have been observed to crack open after

the passage of the scratching pyramid.

Fig. 23. Scratch on NiZn ferrite, fired

at 1230°C; load 1 Nj width of picture:

150 urn.

Fig. 2k. Scratch on NiZn ferrite, after hot

isostatic pressing at 1300°C; load 0.5 N;

scratched after thermal etching; width of

picture: 150 urn.

Notice that the chips remain fastened to the surface. A greater removal of material

is expected to occur when a second scratch interferes with the chips of the first one. An example of this effect has been found on a MnZn ferrite that shows large chips. The

following experiment was done: scratches were made at various distances (fig. 25) and the

53
material that was removed was determined by weighing. A pronounced increase was found below

a certain distance that still depends on load (fig. 26).

Fig. 25. Photographs of parallel scratches on MnZn ferrite at various distances between the

scratches: 250, 175, 200 and 125 ym.

Another item of interest in the context of material removal is the number of lateral

cracks and chips per unit length of the scratch. Both increase with increasing load; the

chips remain smaller until at higher loads the lateral cracks have disappeared in the area with large damage and a few large chips remain, as shown schematically in figure 27.

Fig. 26. The effect of the distance between Fig. 27. The number of lateral cracks per

parallel scratches, at constant load, on the unit scratch length, as measured on a MnZn

weight of the material that was removed. ferrite, and, schematically, the number of

Weight per unit scratch length vertical, chips, both as a function of load

interscratch distance horizontal. Material (vs = kOO ym/s).

Mn Zn ferrite (20 = 136° vs = 50 ym/s).

5. Liquids

In ferrites the effects of water are similar to those in glass. Scratching with a pyramid in a droplet of water the results of figure 28 were obtained: a slow scratch in

normal, humid, air has a cracked groove, whereas a faster one is smooth. The influence of

water on the strength of ferrites has been noticed before by Lessoff [ 30] . Liquids on the surface of an oxide may have a softening effect as well as an influence

on crack propagation. We used a variety of liquids during the scratching of MgO crystals,

such as n-alcohols, di-methyl-formamide (DMF), di-methyl-sulfoxide (DMS0) and a k0-60%

mixture of the latter two. Neither force ratio nor the scratching hardness were found to

be influenced. Underneath the surface, however, the dislocation pattern extends farther

for hexanol than it does for methanol or ethanol, in agreement with Westwood's findings

54
[ 1U] , taken by indentation at much longer times than in scratching. Two other experiments

also gave negative results for a possible influence on the force ratio: one on quartz

crystals in the presence of dodecyl-trimethyl-ammonium-bromide (DTAB) and one on a barium

titanate in DMF and DMSO. The latter combination of a piezoelectric material with a liquid

having a high dielectric constant looked promising, but gave no result. A real lubricant,

DDDS - dioctadecyl-disulfide-, was used on a glass ceramic, dissolved in di-ethyl-ether

and in ethanol; neither in solution nor after the evaporation of the solvent was there any

change in the force ratio (fig. 29). The lubricating action during grinding, however, is

in fact present, because when the solution is sprayed on a grinding wheel, the grinding

force drops to a low (friction) level and material removal stops!

The lubricating effects of alcohols etc. on oxides have been investigated by Wiederhorn

[ 31] , using a spherical indenter; a sphere allows a more direct determination of the friction coefficient (cf. eq. (2)). Again no effect was found. The only positive result is the

softening of MgO in hexanol, but only after the samples had been annealed. A possible

influence of alcohols on the ease of crack propagation was investigated by Freiman [3,4

Fig. 28. Scratches on SrFe-|20l9 under various circumstances: slow (vs = 50 um/s) and dry

(left); fast (750 um/s) and dry (middle); slow and wet (right). Load 5 N.

A slightly higher crack speed is only found for hexanol.

0.6

05

0.4

0.3

„ 02

01

Glass ceramics 29 = 100"

Vs = 200 >jm/s

Fig. 29. Force ratio f versus load for

scratching on a glass ceramic with and

without DDDS (see text).

0.1 02 03 0.L 05

Normal load FN (Nl

55
100

150

Fixed in upper plate

Loose sapphire spere

Vp= 8 >jm s"'

Fn = 1.0 0 5 0 2 0.1 005 N

Fixed in lower plate

0 50 100 Vickers hardness Hv I10

7 Nm'2

) »

Fig. 30. Scratch length on lower metal divided by total scratch length, the rolling

fraction, as a function of the hardness of

the lower metal for various loads

;

v

g = 8 Um/s; brass upper plate; sapphire

sphere

.

6 . Rolling

Scratches were made with a loose particle between a brass plate on top and a variety

of materials underneath. For a sapphire sphere (0.1 or 0.2 mm diam. ) some of the results

are simple. On soft metals (in, Sn, Pb, Al, Zn, Cd, Cu) a transition occurs between fixation in the soft metal and rolling. This transition depends on load (fig. 30 )

, the harder

substrates needing less load for rolling to occur. On harder metals, like annealed and

quenched steels or soft cast iron, rolling and fixation in the upper or the lower metal

plate occurred in one scratch. The effect of increasing scratching speed is also confusing

in that fixation and rolling seem to happen irrespective of the circumstances. At high

speeds, however, rolling seemed to be the dominant phenomenon.

For a sharp diamond grain of about 250 urn, the degrees of freedom include that of

rotation around several axes, as can be seen in the pictures taken successively during

a scratch (fig. 31). The trace left on the workpiece surface gives the impression of a periodic but random walk (fig. 32). We conclude that, although scratching with a fixed

pyramid seems to present a number of problems, a loose particle is much worse. The inefficiency of lapping with loose particles is clearly visible in these single-point loose-scrat

ching experiments.

Fig. 31. Subsequent photographs of rolling Fig. 32. Track on SrFe-^O-io. after a rolling

diamond (250 ym diam.) and its two images in diamond has passed over its surface. Load

the upper and lower plate (brass). During the 0.2 N; upper plate: brass,

motion of the lower plate the particle ro- tates around a horizontal and the vertical

axis. Load 0.5 N; v., = 8 um/s.

56
6. Discussion

Scratching experiments are a source of information on three types of problems: what

happens during scratching, what can be learnedabout the material and in how far can grinding behaviour be predicted?

1. The problems of single-point scratching can be summarized as follows:

a. Can one ever expect a detailed description of the stress field around the point and of

the deformation? Electron-microscopy (Hockey, Lawn [32]) provides an enormous amount of

interesting information, but it is qualitative in nature with respect to a description

of the deformation.

b. From the groove dimensions under a known total force, characteristic average stresses

have been derived (e and Hs) which are related to the indenter hardness. The hardness values

show a dependence on depth, the "size effect". It is especially large in AI2O3 and rather

small in ZnO (fig. 13). The same result was found for the soft polycrystalline ZnO

(Hv

= 2.109 N/m2 ) as for the harder ZnO crystal (H^ = 10.109 N/m2 ).

We consider two possible explanations for the size effect: one is based on a reduced

stiffness, the other uses the speed dependence of the hardness. The first argument notices

an important difference between the extreme cases of ZnO and A^O-^: the number of cracks in

ZnO is much smaller. A cracked material has a larger compliance; similarly, a scratching particle at the end of a sample encounters a more compliant material. A lower

resistance to bending may lead to wider grooves and a lower hardness. This argument might

also be used to understand the variation in length of consecutive cracks along the scratch

(fig. 15b).

A quite different explanation of the size effect relies on the speed dependence of

the hardness (fig. 1M, relating it to an effect of the strain rate, e = vs /d and giving

equal weight to a change in speed vs as to depth d. This model could be checked by careful

measurements of the size effect at various scratching speeds.

c. A third problem is the sign of the interface friction coefficient y in eq. (2). The sign

depends on the material [15, 17] for scratching with a leading plane, and it is rather

large for the case of a leading edge. The difficulty is to predict the relative motion

of the pyramid and the material that is emerging from the upper layer of the workpiece.

d. Cracks are formed on the surface and underneath the scratching point. Some cracks

(median, lateral) are formed in the contact area between indenter and material, others

are formed( subsurface , chips) after the point has passed. All cracks propagate in the

combined stress field of the loaded point and the groove that was just made. Since the

groove region is highly deformed this is a complicated problem. Correct analytical ex- pressions are not expected.

2. Scratching provides an easy way to obtain the following information about a material.

a. The scratch hardness can be determined, similar to the indenter hardness; it can be

used as a test of homogeneity as well.

b. The extent of cracking is an indication of the fracture resistance of the material

and by using eqs. (7,8), Kjq can be derived. In a polycrystalline material the strength

of the grain boundaries can be tested qualitatively; the same holds for the possible

weakening effects of second phases and pores.

c. Speed-dependent fracture phenomena can be studied on a microscopic scale by direct

observation. The influence of special environments is easily studied. The problem, however

,

is that scratching is no more than a handy mechanical device, but cannot provide details

of the mechanisms involved.

d. A minimum load Fmj_ n can be derived from eq. (8) below which the cracks are confined

to the groove [ 9] • Fm.; n is proportional to KVH~ and may be used in considerations about

lapping, for example.

3. Scratching can be used as a basic model for grinding and lapping by comparing equivalent

quantities such as the specific energy and the force ratio [5 S 6,22] . For some materials

the agreement is quite good, for others the extent of cracking must be considered, because

57
that was omitted in calculating the groove energy from eq_. (k) . The experiments with parallel scratches at various distances have shown that cracking plays an important role at the

higher loads (fig. 26). Damage made by grinding can he studied in detail "by scratching,

where especially the depth of cracking is relevant.

7. Conclusions

Scratching has proved to he a useful basis for understanding qualitatively the deformations that are encountered in grinding and other abrasive processes. The experiment

allows detailed investigations of the dependence on load, speed and environment for those

quantities that characterize a scratch. The results are as follows.

1. The force ratio is independent of load, material, speed and environment; it is determined by the geometry of the point. The interface Coulomb friction is a minor correction

for pyramidal points, whereas it is the major effect for a sphere.

2. The specific energy depends on the material and is proportional to the Vickers hardness.

No dependence on environment was found, but for several materials its value depends on the

load, the so-called "size effect", and on the speed, at least in the present range of (low)

scratching speeds.

3. The type of cracking depends heavily on the type of material. In glass^ four types have

been found; in crystalline materials, three . The extent of cracking depends on load according

to eqs.(T and 8), which contain the material-dependent fracture toughness Kj. Especially

in humid environments containing water the extent of cracking in several materials becomes

speed-dependent

.

The link with grinding parameters, such as the force ratio, the specific energy, the

surface roughness and the depth of damage, can now be made.

References

[ 1] Charles, R.J. , Static fatigue of glass, J.Appl.Phys. 29, 15^9-1560 ( 1958).

[2] Davidge, R.W. and Tappin, G. , Strength of polycrystalline aluminas, Proc . Brit. Cer.

Soc, _1_5, ^7-60 (Jan. 1970).

[3] Freiman, S.W. , Effect of environment on fracture of ceramics, 2, [3] , 111—118 (1976).

[ k] Freiman, S.W. , Mulville, D.R. and Mast, P.W. , Crack propagation studies in brittle

materials, J. Mater. Sci. 8, 1527-1533 (1973).

[5] Gielisse, P.J., Kim, T.J., Goyette , L.F. and Nagarkar, R.V. , Fundamental and practical

evaluation of the ceramic finishing process, Report (Univ. of Rhode Island), Kingston

R.I., 1975.

[6] Prins, J.F., A study of single diamond particle interaction on alumina, Ind. Diamond

Rev., 1+97-503 (Dec. 1971).

[7] Busch, D.M. and Prins J.F., A basic study of the diamond grinding of alumina, (Proc.

Symp. Science of Ceramic Machining and Surface Finishing, Gaithersburg, Md. , November

2-h, 1970), pp. 73-87 (NBS Spec. Publ. 3I+8 )

.

[8] Peter K. , Sprodbruch und Mikroplastizitat von Glas in Eindruckversuchen, Glastechn.

Ber., 37, [7] , 333-31+5 (July 196U).

[9] Veldkamp J.D.B., Hattu N. , and Snijders, V.A.C., Crack formation during scratching

of brittle materials, in Fracture Mechanics of Ceramics , Vol. 3, R.C Bradt , D.P.H.

Hasselman and F.F. Lange, eds., pp. 273-301 (Plenum Publ. Corp. 1978).

[10] Veldkamp, J.D.B., Crack formation during scratching, (Proc. 9th. Int. Conf. on Science

of Ceramics, Noordwi jkerhout , November 13-16, 1977), Book, Science of Ceramics , 9., K.J. de Vries ed. , pp. 1+38-1+1+6 (Nederlandse Keramische Vereniging, 1977).

[ 11] Lawn, B.R. and Fuller, E.R. , Equilibrium penny-like cracks in indentation fracture,

J. Mater. Sci. JO., 2016-2021+ (1975).

[ 12] Swain M.V. and Hagan J.T. , Indentation plasticity and the ensuing fracture of glass,

J. Phys. D.: Appl. Phys . , 9, 2201-2211+ (1976).

[13] Westbrook, J.H. and Jorgensen, P. J. , Effects of water desorption on indentation microhardness anisotropy in minerals, The Am. Mineral., 53, 1899-1909 (Nov. 1 968 )

.

[ 11+] Westwood, A. R.C, Control and application of environment-sensitive fracture processes,

J. Mater. Sci. 9, 1871-1895 097*+).

58
[ 15] Broese van Groenou A., Maan N., and Veldkamp, J.D.B., Scratching experiments on various ceramic materials, Philips Res. Repts., 30_, 320-359 (1975).

[ 16] Stokes, R.J., Johnston, T.L. and Li, C.H. , Crack formation in magnesium oxide crystals,

Phil. Mag., 3, 718-725 (1958).

[ 17] Maan, N. and Broese van Groenou, A., Low speed scratch experiments on steels, Wear,

h2, 365-390 (1977).

[ 18] Hisakado, T. , On the mechanism of contact "between solid surfaces, Bull. JSME, J_3, [55] , 129-139 (1970).

[ 19] Bowden, F.P. , and Tabor, D. , The friction and lubrication of solids (Clarendon Press,

Oxford, 195M . [20] Bowden, F.P. and Tabor, D., Friction, lubrication and wear: a survey of work during

the last decade, Brit. J. Appl. Phys., J_7, 1521-15^ ( 1966).

[21] Shaw, M.C., Fundamentals of grinding, in New Developments in Grinding , M.C. Shaw,

ed. , pp: 220-258 (Carnegie Press, Pittsburgh, Pa, 1972).

[22] Broese van Groenou, A., Veldkamp J.D.B. and Snip, D. , Scratching and grinding parameters of various ferrites , J. de Physique, C1, 285-289 (1977).

[23] Broese van Groenou, A., Experiments on grinding of ceramic ferrites, see ref. 9,

pp. U86-U93.

[ 2k] B.G. Koepke and R.J. Stokes, Grinding forces and the machining of magnesium oxide

crystals, J. Mater. Sci. 7, U85-U93 (1972).

[25] Lawn, B.R., Partial cone crack formation in a brittle material loaded with a sliding

spherical indenter, Proc. Roy. Soc . (London) 299

, 307-316 (1966).

[26] Busch, D.M. , Verformen von Glas mit Ritzwerkzeugen, Films 1328-1330, Inst. Wiss.

Film., GSttingen 1967.

[27] Dick, E., Plastische Verformung, Anrissbildung und statische Ermudung bei Ritzversuchen

an Tafelglas, Glastechn. Ber. , U6 [ 6] , 120-128 ( 1973).

[28] Roessler, F.C., Brittle fractures near equilibrium, Proc. Phys. Soc. B69 , 981-992 (1956).

[29] Sugita, T., Suzuki, K. and Kinoshita, S. , Wear characteristics of magnesium oxide single

crystals on steel, Wear U5, 57-73 (1977).

[30] Lessoff, H. , Effect of moisture on strength of ferrite memory cores, J. Am. Cer. Soc,

5^, [7] , 353 ( 1971 ).

[31] Widerhorn, S.M. and Roberts, D.E., Influence of normal alcohols on the abrasive wear

of glass, Wear 32, 51-72 (1975).

[32] Hockey, B.C., and Lawn, B.R. , Electron microscopy of microcracking about indentations

in aluminium oxide and silicon carbide, J. Mater. Sci. _10_, 1275-128U ( 1975).

Discussion

BENNETT

Did you investigate scratch behavior for increasingly light loading?

Your slide on the force ratio seemed to indicate the beginning of a deviation from a constant value for light loading. Our results on much softer

materials suggest that this deviation would continue to increase as the

loading decreases and is not simply a result of the error in shape of the

diamond point

VAN GROENOU

Our experience is the following: every time we found such a decrease in

force ratio on decreasing the load, we looked at the diamond under high

magnification (dark-field) and found something suspicious at the point,

usually a tiny flat spot. After regrinding the pyramid the force ratio was

found to be increased.

CHOUDRY

The tangential force (Ft) for high speed grinding in your model is strictly governed by the geometrical volume removed and this volume is calculated

by the geometry of the tool (a) , depth of cut (d) and the ratio of the work59
piece speed (Vw ) & wheel speed (V s ); all of which are experimentally adjusted

parameters and do not depend on the elastic (or materials) property of the

workpiece . Does it mean that for a given set of a, d, Vw , V„ , one would always get the same F^ irrespective of the workpiece material be it diamond or

marble. On the other hand the dynamic elasticity of the impact would inescapably generate a force which among other variables depends on the materials properties. Is there a need in your model to take this into account?

VAN GROENOU

In the paper the specific energy is introduced by the relation power =

e. Z, also power = -Vs from which follows F^- = e Z/Vs „ The material properties relevant for grinding are contained in the factor e.

60
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

EXPERIMENTS ON HIGH-SPEED GRINDING OF FERRITES AND GLASSES

A.Broese van Groenou and R.Brehm

Philips Research Laboratories

56OO MD Eindhoven, The Netherlands

At wheel speeds of 60-100 m/s measurements have been made on a horizontalspindle grinder with a nickel-plated diamond wheel. The grinding force components

are unique functions of the continuum chip thickness he ^

= avw

/vs , over a large

range of the depth of cut a, the workpiece speed vw and the given range of wheelspeed vs

. The vs dependence is different for those glasses where melting occurs. Where a specific energy can be determined, it shows a dependence on the depth of

diamond penetration. This "size effect" is strong in the case of low- toughness

materials. The energy data are compared with those obtained for grooving in lowspeed scratching experiments. Deviations occur where groove formation is accom- panied by heavy chipping. Finally, for brittle materials, the requirements of a grinding model are considered, with emphasis on the amount of chipping and the

diamond distribution on the grinding wheel. In this model the specific energy

is a characteristic of the grinding system and not of the workpiece alone.

Keywords: Chipping; ferrites; glass melting; grinding; grinding model; high-speed

grinding; pyrex; size effect.

1 . Introduction

The first meeting on "Ceramic Machining and Surface Finishing" [ 1] ' showed clearly

that grinding is a complicated process, in which many components take part. Macroscopically

one finds: motor, grinding wheel, frame, workpiece, cooling. On a microscopic scale : diamonds

,

wheel and workpiece microstructure , frame displacements and wheel unroundness. This complexity stimulated us to set up experiments and collect data in order to find a description of

the process on both a large and a small scale. Medium hard materials, ferrites and glasses,

have been studied in some detail. The results will be reviewed in this paper; some of the

results have been published elsewhere [ 2-h] .

A macroscopic quantitative description considers the components F^ and Fn

, rate of vo- lume removal, Z, the power W, the force components F-^ and Fn , the workpiece-wheel contact

width b, the depth of cut a, the workpiece feed vw , the wheel speed vs and the contact

length lc (fig. 1). For horizontal surface grinding the following relations are useful [5] :

z = a.b.v

w

(1)

w (2)

= e.Z (3)

z = F .v /e

t s' (k)

= F .v /H

n s' g

(5)

f = F./F

t' n

(6)

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

61
DIAMOND WHEEL

Fig. 1. Schematic view of grinding wheel in

horizontal surface grinding.

where, by analogy with the grinding of metals [ 5] , the force ratio f, the specific energy

e and the grinding hardness H were introduced. Both have the dimensions of a stress, or of

an energy per unit volume. Together with the wheel wear and the surface damage, e.g. expressed in terms of c.l.a. roughness, these quantities usually characterize a grinding operation.

In grinding steel a useful quantity is the continuum chip thickness, hg^, defined by [6,7]

Z = b.v.h n (T)

s eq

which, together with eqs. (U) and (5), gives

F , = b . e . h (8)

t eq

F = b H .h (o) n g eq Ky 1

It would be a considerable simplification if these relations were to hold for grinding ceramics and glass. If that were the case, different grinding machines could be compared, although used at different values of a, v and v . In grinding steel [5] the linear relation

between force and rate of removal that is suggested by eqs. (8) and (9) is not found. The

deviations are usually described by having e and H vary with, for example, h

e(^. Such a "size effect" has also been found in grinding MgO f8] and forms an important part of the

discussion on the present results for several ferrites (sections 3.1 and 3.2) and glass

(section 3.3). From eqs. (k) and (5) the advantages of using high-wheel speeds are obvious,

as long as the specific energy is independent of v

s

. A comparison between e values found in

grinding steel at 20-30 m/s [5] and low-speed scratching at 0.U-U00 mm/s [9] shows that the

high-speed e values are higher and show a stronger size effect. The experiments to be described were done at 60 to 100 m/s (section 2) in order to see whether such discrepancies

are also found in oxides; where possible the results obtained at 5 m/s will be included { 2] . The corresponding scratching experiments have been described in refs, [ 10,11] . Some highspeed grinding experiments were done on a number of glasses, one of which, a borosilicate

glass, will be treated here because of its deviating behaviour.

The reason for comparing grinding data with scratching becomes clear when the processes

are compared on a microcopic scale. At the beginning of contact between a grinding wheel

and a slightly tilted sample (e.g. 1:100) one observes first a fine scratch at the start of

the tracks made by the diamonds; then, the track widens and chips appear, which finally

merge with the deformation made by neighbouring particles (fig. 2).

From the figure another complication is clear: different diamonds touch the tilted sample at

various places. This makes it impossible to speak about the depth of cut for every particle

and statistical considerations are necessary for any grinding model [ 12,13] . Figure 2 also

shows the two basic features of the grinding models that have been proposed, grooves and

cracks. The grooves have led to the ploughing type of models [5] , which give the resistance

to grinding or lapping as proportional to the hardness of the workpiece, in agreement with

experiments on lapping of glass [ 1 U , 15] and ceramics [ 16] .

62
Fig. 2. First contact between grinding wheel and slightly tilted sample of MnZn ferrite

(v

s = 75 m/s).

The cracks and chips have been used in models that are based on the fracture underneath

a statically loaded spherical indenter, where ring and cone cracks are formed. These ideas

have been applied to erosion [ 17] and lapping [ 18] . A difficulty is estimating the volume

removed by a single indentation, and opinions vary considerably [ 12, 17, 18] ; it seems essential to take into account the interaction between cracks made by different indenters because

no material is loosened from the surface by a single indentation [ 18] . One result of this

model is a dependence on the fracture strength, which agrees with Gielisse's results on grinding various types of alumina [ 19] • Figure 2, however, does not show separate indentations

but long scratches. A grinding model based on the phenomena around a scratch has been given

by Veldkamp [ 2] and we will return to such models in the discussion. First a description

is given of the experimental set-up; results for two ferrites and one glass are given in

section 3. The relevance of the continuum chip thickness is discussed in section h, followed by considerations about grinding temperatures, the size effect and the requirements

for a grinding model for brittle materials

2. Experimental

2

The measurements were made on a flat bed grinder using a horizontal axis motor , running at

20000 to 36000 r.p.m. , with a nickel-plated diamond wheel (diamond size 220 y.m) . The horizontal traversing carriage is hydrostatically supported and hydraulically driven. The. oil

unit is separate from the grinder (fig. 3). The following values were used: the feed vw

from

0.1 to 60 mm/s, the depth of cut a up to 500 urn, the wheel speed v from 58 to 102 m/s

(in continuous operation, lower values were obtained upon starting the motor), the sample

length 10 to Uo mm and the contact width b from 2 to 5 mm. Mist cooling was applied. The

apparatus and a detail are seen in figure 3. The force components were measured by means of

a piezoelectric transducer and the necessary electronics, including a multichannel tape

recorder^ for the data on force components, speed v

g and radial vibrations^.

The apparatus was carefully balanced. The amplitude of the vibrations of various parts of

the grinder was measured as a function of frequency, both by external excitation and with

the motor running. The amplitude u during grinding remained constant at 330 Hz (vs = 58 m/s),

increased at 1+95 Hz (86 m/s) and decreased at 550 Hz (96 m/s) with respect to the amplitude

before contact (fig. 4).

SJF-HE 1+0, Fischer, Herzogenbuchsee , FRG

Techno Diamant , Amsterdam

Kistler 9259, Winterthur, Switzerland

Analog 7 j Philips

PM 61+31 , Philips

63
Fig. 3. Grinding apparatus: electronics (left, grinder and oil-and mist supply right. Detail

showing wheel, motor and the contactless vibration transducer above the extended shaft

Frequency 330 Hz 1+95 Hz 550 Hz v

s 58 m/s 86 m/s 96 m/s

h 0. 17 ym 0.09 um 0.28 ym eq

Fig. k. Amplitude of axial vibration u and normal force Fn

as a function of time, showing

the changes upon the start of grinding.

From such results we have the impression that vibrations do not contribute much to the

force level during grinding. More important is the finding (also on other grinders we have

used) that the force level is pulsed, with a repetition time equal to the period of the

wheel (fig. 5). This was due to the unroundness of the wheel. During about 18 months of

experimenting the wheel showed very little wear, as measured by making replicas of those

parts that had a slightly larger radius. This was known from a roughness measurement, done

by rolling a small cylinder over the wheel surface. From the roughness plot the distribution

of the heights of the diamonds was obtained.

Fig. 5. Force pulses during several revolutions

of the grinding wheel; frequency 330 Hz, v = 58 m/s, vw

= 60 mm/s, h = 0.84 ym; 20 N/div

vert, 1 ms/div horizon; material: SrFe-^O^g.

64
The following samples were used:

a. SrFe-ipOip, a hexagonal ferrite, polycrystalline, porosity 8%, aligned, grain size about

5 ym,

b. (Mn, Zn, Fe) Fe20|

+

, a cubic ferrite, polycrystalline, porosity Q%, grain size about 15 ym,

c. A borosilicate glass with an ASTM softening point of 800°C

d. Quartz glass.

3. Results

3.1 SrFe

l2

0

19

The results for three wheel speeds are given in figure 6 for the normal force component

Fn

as a function of the volume rate of removal Z; both quantities have been reduced to unit

width of contact. The whole range of variation for a and vw

were used. The three v

g ranges

show clearly different curves. From eqs. (7 and 9) a single curve is expected if he ^

= Z/bvs is plotted along the abscissa. For both F^ and F

n this is indeed the case (fig. 7); there

is no longer any systematic variation with vs

. This result means that one combination of the

machine parameters, a, vw

en vs

, is sufficient to describe the grinding force for this material. The same holds true for the (c.l.a.) roughness, which shows a relatively small variation, from 0.7 to 1.5 ym, over the h

g

range in fig. 7- The force ratio, f = F^/F

n

, varies

range

from 0.1 to 0.3 over the same heq

m/s

| 30 i

10 20 3p «j 50

° 76

• 81 m/s

a 94

» 102 m/s

20 30 40 50

-»• Z' (mm 3 /sec mm)

£

e

xo

5

z

OCL

O

o

o

z

Q

z

oco

heq (urn)

Fig. 6. Normal component of grinding force Fig. 7. Grinding force components, F and Ft , versus continuum chip thickness he(^; data

from fig. 6.

Fn

as a function of rate of removal Z, both

per unit width of contact zone. Three ranges

of wheel speed v

g are shown; material:

SrFe

12

0

l9

.

In these results the peak values of the force components were used, whereas Z from eq.

(1) and heg from eq. (7) are average values. For the calculation of the specific energy

and the grinding hardness from eqs. (8) and (9) a factor of three was used to change from

peak to average values. The results are shown in figure 8, where data from ref. [2] taken

at vs = 5 m/s are shown too. The slope in the log F - log h plot of fig. 7 was less than

one for F

n and just above one for Ft , which corresponds in fig. 8 to a decrease of H and a slight increase of e with h . ?

eq

65
100

o

-5 W0

X

10

GRINDING. PLATED WHEELS

a

Vs = 58-102 m/s

V-7-V.V—- NiZn

0,1 » hJ^(um1/2)

eq

100

50

I 20

5: 10

l_c

01

u

in

o 25

Vw = mm/s

• 50

Vs =58m/s

>

• •

10-' 5 10'

1 2

Fig. 8. Specific energy e and grinding hard- Fig. 9- Specific energy for MnZn ferrite

ness Hg versus h

e ^ for SrFe-^Oio at vg = 5 m/s at 58 m/s, versus h

g

. ([2,3] and at high vg values.

e<^

3.2. (Mn.Zn.FejFe^

The data on this cubic ferrite show the same independence of wheel speed when the

grinding force is plotted versus h

g as for the hexagonal ferrite. The force ratio is about

0.2 over a large range of h valuei. The specific energy data show a marked decrease at hgi

larger than 0.1 urn (fig. 9). we shall return to this size effect in section k.3.

A qualitative check was made on the behaviour of some materials as a function of wheel

speed. Again using the first contact with the diamonds on a tilted sample, of the same MnZn

ferrite, the photographs of fig. 10 were obtained, giving the grooves due to the same group

of diamonds. There seems to be more variation due to local differences in the material than

due to the fivefold increase in vs

.

Fig. 10 Start of grooves on MnZn ferrite, as observed on a tilted sample. From left to right vs = 22, 58, 81 and 10U m/s.

66
When the grooves become deeper, more chips appear along the scratches, which finally merge

with those from neighbouring scratches. We have studied this interaction between parallel

scratches at low speed in a single-point apparatus by varying the distance in between.

The smaller the distance, the more damage is found and the weight of the material removed is a strong function of distance [11].

3.3. Glasses

Quartz glass behaves in the same way as the ferrites at these high speeds. The grinding

hardness decreases from kO to 30J09/Nm2

, the specific energy from 15 to 7.1C>9j/m3 when he ^ increases from 0.1 to 0.6 urn. The data for vs = 58 and 9*+ m/s are equal within the scatter.

The e values are about a factor 3-5 larger than Huerta and Malkin's results at v = 19 m/s

[ 20] .

The borosilicate glass, however, deviates in several respects when v„ = 9^ ni/s is used.

• • • V The time dependence of the force during the passage across the workpiece is an example.

Normally the force increases upon contact and quickly reaches a constant value; at low he ^ this is indeed found (fig. 11a). At higher h it takes much longer to reach a constant value

(fig. 11b). The damage changes too, from a normal, rough, surface to a shining surface with

dikes of molten material, cracking and sometimes large holes where big lumps have disappeared

(fig. 12). In spite of this, the F

n-h plot is quite normal (fig. 13a). The tangential force,

however, is completely irregular at high v

g

; at all h values above 0.2 urn melting occurs and the correlation between F-fc and he(^ is lost (fig. 13b) For a molten surface the continuum

chip thickness is no longer a useful parameter to describe grinding; the same holds for the

specific energy. At lower wheel speeds, up to ho m/s, Huerta and Malkin [20] found a normal

F - Z behaviour, as in our results at 58 m/s in fig. 13.

k. Discussion

U.1. Significance and use of h

eq

For the results on the grinding of ferrite, shown above, and a number of other materials, the continuum chip thickness is a valuable parameter. It can be calculated from the

machine settings with eq. (7) and is a direct measure of the grinding forces for a given

material, i.e. for given e and Hg values (eqs. 8 and 9). Measurements on dry grinding of

MgO reported by Koepke [ 8] can be presented in a simpler way by using h

e

. It is also a measure of the thickness of the debris layer that would be left on the grinding wheel if no cleaning were used; in fact, this is where h got its name of continuum chip thickness

[6,7] • In this application he ^ should be compared with the height of the abrasive particles

above the binder level. If the wheel is clogged after a small number of revolutions, each

of which add a layer of h

g

, the cleaning action should be improved. For small diamonds or

a large contact area this is sometimes solved by segmenting the grinding wheel, e.g. by

using pellets. One should distinguish between hg^ and the average depth of cut of the diamonds, which will be discussed in the next section.

The surface roughness is only a weak function of he ^, much weaker than the force.

Since grinding force and workpiece damage are certainly related, this result is surprising,

at first sight. Two reasons, however, may be given. Firstly, as mentioned before, he q is not

a direct measure of the depth of cut of the average diamond; it can even be shown that for

a distribution of heights of the diamonds, its value depends only weakly on the depth of

cut of the outermost diamond. Secondly, referring to figure 1, the maximum depth of cut for

each diamonds occurs at the end of the contact length (at the beginning for down-grinding);

this area, however, is removed to a depth a by the action of the following diamonds. This

probably leaves the damage at first contact (last contact for down-grinding) as the main

contribution to the roughness. This part of the contact zone, however, hardly contributes

to Z or he ^ and certainly not to the peak value of the grinding force. It is therefore not

surprising that roughness and h (or F ) are not strongly correlated.

67
: t J

F

n 1

he q=0 . 1 5um

hi

n I

heq =0.28um

Fig. 11. Borosilicate glass (v = 9^ m/s).

Time dependence of the force components F. and F .

Examples of surfaces for h = 0.15 ym (a),

normal) and h = 0.28 ym f§) , melting.

Fig. 12. Surface of ground borosilicate glass (v = 94 m/s). Left: where the grinding was interrupted surface melting stops (left) and a normal surface from a previous run remains

(right). Width of picture: 1 mm. Right: molten chips, showing thick glass filaments at the

end of molten particles; width of picture: 2 mm.

Borosilicate glass

58 m/s

melting

02 OA 08 01

Continuum chip thickness h eq (um]

02 08

Fig. 13. Borosilicate glass. Force components F

t

, F^ versus h

g ^ for v

g = 58 and 9^ m/s,

68
k.2. Melting

For several glasses a quite different result has been obtained: at high wheel speeds

melting occurs and he ^ is no longer a useful parameter. The heat input q into the workpiece

can he derived from the power in eq.(2) and the ratio of the thermal "impedances" of wheel

and workpiece [ 21-2U] ,

^wheel

'"'workpiece

wheel

(kpc) ~~

workpiece

(10)

where k is the heat conductivity, p is the density and c the heat capacity per unit mass. For

a steel wheel on a ceramic workpiece 6 to 8% of the total heat generated enters the workpiece, the rest goes into the wheel and the cooling water. The temperature at a depth y can be calculated from a model where the heat source, W=F^.v

s

, moves over the surface at a speed vw ; the result is [21-23]

T(y) = T(o)exp (-y/y

Q

). (11 )

In this expression T(o) and y are functions of l

(,vv

./qt, where a is the thermal diffusivity

a = k/pc; the contact length 2 C is given by lc = faD) 5 for horizontal surface grinding

(fig. 1), where D is the wheel diameter. Without going into further details , it can now be seen that the equivalence of depth of cut a and feed v , which was present in eqs.(l) to (9)>

is lost when the temperature affects the behaviour of ¥he material. In the most extreme circumstances of our grinding apparatus the surface temperature is calculated to be 1CrK, and

at a depth of cut of hOO um the borosilicate glass reaches its ASTM softening point (800°C).

The appearance of glass surface can thus be understood in a qualitative way. As in the roughness problem discussed earlier, not all of the damage due to melting and cooling will be

present on the surface afterwards.

If melting or, less severe, softening occurs, a shear deformation of the hot surface

material must be considered. The necessary shear stress x probably contributes more to F^. than to Fn

, which agrees with the normal behaviour of the Fn

- he^ curve (fig. 13). With

the viscosity of the flowing material, T is given by

dv v

T. = T) - ~ - ri — (12)

x is related to F by

dy yQ

F, = T 1 b. ( 13) t c

The specific energy is derived from {k) , (12) and (13),

e = F+ v /Z = n v

2

1 b/y Z ( 1U)

t S SCO

In this simplified model a strong influence of vs on e is therefore expected. This may account for the relatively high e values found here at high v , even in quartz glass, in com- parison to Huerta's results at 19 m/s [20] . The increase of e with the softening temperature

[20] , is also in agreement with the r) dependence in eq. (1U).

k.3. Size effect

The specific energy depends on the rate of removal Z, or, as in figs. 8 and 9» on he ^.

This "size effect" is usually attributed to the behaviour of the material and must then

69
be considered on a microscopic scale, where the depth of cut of each diamond, d, is used

as a relevant variable [ 5 9 8] . To derive d from Z or h

g(^ one assumes that each of the active

diamonds (total instantaneous number N^, speed vs

) makes a groove with cross-section A,,

Z = N-A^v . d d s (15)

For pyramidal particles (apex angle 6) the grooves are triangular. In horizontal axis grinding the groove starts at a point and widens over the contact length lc

. The cross-section,

averaged over l

c

, is

A, = d /3 cot 6

d max

(16)

where d^ is the depth at the end of the groove. With eqs.(7), (15) and (16), dmax is given

by

d = (3 cot 6 h .b/N, max eq d

(17)

For groove formation at low scratching speeds [9, 10] the force ratio f equals cot 6;an

estimate of cot 0 at high speeds can also be made by using f . A number of methods are known

for obtaining the N- value [ 25] ; here we used a roughness measurement by a roller over the

circumference of the wheel. The result conforms with what the figures 2 and 10 already indicated, i.e. a distribution of heights of the diamond type must be used. Introducing the

probability density N(y), A^ is given by

max

A

d

=

1 3 cot 9 N(y)dy (18)

0

where y runs from the outermost diamond tip to the depth d^^. In the roughness experiment

a discrete distribution was obtained, counting the number of peaks at intervals of Ay = 5um

(no shadow correction was applied [26] ). The integral in eq.(l5) was evaluated as a sum for

dmax up to 60 urn. Together witheqs.(7) and (15) a numerical relation between dmax and he ^ is

found which replaces eq.(l7). The specific energy can now be taken from figures 8 and 9 and

replotted as a function of dmax (figs. 1^ and 15). For the hexagonal ferrite we have added

data taken from measurements at

100

depth ( u m )

Fig. 1H Specific energy as a function of calculated grit depth for SrFe. ?0 and (Ni,Zn)

50

E

£

Ol

u

(0

Xi

o slow scratch

GROOVE

1 1

\ GRINDING

5 10 20 —» d (pm

)

50 100

Fig. 15. As Fig. 1 U for (Mn,Zn)Fe

2 0^

,

Fe 2°U70
vg = 5 m/s (an average N, was used [ 2) ) and from low-speed scratching [ 10,11] . The agreement

is remarkably good. For the cubic ferrite the results show again the strong decrease of

the specific energy at the larger depths of cut, much more than found upon scratching the

same material (fig. 15). It should be remembered, however, that the value of the specific

energy for scratching has been derived from the groove parameters only, disregarding the

chips that were made and may still be fastened to the surface. In grinding the chips are loosened too and must be included in the volume removed. This, however, has not been taken

into account in the derivation of eqs. ( 1 6— 1 8 ) . Formally, one could introduce a wider groove

by using a lower value for cot 8. The result is an even stronger decrease of e with increasing d. On the other hand, the agreement between e values from grooving and scratching for

SrFe^O-i^ (fig. 1M should correspond to the absence of large-scale chipping in this material. Such is indeed the case: when the extent of cracking and chipping in the hexagonal

ferrite is compared with that in the MnZn ferrite one finds that the latter is much larger.

By analogy with the half-penny crack model [ 27] the extent of chipping can be expressed in

terms of the fracture toughness value, K^, by

c « (Fn/TTK

].) 2/?

(19)

<

where Fnis the load on the scratching diamond. The Kj values obtained from a three-point

bend test [2] , are

SrFe

12

0

19

K^ = 2 . 1 x 10

6 N/m2 ^ 2

(Mn, Zn) Fe

2

0

1+ ^ = 1.1 x 10^ N/m3//2

.

We conclude that the different chipping of these materials may be responsible for the

different size effects. A grinding model for brittle materials should therefore include the

consequences of fracture on a microscopic scale.

k.k. Grinding models

Characteristic quantities for grinding are the force ratio, the depth of damage and the

specific energy. The force ratio has a simple meaning for scratching at low speeds:

f = cot(Q + a), where 28 is the apex angle of the scratching pyramid and tan ot is the Coulomb

friction on the interface between diamond and workpiece [9,10] . At high speeds heat is developed that may modify the material: softening for metals [28] , melting in glass; the simple

"ploughing" relation between f and 0 is then no longer valid. The force ratio seems to

be a good measure of the sharpness of the abrasive particles, which need not be the same when the wheel is used in fine and coarse grinding, because more and different particles are active in the latter case.

The depth of damage, as measured by the surface roughness, Ra

, does not allow a simple

interpretation as was argued in section Compared with the dmax values derived from

h

e(^ in section U.3. , Ra

is much smaller. A much better description of damage would be given

by the subsurface cracking, but this requires much more experimental work and is suitable

for laboratory investigations only.

The specific energy, the third characteristic quantity of grinding, measures the resistance to material removal. Under the action of the diamonds the material is heavily deformed

and grooves, cracks and chips are formed. The groove is found at all loads on the diamonds,

but cracks appear outside the groove at a force level that is given by [ 29]

k q F . * K/H . (20) mm I v

For a grinding force below this value one could speak of light grinding. Apparently, this

definition depends on the material: for metals no chips are formed, while for a ceramic

71
like SrFe^O-ig grooves are also the predominant deformation. For such materials the value

of the specific energy may he taken from a single-point scratching experiment.

For coarse grinding of brittle materials, where chipping does occur, the effects of

fracture must be included. Early work by Gielisse [ 21] stressed the importance of this by

taking the energy density for macroscopic fracture (o f2/2E), with fracture strength and

elastic modulus E, as the only measure of grinding resistance. Although the strength may

be determined by surface flaws the situation seems to be more complicated: microscopic crac- king differs from macroscopic in that the latter needs only one (large) crack for fracture

whereas for material removal by chipping one needs, firstly, the nucleation of cracks along

an arbitrary scratch, secondly the interaction of such cracks for chips to form, and thirdly, the loosening of the chips from the surface. Due to the large stress near the groove

the crack nucleation seems to be no problem. On increasing the load on a scratching point

first lateral cracks appear, which, at higher loads, combine to form chips [11]. When the

scratching point is loaded by more than the minimum given in eq. (20), the lateral cracks

grow in size and number until chips are formed at higher loads. Their extent on the surface

is about that of the lateral cracks, their depth and number, however, are much less. Chips

become effective only when loosened and this seems to be the bottleneck in low-toughness

materials like the MnZn ferrite used here.

The loosening action has been considered before by Wiese and Wagner in an identation

model for lapping. The cracks from neighbouring indentations interact in a random way to

loosen lumps of material [18] . For scratching Veldkamp f 2] made an estimate of the probability of interaction by considering the newly created surface. Neither model is concerned with

the effect of distance between neighbouring scratches on the volume taken away. This effect

was determined experimentally by making parallel scratches under a given load at various

distances and measuring the volume removed [ 11] (fig. 16).

E7 10

Weight loss per Mn Zn ferrite

unit scratch length

Load 4.9 N x

1.2 N o

\X\

0 100 200 300

Distance between scratches (um| 400

Fig. 16. Mass removed by making parallel

scratches on a polished surface of a MnZn

ferrite, as a function of the distance

between the scratches [11].

This result means the following: at large distances each scratch acts independently, the

volume that is removed is practically that of the groove, while the chips that are present

with a certain density remain fastened to the surface. When the distance is decreased some of the chips from one scratch interact with those from the neighbour and more material is

removed. At a distance of about twice the average extent of chipping, the volume removed

rises steeply. At a distance of about the groove width, the separate scratches begin to

act as one scratch and the loosened volume is expected to decrease. The extent of chipping

at a given place along a single scratch and the number of chips per unit length are random

quantities in that material irregularities are unpredictable. For two neighbouring scratches

it may be that the cracking and chipping along the second scratch is influenced by the presence of the damage due to the first one, but this is not known at present.

It seems that the best one can do is to collect data about interacting scratches and

use these as a basis for a multipoint grinding model. Here the removal due to grooving and

chipping are added; eq. (15) becomes

Z = Z + Z . = N,v (A + A , ) gr ch d s gr ch

72
where is the volume removed per unit length of scratching* as plotted in fig. 16. In

view of the distribution of heights and distances (the Z- and X-distributions of ref. 13)

Ac j1 is a sum over curves as shown in fig. 16. The power W necessary for the removal rate

in eq. (21) is mainly determined by the formation of the groove, because cracking requires

little energy and can be ignored:

W = e Z . (22) gr gr v '

The specific energy follows

e Z e A

W gr gr gr gr , } Z Z +Z A +A * ; gr ch gr ch

With increasing depth of cut, A ^ increases sharply for two reasons: firstly, the chip

volume increases and secondly, the interaction between chips increases (fig. 16). If this

interpretation of the size effect in fig. 15 is correct, this means that fairly detailed

information about the grinding wheel is needed. In that case the specific energy is not

a material property, as in groove formation, but a property of the grinding system.

5. Conclusions

Grinding experiments at wheel speeds of 60-100 m/s have shown that the force components

are unique functions of the continuum chip thickness, heCp defined by eq. (7). For a quick

comparison of grinding operations he q is useful (e.g. hec, = 0.01 urn for light grinding,

he q

= 1 um for coarse grinding). Problems are encountered when melting occurs, as in the

borosilicate glass at vs = 100 m/s. Depth of cut and feed are then no longer equivalent and

he q loses its meaning.

For microscopic details, however, more information is necessary. The surface quality

as measured by the c.l.a. roughness, is a weak function of he q. This may be related to the

variation of the grinding force over the contact length, the maximum damage being removed,

the minimum being left behind (sect. It would be better, but more cumbersome, to measure the subsurface damage, especially cracking. This is also relevant for the matching

of subsequent stages in a grinding operation.

The specific energy has been derived from the data and considered in some detail. The

dependence of e on the grit depth d^x is found to be rather weak in a hexagonal ferrite.

The e values are in good agreemetn with those found at a lower grinding speed and in scratching at very low speeds. This material shows small chips only, and the groove energy is a good measure of the grinding energy.

Quite a different behaviour was found in a (Mn,Zn,Fe) ferrite, where a strong size effect is found. This material shows large chips, and arguments have been presented that the

interaction of neighbouring scratches is responsible for the increase in removal.

A grinding model that goes beyond simple grooving must be based on the workpiece behaviour, i.e. the extent and frequency of chipping and the more or less random interaction

between neighbouring scratches. It seems that mult i scratching experiments can provide the

necessary information. The model cannot avoid rather detailed information about the diamond

distribution on the grinding wheel. The height distribution is needed for the number of

active diamonds at a given depth of cut, while the transverse distribution (pairwise) allows the scratch interaction to be treated. For brittle materials of the second kind the

specific energy, especially the size effect, is a property of several components of the

grinding system.

The authors wish to thank a number of colleagues for their contributions to this work:

J.C.G. Teunissen and M.A. Brebels for their help in building the grinding apparatus;

G.F.M. Gasseling for his contributions to the experimental results and D. Snip for many

discussions during several years of close cooperation on the grinding of ceramics.

73
References

[I] Proc. Symp. on Ceramic Machining and Surface Finishing, Gaithersburg, Md., November 2-k

1970, Book, The science of ceramic machining and surface finishing, S.J. Schneider Jr.

and R. Wo Rice eds. (NBS Spec. Publ. 3^8, 1972).

"

[2] Veldkamp, J.D.B. and Klein Wassink, R.J., Grindability of brittle materials, Philips

Res. Repts., 3J_ 153-189 (1976).

[3] Broese vanGroenou, A., Veldkamp J.D.B. and Snip, D. , Scratching and grinding parameters

of various ferrites, J. de Physique, C1, 285-289 (1977).

lh] Bro ese vanGroenou, A., Experiments on grinding of ceramic ferrites, (Proc„ 9 th Int.

Conf. on Science of Ceramics, Noordwijkerhout , November 13-16, 1977) » Book, Science of

Ceramics 9 » K.J. de Vries, ed., pp. U86-^93 (Nederlandse Keramische Vereiniging, 1977).

[5] Shaw, M.C., Fundamentals of grinding, in New Developments in Grinding , M.C. Shaw,-ed.,

pp. 220-258 (Carnegie Press, Pittsburg, Pa., 1972).

[6] Snoeys, R., The mean undeformed chip thickness as a basic parameter in grinding, Annals

C.I.R.P., 20[2], 183-186 (1971).

[7] Snoeys, R., Peters, J. and Decneut, A., The significance of chip thickness in grinding,

Annals C.I.R.P. , 23 [2] , 227-237 097*0.

[8] Koepke, B.G. and Stokes R.J., Grinding forces and the machining of magnesium oxide

crystals, J. Mater. Sci. 7, U85-^93 (1972).

[9] Maan, N. and Broese van Groenou, A., Low speed scratching experiments on steels, Wear

M» 365-390 (1977).

[10] Broese van Groenou, A., Maan,N. and Veldkamp J.D.B., Scratching experiments on various

ceramic materials, Philips Res. Repts, _30_, 320-359 0975).

[II] Broese van Groenou, A., Maan, N., and Veldkamp, J.D.B., Single-point scratches as a basis for understanding grinding and lapping (Proc. Conf. Ceramic Machining and Surface

Finishing II, 1978).

[12] Vaidyanathan, S. and Finnie I., The grinding of brittle solids, in New Developments in

Grinding , M.C. Shaw, ed. , pp. 813-831 (Carnegie Press, Pittsburgh, Pa., 1972).

[13] Law, S.S., Wu, S„M. and Joglekar, A.M., J. Eng. Ind. (Trans ASME), B95, 983-991 0973).

Izumitani, T., and Harada, S., Polishing mechanism of optical glasses, Glass Techn.,

J2[5], 131-135 (Oct. 1971).

[15] Izumitani, T. and Suzuki, I. Indentation hardness and lapping hardness of optical glass,

Glass Techn. jk[ 2] , 35-^1 (Apr. 1973).

[16 ] Broese van Groenou, A., Grinding of ferrites, some mechanical and magnetic aspects, IEEE

Trans, on Magn„ , MAG-JM[ 5] , lUl+6— 1U5 1 (Sept. 1975).

[17] Oh, H.L., Oh, K.P.L., Vaidyanathan, S. and Finnie I., On the shaping of brittle solids

by erosion and ultrasonic cutting, see ref. 1 of this paper, pp. 119-132.

[18] Wiese, G.E. and Wagner, R.E., Physical model for predicting grinding rates, Appl. Optics, _3 [11] 2719-2722 (Nov. 197*0.

[19 1 Gielisse, P.J., Kim T.J., Goyette L.F., and Nagarkar, R.V., Fundamental and practical

evaluation of the ceramic finishing process, Report, Univ. of Rhode Island, Kingston,

R.I., 1976.

[20] Huerta, M. and Malkin, S., Grinding of glass: the mechanics of the process, J. Eng. Ind.

(Trans. ASME), I159-U67 (May 1976).

[21] Gielisse, P„J„, and Stanislao, J. Mechanical methods of ceramic finishing, ref. 1 of

this paper, pp. 5-35.

[22] Takazawa, K. , Thermal aspects of the grinding operation, Ind. Diamond Rev., lU3-1*+9

(Apr. 1973).

[23] Maris, M. and Snoeys, R., Heat affected zone in grinding operations, (Proc. lUth Conf. on Mach. Tool Design and Research, Manchester 1973), paper Soh.

[2h] Maris, M., Thermische aspecten van de oppervlakteintegriteit bij het slijpen, Thesis

Univ. of Leuven, Belgium, 1977 (in Dutch).

[25] Horn, K., Uberblick uber moderne Verfahren der In-Prozess-Messung in der Fertigungstechnik, Annals C.I.R.P., 2k, 257-262, (1975).

[26] Tigerstrom, C.L., A model for the determination of the number of active grinding edges

in various grinding processes, Annals C.I.R.P., 2k_, 271-275 (1975).

[27] Lawn, B.R. and Fuller, E.R., Equilibrium penny-like cracks in indentation fracture,

J. Mater. Sci. _10_ 2016-202U ( 1975).

[28] Broese van Groenou, A., The high speed size effect in grinding: the role of heat generation, Wear hh[2] , 203-211 (1977).

[29] Veldkamp, J.D.B., Hattu, N. and Snijders, V.A.C. , Crack formation during scratching of

brittle materials, in Fracture Mechanics of Ceramics , Vol. 3, R.C. Bradt, D.P.H. Hassel- man and F.F. Lange, eds., pp. 273-301 (Plenum Publ. Corp. 1978).

74
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

EFFECT OF WORKPIECE PROPERTIES ON

GRINDING FORCES IN POLYCRYSTALLINE CERAMICS

B. G. Koepke and R. J. Stokes

Honeywell Corporate Material Sciences Center

10701 Lynda le Ave. S. Bloomington, MN 55420

Grinding forces have been measured on a number of polycrystalline ceramics

and correlations between the forces and workpiece properties have been attempted

in an effort to gain insight into mechanisms of chip formation. The closest

correlations were with a parameter related to the indentation fracture characteristics of each workpiece. This suggests that a major mechanism of chip formation

during multipoint abrasive machining of ceramics is fracture resulting from the

plastic indentation of the workpiece by the abrasive grains. Grinding forces

during up and down grinding were different but both correlated with the indentation fracture parameter when the comminution characteristics of the chips

during down grinding were accounted for.

Key Words: Chip formation; grinding forces; surface condition; surface grinding.

1. Introduction

The quality and usefulness of a precision ceramic component may depend entirely on the

abrasive machining operations used to shape and finish its surfaces. Yet it has been noted

that abrasive surface finishing is still an "art based on experience " [1,2] 2 . The re are good reasons for this comment. First, the mutual interactions between the components of a surface finishing system (e.g., machine, tool, workpiece, environment, etc.) are complex and

not well defined[2]. Second, the mechanisms whereby material is removed are not completely

understood. For example, finish surface grinding causes chip formation on an extremely fine

scale at high rates under conditions of stress and temperature that are difficult to measure

[2,3].

Nevertheless considerable progress has been made since the first Symposium on the

Science of Ceramic Machining and Surface Finishing [4]. It is now known that significant

plastic flow may accompany chip formation during grinding of even the most brittle ceramics

[5-8]. Scratch test studies [9,10] and work using model grinding systems employing single

point diamonds mounted on a grinding wheel [2,11] have clearly shown that the mechanism

of material removal depends on the normal force on the abrasive grain or, equivalently , on

the depth of cut. At low loads material is removed by plastic flow, as the load increases

cracking occurs and finally chips with dimensions larger than the abrasive grain form [10].

In a multipoint cutting process such as surface grinding, the subsurface interaction of the

fractures due to chipping represent a major mechanism of material removal[l0]. Single

point grinding studies have further shown that chips can form in front of and behind the

abrasive grain [3,8], the latter presumably due to relaxation of stored elastic energy.

In view of the complexity of the situation the question arises as to whether there is

any physical property or combination of physical properties which can be related to the

forces involved in machining a particular ceramic. From the scratch test and single point

grinding work it is recognized that the resistance to plastic flow and to fracture must

both be included in any model of chip formation [2, 7-9]. Low speed scratch tests have

-•-Work carried out under Contract N00014-69-C-0123 to Office of Naval Research.

2Figures in brackets indicate the literature references at the end of this paper.

75
shown that the forces vary with hardness when plastic flow is the dominant mode of material

removal [9]. Gielisse et al. [11] have demonstrated in single point grinding that the

forces vary as 0

2 /2E, where a is the fracture stress and E Young's modulus, when brittle

fracture is the dominant mode of material removal.

Recent analyses of the relation of material properties and chip formation during

ceramic machining have benefited from the indentation fracture theories of" Lawn and co- workers [12,13]. Evans and Wilshaw [13] have modelled abrasive machining as a series of

sliding parallel indenters and have shown that the chip volume, V, is V ^ p5/4 h~1'2kc ~3/4

where P is the vertical load on the abrasive grain, H is the hardness and K

£

is the

fracture toughness. Sawing rate data of Rice and Speronello [14] fit this relation well.

Kirchner et al. [15] have measured the depth of damage, c, introduced in Si3N4 during single

point diamond grinding and have shown it tc be related to the vertical grinding force, P,

according to a relation derived by Lawn and Marshall [16] based on identation fracture

theory P ^ Kc c\l ^- . Thus, studies using single point abrasive systems and analyses incorporating the relatively new identation fracture theories have both led to an improved

understanding of the important physical parameters determining material removal during

abrasive machining of ceramics. Whether these results can be directly applied to multipoint machining operations such as surface grinding is questionable. The simultaneous

interaction of adjacent abrasive grains during chip formation will alter the relative

amounts of plastic flow and fracture and thus the resulting forces. Which dominates and

under what conditions are not clear.

In this paper we attempt to establish correlations by examining the effects of workpiece properties on forces generated during the surface grinding of a number of different

ceramics. It was our hope that a clear correlation would establish the predominant

mechanism of material removal under given conditions. Correlations are attempted between

the forces and the workpiece properties to determine, at least, what properties most

affect the mechanism of chip formation.

2. Experimental Procedure

2.1 Grinding

Grinding was carried out using a DoAll Model D618-7 high speed precision surface grinder

with a fixed spindle speed of 3540 rpm. A 100 grit resinoid bond diamond wheel^ 15.2 cm

(6 inches) in diameter and 1.32 cm (1/8 inch) wide was used for all measurements. The

cutting point density on the wheel face was 480 cm

-

^ and the mean width to depth ratio for

a groove made by a single abrasive grain, as determined by techniques described by Backer

et al. [17] was 1.6. Grinding forces were measured with a strain gauge dynamometer based

on a design of Cook et al. [18]. The dynamometer is similar to those used by others [2] and

will not be discussed. With a sample mounted on the dynamometer, the horizontal and vertical grinding forces induced in the workpiece during a grinding pass can be measured independently. The elastic stiffness of the dynamometer is 1.2 x 105 gm/ cm in the horizontal

direction and 4.5 x 10^ gm/ cm in the vertical direction. The sensitivity of the instrument

is 3 gm in both directions. Grinding was carried out both with the cutting face of the

wheel moving in the same direction as the workpiece (down or climb grinding) and with the

wheel moving in the opposite direction (up grinding) . Different parts of the same workpiece were used for each case, i.e. one part of the surface was up ground and a separate

part was down ground. Samples were ground at feed rates of 0.042, 0.21, and 0.42 cm/sec

(1, 5 and 10 inches/min) . The depth of cut was 25.4 ym (10~3 inches) except in one case where it was half this value. At least 5 passes were made at each feed rate and each

sample was always ground at increasing rates of material removal. The wheel was dressed

by grinding the face of the wheel with a silicon carbide wheel before each sample was

ground at the lowest rate of material removal.

The theoretical chip geometry (i.e. chip length and maximum chip thickness) [17] for

the grinding conditions used here are listed in table 1.

^DoAll Type. DlAl-Mdl00N/lDoBl/4B3

76
From table 1 we note first that the grit depths of cut, t, are small and are on the

order of 10

-

^ of the chip lengths. An indentation made by an abrasive grain on a ceramic

workpiece is, therefore, no more than a long thin scratch [5,9] provided material is

removed by plastic flow. As noted above, however, if brittle fracture takes place during

chip formation the actual chip geometry may in no way resemble the dimensions in table 1.

Note also that the small grit depth of cut indicates the advisability in many instances of

using small wheel depths of cut. Otherwise a large amount of the binder in the wheel face

may contact the workpiece.

Table 1. Theoretical maximum

diamond wheel at the

chip thickness, t, and

machine settings used

chip length, I, in this study.

for the 100 grit

Wheel Depth of cut Feed Rate

(in/min)

t (ym)

I (mm)

100 grit 25.4 urn .042 cm/sec .32 2.0

.21 cm/sec .71 2.0

.42 cm/sec 1.0 2.0

Samples used for grinding were approximately 25.4 x 6.4 x 3.2 mm (1 x 1/4 x 1/8 inches)

Grinding was carried out on the 25.4 x 6.4 mm side in the short (6.4 mm) direction. All

grinding was carried out dry.

2.2 Characterization of the Ground Surface

Both optical and scanning electron microscopy were used for routine characterization

of the surfaces produced by grinding. Most observations were made after the last pass at

the highest feed rate of 0.42 cm/sec.

2.3 Workpiece Materials

The materials machined in this study are listed in table 2, are the supplier, method of manufacture and grain size.Also included in the table

Table 2. Materials used for Grinding Force Measurements

Material Supplier Manufacturing

Method

Approximate

Grain Size

(Vim)

PZT Honeywell Cold Pressed and

Sintered

MgO IITRI Cold Pressed and

Sintered

18

Ferrite NiZn

Ferrite NiZn

Silicon Nitride0

lion

Honeywell

Honeywell

Westinghouse

Hot Pressed

Hot Pressed

Hot Pressed

13

3

2

Silicon Nitride a

Silicon Nitride3

Norton

Westinghouse

Hot Pressed

Hot Pressed

Lucalox Alumina

Alumina''

AD 999 Alumina

Boron Carbide*5

G.E.

AVCO

Coors

AVCO

Cold Pressed and

Sintered

Hot Pressed

Cold Pressed and

Sintered

Hot Pressed

38

2

2

aCourtesy Dr. F. F. Lange

''Courtesy Dr. J. Niesse

77
The Westinghouse silicon nitrides listed in the table differ as follows [19,20].

"3" Si3N^ was fabricated from powders containing 90% 3~phase Si^N^ and 10% a-phase

Si3N^. "a" Si3N4 was fabricated from powders containing 90% a-phase and 10% 3 -phase

powder. All three materials were 3 phase after hot pressing. The three different Si3N^

materials are interesting because, as noted later, they exhibit measurably different

properties

.

3. Experimental Results

3.1 Grinding Force Measurements

In order to conserve space the results are presented in tabular rather than graphical

form although a few curves will be presented for illustrative purposes. The grinding

force measurements are listed in the Appendix. Each value of force is normalized by the

width of the cut, b, and is the average of five passes made at the depths of cut and feed

rates indicated. Some brief comments concerning the force measurements are given below.

Grinding forces on PZT, MgO and the ferrites were extremely low and without exception

increased with rate of material removal. The forces measured during up and down grinding

were about equal. The vertical forces measured on the smaller grained ferrite exceeded

those measured on the larger grained material. The horizontal forces measured on both

ferrites were not a function of grain size.

The forces measured on Si3N4 during 25.4ym cuts were the highest measured in the series.

In order to minimize damage to the dynamometer or wheel, measurements of the effect of rate

of material removal on grinding forces were made using a cut of 12.2Um. We note here that

in a number of cases the vertical forces measured during down grinding exceeded those

measured during up grinding in contrast to results on other hard ceramics such as B4C and

AVCO AI2O3. Examples of this behavior are shown in figure 1 in which the grinding forces

for the Westinghouse "a" Si3N4 are plotted as a function of feed rate during cuts of 12.2lim.

O DOWN

A UP

•

I 2 .3 .4 .5 .1 .2 .3 .4 .5

V ( cm/sec ) V ( cm/sec )

Figure 1. Horizontal, FH

/b , and vertical, F

y

/b , grinding forces measured on Westinghouse

"a" Si 3 N4 during up and down grinding plotted as a function of workpiece feed rate, b

is the width of the cut. The depth of cut, d, is 12.2ym (0.0005 inch).

The forces measured on alumina depended on fabrication. Sintered material such as Lucalox and Coors AD 999 showed low grinding forces at low feed rates but the forces

increased rapidly with feed rate. AVCO hot pressed alumina on the other hand sustained

a relatively high force level at all feed rates as indicated in figure 2. The low

78
sensitivity to feed rate during down grinding is not understood but may be related to the

comminution characteristics of the chips as they travel between the wheel and the workpiece.

This point is discussed later.

4 00T

300 -o

o

~ 200

100

AVCO Al

2

0

3

A—

o - DOWN

A - UP

GM/CM

I200r° oo

100 GRIT *§ 600 -

d = 0.025 mm ^

400 -

200 ~

.2 .3 .4 .5

100 GRIT

d = 0.025 mm

J I L

.2 3 4 .5

V ( cm/sec

V ( cm/sec)

Figure 2. Horizontal, %/b, and vertical, Fv/b, grinding forces measured on AVCO hot

pressed AI2O3 during up and down grinding plotted as a function of workpiece feed rate,

b is the width of the cut. The depth of cut, d, is 25.4ym (0.001 inch).

At the lowest feed rate the forces on B^C during down grinding were greater than those

during up grinding. This behavior reversed at the higher feed rates.

3.2 Examination of the Ground Surfaces

The ground surfaces of every sample were examined following the last grinding force

measurement which, as noted earlier, was made at the highest feed rate used in this work,

0.42 cm/sec. The nature of the surfaces will be commented on and illustrated in a few

cases with scanning electron micrographs.

The ground surfaces of PZT were smooth and severely burnished. This accounts for the

high grinding forces measured on this relatively soft ceramic.

The MgO surfaces were similar to the ferrite surfaces and contained plastically flowed

regions and regions where material had been removed by brittle fracture. These surfaces

are similar to those examined in an earlier study [6].

The large grained ferrite showed more intergranular fracture than plastically flowed

material. The fine grained ferrite showed large flowed regions with some patches where

material was removed by intergranular fracture. Examples of these surfaces are shown in

figure 3. Recall that the grinding forces on the fine grained material were slightly

higher at the highest feed rate.

Without exception all silicon nitride surfaces were quite smooth and burnished and

consisted mainly of plastically flowed regions. Patches where material was removed by

what appeared to be intergranular fracture were seen scattered over every surface.

79
The sintered aluminas had surface features resembling the Si3N4 samples. A typical

example is shown in figure 4(a) for Coors AD 999 alumina down ground, 0.42 cm/sec. The

irregular patches are where material has been removed by intergranular fracture. In

contrast, the hot pressed alumina showed fewer burnished regions as shown in figure 4(b).

Most material has been removed by transgranular fracture. The B^C surfaces were similar

to those of Si-jN^ and the sintered aluminas and contained regions of plastically flowed

material and patches where material had been removed by brittle fracture. Due to a large

amount of debris on the surfaces the nature of the fractures could not be assessed.

GRINDING DIRECTION

(a) 1 20 ym ' ( b )

' 20 ym '

Figure 3. Scanning electron micrographs of the ground surfaces of NiZn ferrite with mean

grain sizes of (a) 3 and (b) 13ym.

GRINDING DIRECTION

Figure 4. Scanning electron micrographs of the ground surfaces of (a) Coors AD 999

alumina and (b) AVCO hot press.

80
A few general comments concerning the nature of the ground surfaces are appropriate

at this time. First, no significant differences were noted between surfaces produced

by up and by down grinding. This was unexpected in view of the differences in forces

observed. Second, all surfaces contained significant amounts of plastic flow. The large

grained ferrite and hot pressed alumina contained less plastic flow than the others.

Finally, it must be cautioned that the conditions observed here may not be identical to

those obtained in actual practice. The samples were ground dry and were mounted on a dynamometer whose compliance is significantly greater than that of the machine table itself.

Both dry grinding and a lower compliance may enhance burnishing.

4. Grinding Force Correlations

The properties used for the correlations are listed in table 3. Note that many of the

values in the table were not measured directly on the material used in this study, but were

taken from the literature or supplied by individuals noted in the references to the table.

In all cases the vertical and horizontal grinding forces measured during both up and down

grinding during a 25.4ym cut at a workpiece feed rate of 0.21 cm/sec were used for

correlations

Table 3. Representative physical and mechanical properties of ceramics.

Bend Poly crystalMaterial

Grain

Size

(ym)

Melting

Point (°C)

Knoop

Hardness3 (kg/mm 2

)

Tensile

Strength

(MN/m2

)

line Young's

Modulus

(GN/m2

) 3/2 K

IC (MN/m )

PZT 3 1300 200 * 55

b 83b .75h

MgO 18 2850 400 * 28 c 303 c 2.5 1

NiZn Ferrite 13 1660 550 # 166 d 2006

NiZn Ferrite 3 1660 650 # 124d 200 8

"3" Si

3 N

4

2 Sublimes

at 1870

1450 * 375 c 307e 3.1e

Norton Si3N^ 1 Sublimes

at 1870

1500 * 676 c 30 7 e 5.1e

"a" Si

3 N4

1 Sublimes

at 1870

1550 * 656 c 30 7 e 6.5 e

Lucalox AI2O3 38 2050 1800 * 440 c 400 c 4.lJ,k

AVCO A1

2 0 3 10 2050 1850 * 407 f 400 c 6.5l

AD999 A1 2 0

3

2 2050 2000 * 469 c 400° 4.0m

B

A C 8 2350 2600 * 448 f 448 c 4.2n

a. All values measured in this laboratory. (// lOOgm load, * 500gm load).

b. Honeywell data sheet.

c. Engineering Properties of Selected Ceramic Materials , American Ceramic Society,

Columbus, Ohio (1966).

d. M. Braitburg, Honeywell Inc. private communication (1972).

e. F. F. Lange, Journ. Amer . Ceram. Soc, 56^, 518 (1973).

f. AVCO data sheet.

g. J. Smit and H. P. Wijn, Ferrites , Wiley & Sons, New York (1959) p. 225.

h. J. G. Bruce, et al. Fracture Mechanics of Ceramics, Vol. 4, R.C. Bradt et al. ed.

Plenum, N.Y. (1978) p. 696.

i. P. Gutshall and G. Gross, Eng. Fract. Mech., 1_, 463 (1969).

j. R. F. Pabst, Fracture Mechanics of Ceramics, Vol. 2, R. C. Bradt et al. ed., Plenum

N.Y. (1974) p. 557.

k. S. W. Freiman et al. ibid p. 666.

1. Measured in this laboratory by hardness indentation technique of Evans and Charles,

Journ. Amer. Ceram. Soc, 59_, 371 (1976).

m. L. M. Barker, Fracture Mechanics of Ceramics, Vol. 4, R. C. Bradt, et al. ed. , Plenum,

N.Y. (1978) p. 490.

n. G. Gazza, Army Research Center, Watertown, MA., private communication (1972).

81
4.1 Melting Point

No obvious correlations with melting point were noted.

4.2 Hardness

Since hardness is an indication of resistance to plastic flow an examination of the

relation between grinding forces and hardness is warranted. The horizontal and vertical

forces measured during down grinding are shown as a function of workpiece hardness in

figure 5. There is a trend toward increasing vertical grinding forces with hardness but

the horizontal forces drop in the harder materials. In both cases the highest grinding

forces are those measured on ceramics exhibiting intermediate values of hardness. Similar

trends were observed during up grinding but the forces measured on ceramics with intermediate hardnesses were lower than those shown in figure 5.

i- o £j

a. 2

ro

ro o

OD 2000 1 I I 1

1

o

(S>

1800 -

1600

1400

12 00

1000

800

600

400

200

TT1—ITT

DOWN

GRINDING

1000 2000

KNOOP HARDNESS

3000

Figure 5. Horizontal, Fg/b and vertical, Fv /b, grinding forces measured during down

grinding on different ceramics plotted as a function of workpiece hardness. The depth

of cut and feedrate were 25.4um and 0.21 cm/sec respectively.

82
4.3 Tensile Strength

Grinding forces measured during down grinding are shown as a function of tensile strength

in figure 6. There is a trend for the forces to increase with strength but again the forces

measured on materials with intermediate strengths are high. Similar behavior is observed

when the forces measured during up grinding are plotted versus tensile strength.

4.4 Poly crystalline Young's Modulus

No correlations were noted but again the forces were highest on materials with intermediate values of modulus eg. Si^N^

4.5 Strain Energy Density, a

2 /2E

In single point diamond grinding studies on alumina workpieces Gielisse et al. [ll]

obtained excellent correlation between grinding forces and the strain energy density

0^1 2E,

where G is the tensile strength and E is Young's modulus. No correlations were noted when

the grinding forces measured in this work were plotted versus a^/2E.

2000

1800

1600

14 00

5 1200

o

s

o- 1000

U.

I JO

800

600

400

200

100 200 300 400 500 600 700

BEND TENSILE STRENGTH (MN,m 2

)

Figure 6. Horizontal and vertical grinding forces measured on different ceramics plotted

as a function of the bend tensile strength of the workpiece materials. The depth of cut

and feedrate were 25.4ym and 0.21 cm/sec respectively.

83
4.6 Fracture Toughness

It can be argued that, if a major mechanism of material removal during grinding is the

initiation and/or propagation of brittle cracks, then grinding forces should correlate with

fracture toughness. Figure 7 shows that the grinding forces measured during up grinding

do increase with the fracture toughness, but not linearly.

The grinding forces measured on Lucalox were low with respect to the other aluminas

particularly during up grinding. The vertical forces measured on Lucalox are indicated

on the plot. The low forces are attributed to the large grain size of this material.

Gielisse et al. [11] have noted a similar sharp decrease in forces during single point

grinding of aluminas when the grain size exceeded about 20]im.

No conclusions between grinding forces measured during down grinding and KIC were

ob tained.

2000

1800 -

1600 "

UP GRINDING

1200 "

Figure 7. Horizontal and vertical grinding forces measured during up grinding on different

ceramics plotted as a function of the fracture toughness, Kjq, of the workpiece materials.

The depth of cut and feedrate were 25.4pm and 0.21 cm/sec respectively.

84
4.7 Indentation Fracture Parameter

As noted in the Introduction, recent theories of plastic indentation fracture of brittle

materials have led to increased understanding of the mechanisms of material removal during

the high speed impact of ceramics. A major result of the plastic indentation of a brittle

solid with a sharp indenter is the formation of subsurface lateral cracks lying roughly

parallel to the surface. When these cracks intersect the surface or other fractures a chip is formed and material is removed. The volume of material removed is proportional to

the size of these cracks which, in turn, is governed by the vertical force on the particle,

P, the hardness, H, and the fracture toughness, Kq. When material removal during abrasive

machining is assumed to take place by a parallel array of sliding indenters , Evans and

Wilshaw [13] have shown the upper limit for the volume removed to be V ^ P '^K "3/4 H

--'-'^. Thus, under conditions of a fixed rate of material removal (ie. fixed feed rate and depth

of cut) the grinding forces are expected to be close to a linear function of Kq^M Hl/2.

Figure 8 shows this to be the case for data taken during up grinding. The low data point

represents the large grained Lucalox data discussed above. The high point is Norton Si3N4-

Figure 8 represents by far the best correlation we have found between material properties

and grinding force measurements

2000

1800

1600

1400

£ 1200

o

s

2. 1000 -

U-|jO

800

600

400

200

UP GRINDING

100

K

|G

H

200

Figure 8. Horizontal and vertical grinding forces measured during up grinding on different

ceramics plotted as a function of an indentation fracture parameter derived by Evans and

Wilshaw [13]. K „ is the fracture toughness and H is the Knoop hardness. The depth of

cut and feedrate were 25.4um and 0.21 cm/sec respectively.

85
When the grinding forces measured during down grinding are plotted in a similar fashion,, no correlation is obtained. In what follows we attempt to rationalize this lack of correlation by considering a basic difference in the mode of material removal during up and down

surface grinding of ceramics on a dynamometer. During up grinding chips are formed at the

leading edge of the contact area between the wheel and workpiece and are immediately ejected

from the system. During down grinding the ships formed at the leading edge of the wheelworkpiece interface must first travel between the workpiece and the wheel to the back edge

of the contact area before they can be ejected. The initial chip size in both cases is

larger than the theoretical chip size and is presumably determined by the indentation fracture characteristics of the workpiece and the cutting point geometry on the wheel surface

as discussed by Evans and Wilshaw [131. During up grinding the vertical grinding forces

measured by the dynamometer depend mainly on the forces necessary to form the chips. During

down grinding, on the other hand, the measured vertical grinding forces must also be related

to the impact fracture or comminution characteristics of the isolated chips. If the chips

break up readily as they travel between the wheel and the workpiece the vertical grinding

forces will be reduced. If the chips do not break up readily, high vertical grinding forces

2000

1800

Jo

DOWN /

GRINDING /

NORTON

Si3 N 4

1600 /

1400 «=•< Si 3 N4 O

t

CM) in Aft / AVCO

/

A, 2°3 O

CD

Ll. |-0

10 00

800

e si 3

N4 o /

/ /° B 4

C

/LUCALOXq /

600 / / AD999 Al 2°3

400

200 "O / °MgO

PZT

i

i i 1 i i 2 4 6 8 10 12 3/4 1/2

K H

Ln N

Figure 9. Vertical grinding forces measured during down grinding on different ceramics

plotted as a function of the indentation fracture parameter of Evans and Wilshaw [13]

divided by a toughness parameter suggested by Clarke et. al [21] that can be related to

the comminution characteristics of the chips.

86
are expected. A parameter relating to the ability of a ceramic to resist impact fracture

has been suggested by Clarke et. al. [21]. The number of pieces a cubic centimeter of

material would break up into under the influence of an elastic-energy density, ^Vl, is

(W/6y)

J where y is the fracture surface energy and W can be approximately by a /2E. a is

the fracture strength and E is Young's modulus. Clarke et. al. [21] suggested that a suitable toughness parameter is -In N. If N is large the chips break up readily and their effect

on the vertical grinding forces should be small. By relating this concept to surface grinding we suggest that the vertical grinding forces measured during down grinding vary according to F ^ H^'^/ln N. This relationship is shown plotted in figure 9. Note first

that a correlation is obtained indicating that the comminution characteristics of the chips

do affect the vertical grinding forces measured during down grinding. Note also that the

data are in two distinct groups. The data of figure 9 suggest that a major reason for the

large difference in grinding forces measured during down and up grinding on Si QN is that once formed, the chips break up only with much difficulty and cause high forces during down

grinding. The other group of data including the aluminas and B^C also show good correlation. MgO is low presumably because the chips are easily deformable [5] and do not cause high vertical grinding forces. No correlations were noted when the forces measured during

up grinding were plotted in this manner.

5. Conclusions

A number of points concerning multipoint abrasive machining of ceramics have emerged

from this work. First, grinding forces were highest on ceramic workpieces with intermediate

values of hardness. Secondly, in materials exhibiting high grinding forces, the forces

measured during down grinding exceeded those measured during up grinding. The highest grinding forces measured were on hot pressed Si^N^ and the lowest were on softer ceramics such

as NiZn ferrite and MgO. Furthermore, the general features of most ground surfaces were similar. The surfaces contained regions of highly deformed material and regions where material

had been removed by brittle intergranular and transgranular fracture.

The grinding force correlations strongly suggest that initial chip formation occurs by

indentation fracture. If this is the case then the highly deformed regions on the ground

surfaces presumably occur after the leading edge of the wheel passes. This is a plausible

argument because as the wheel depth of cut decreases (ie. approaches the theoretical chip

thickness) the abrasive grains have more of a tendency to remove material by plastic flow.

This situation is enhanced by the burnishing action of swarf moving between wheel and

workpiece, high vertical grinding forces and by the increased compliance of the dynamometer.

We should point out, however, that we have observed the surfaces of ceramics ground on stiff tables to be very similar to those discussed here. Major differences are usually in

the relative amounts of plastic flow and fracture rather than in the absence of one or the

other.

Based on the correlations of grinding forces and workpiece properties discussed in

this paper, we conclude the following:

1. Chip formation during up and down surface grinding of hard ceramics occurs initially

by indentation fracture due to the impact of the abrasive grains on the wheel on the

workpiece

2. The grinding forces during up grinding correlate with workpiece properties according

to the indentation fracture theories of Lawn et al. [12,13]. In particular the

forces vary linearly with a parameter derived by Evans and Wilshaw [13] according to

F i< Kj, ^'^ H where is the fracture toughness and H is the hardness.

3. Grinding forces during down grinding differ from those during up grinding because the

chips formed at the leading edge of the wheel-workpiece interface must travel between

the wheel and workpiece before being ejected from the system. This results in an increased vertical force that can be accounted for in different ceramics by considering

the comminution characteristics of each material.

4. In every case the ground surfaces contained regions that were plastically flowed and

regions where material had been removed by brittle fracture. It is suggested that a large amount of deformation occurred after the leading edge of the wheel had passed.

87
5. Because of the complexity of the cutting point geometry on a grinding wheel, grinding

force measurements can only be used in a qualitative manner to examine mechanisms of

material removal. Possibly a better and more direct approach is to measure the rates

of material removal on different materials at constant force settings.

6 . Acknowledgements

The authors are indebted to D. Woodward for carrying out the experiments during the course of this work. The continued interest of Dr. A. Diness, Office of Naval Research is

gratefully acknowledged.

7. References

1. Ceramic Processing , Publication 1576 (National Academy of Sciences, Washington, D.C.,

1968) p. 145.

2. Gielisse, P.J. and Stanislao, J., Mechanical methods of surface finishing, in The

Science of Ceramic Machining and Surface Finishing , S.J. Schneider and R.W. Rice, eds

.

pp. 5-35 (Special Publication 348, National Bureau of Standards, Washington, D.C. , 1972)

3. Imanaka, 0., Fugino, S. and Mineta, S., direct observation of material removal

process during grinding of ceramics by micro-flash technique, ibid pp. 37-43.

4. The Science of Ceramic Machining and Surface Finishing , S.J. Schneider and R.W. Rice,

eds. (Special Publication 348, National Bureau of Standards, Washington, D.C, 1972).

5. Koepke, B.G. and Stokes, R. J., Grinding forces and the machining of magnesium oxide

crystals, Journ. Mater. Sci. , [3], 485-493 (1972).

6. Koepke, B.G., An assessment of surface and subsurface damage introduced in ceramics

by semifinish grinding operations, pp. 317-332 in Reference 4.

7. Hockey, B. J. , Observations by transmission election microscopy on the subsurface

damage produced in aluminum oxide by mechanical polishing and grinding, Proc. Brit.

Ceram. Soc. , 20, 95-115 (1972).

i 8. Becher, P.F., Ceramic machining and surface finishing, in Treatise on Materials

Science and Technology, Vol. 9. Ceramic Fabrication Processes , F.F.Y. Wang ed.

,

pp. 217-226. (Academic Press, New York, N.Y., 1976)

9. Broese van Groenou, A., Maan, N. and Veldkamp, J.D.B., Scratching experiments on various ceramic materials, Phillips Res. Rep., 30_, [5], 320-359 (1975)

10. Broese van Groenou, A. and Veldkamp, J.D.B., Scratching and grinding parameters of

various ferrites, Journ. Phys . (Paris), Colloq. [l] 285-289 (1977)

11. Gielisse, P.J., Kim, T.J. and Choudry, A., An experimental investigation of the

dynamic and thermal characteristics of the ceramic stock removal process, in Materials

Science and Research, Vol. 7. Surfaces and Interfaces of Glass and Ceramics , V.D.

Frechette, W.C. LaCourse and V.L. Burdick eds. pp. 137-148 (Plenum Press, New York,

N.Y. 1974).

12. Lawn, B. and Wilshaw, R. , Indentation fracture: principles and applications, Journ.

Mater. Sci., 10_[6], 1049-1081 (1975)

13. Evans, A.G. and Wilshaw, T.R. , Quasi-static solid particle damage in brittle solidsI. observations, analysis and implications, Acta Met,, 24_, [10] 939-956 (1976)

14. Rice, R.W. and Speronello, B.K. , Journ. Amer. Ceram. Soc, 59_, [7-8], 330-333 (1976)

15. Kirchner, H.P., Gruver, R.M. and Richard, D. , Fragment formation and damage penetration during abrasive machining of ceramics, these proceedings.

88
16. Lawn, B.R. and Marshall, D.B., Contact fracture resistance of physically and

chemically tempered glass plates, a theoretical model, Phys. Chem. Glasses, J_8, [l],

7-18 (1977)

17. Backer, W.R. , Marshall, E.R. and Shaw, M.C., The size effect in metal cutting.

Trans. Amer. Soc. Mech. Eng., 74_, [l], 61-72 (1952)

18. Cook, N.H. , Loewen, E.G. and Shaw, M.C., Machine - tool dynamometers: a current

appraisal. Amer. Mach. , 9J3, [10], 125-129 (1954)

19. Lange, F.F., Relation between strength, fracture energy and microstructure of hotpressed Si N Journ. Amer. Ceram. Soc, 56_, [10], 518-522 (1973)

20. Lange, F.F., Strong, high-temperature ceramics, in Annual Review of Materials

Science, Vol. 4 , R.A. Huggins , R.H. Bube and R.W. Roberts, eds. pp. 365-390

(Annual Reviews Inc., Palo Alto, CA, 1974)

21. Clarke, F.J. P. , Tattersall, H.G, and Tappin, G. , Toughness of ceramics and their

work of fracture, Proc. Brit. Ceram. Soc, [6], 163-172 (1966)

Discussion

CHOUDRY

There is a geometrical difference between the undeformed chips in up and down grinding.

Have your results , showing the differences between the two modes (up and down) , been already

corrected for this difference, i.e., is the difference between up and down grinding (under

identical wheel and workpiece velocity and depth of cut) more than what is warranted by the

geometry?

KOEPKE a STOKES

The grinding forces reported in this paper have not been corrected for differences in

the undeformed chip geometry obtained during up and down grinding. We argue that when chip

formation occurs by relatively long-range cracking the actual chip shape bears little resem- blance to the theoretical shape of an undeformed chip. The formation of lateral cracks that are large with respect to the widths of scratches made by single abrasive grains during

diamond surface grinding has been clearly shown on a ceramic workpiece that was tilted

slightly with respect to the table of a grinding machine by van Groenou et al., (1). The

correlation in our paper of grinding forces measured during up grinding with the indentation fracture parameter derived by Evans and Wilshaw (2) is taken as further evidence that

a major mechanism of chip formation during grinding is the intersection of lateral vent

cracks. During down grinding these chips must travel between the wheel and workpiece before

being ejected from the system, thus the vertical grinding forces measured during down

grinding should exceed those measured during up grinding. This was found to be the case for

most of the harder ceramics examined in this work.

GIELISSE

In one of your figures you showed that you found no dependence of grinding force on hardness of the material. In subsequent slides you show that there is an apparent correlation between grinding force and a compound evaluation parameter in which this hardness is

again featured. We do note further that there still remains the need to differentiate

between up grinding and down grinding and subsequently between Si 3^ and "other ceramics."

Do your researches indicate that evaluation of the grinding process in terms of derived

quantities such as hardness carry physical significance, i.e., are fundamentally justifiable

89
KOEPKE § STOKES

The physical significance implied by a correlation of grinding forces and the indentation fracture parameters of Lawn, Evans, and Wilshaw (2,3) relates to the mechanism of

material removal during chip formation. Correlation of grinding forces measured during up

grinding with the parameter derived by Evans and Wilshaw (2) argues that chips form, at

least in part, by the intersection of lateral vent cracks formed as a result of the plastic

indentation of the workpiece by adjacent abrasive grains on the wheel surface. The differ- ence between grinding forces measured during down grinding on silicon nitride and the other

materials emphasizes that during down grinding the comminution characteristics of the chips

contribute to the measured grinding forces.

MALKIN

Could you please explain your conclusion that deformation occurs after the leading edge

of the wheel has passed? Do you mean to imply that plastic deformation during abrasive-work

piece interaction occurs after chip formation by fracture? If so, this is contrary to our observations on grinding of glass, as seen in our paper in this conference.

KOEPKE g STOKES

The correlations shown in this paper, the results of Evans C4) and Evans and Wilshaw (2)

and the particularly revealing photographs of van Groenau et al., pi) and Veldkamp et al.,

(5) , all argue that chip formation during abrasive machining of ceramics occurs at least in

part by the interaction of lateral cracks resulting from the elastic-plastic indentation of

the workpiece by an abrasive grain. Thus, the chip size at the leading edge of the wheelworkpiece interface is greater than the theoretical plastic-chip size. At the trailing edge

of the wheel -workpiece interface, however, the wheel depth of cut approaches the grit depth

of cut and chip formation by plastic deformation becomes more realistic. During down

grinding plastic flow in this region is further enhanced by the burnishing action of the

swarf moving between the wheel and workpiece (6) . Thus we suggest that a large amount of

the observed deformation occurs after the leading edge of the wheel has passed.

LEE

I found the good correlation between some of the monitored parameters and the hardness

of materials Dr. Koepke and also the previous speaker Dr. Evans reported to be very interesting. However, I am troubled somewhat by the term hardness used. In view of the complexity

of determining hardness of a brittle solid, I will appreciate it if the speaker could

clarify the point further.

KOEPKE S STOKES

The hardness values used in this work were, as noted in Table 3, measured with a Knoop

indentor under a 100 gra load following accepted procedures (7) . We agree that the "hardness" of a brittle solid is a complex quantity but suggest that it is measurable and is

related to the elastic and plastic mechanical properties of the workpiece (8,9). This

point is discussed in detail by Rice and Evans in this conference (see p. 185)

90
8. Appendix

Horizontal, F^/b, and vertical, Fy/b, grinding forces measured on a number of ceramics

during down (wheel moving with workpiece travel) and up (wheel moving against workpiece

travel) surface grinding with a 100 grit diamond wheel as a function of feed rate, b ±s

the width of cut. The workpiece materials are listed in order of their hardness.

Material Depth of

Cut

Feed Rate Down

Grinding

Up

Grinding

(mm) (cm/ sec)

b b

(gm/cm) (gm/cm)

_H _V

b b

(gm/cm) (gm/cm)

PZT

MgO

NiZN Ferrite

(13ym g.s.)

NiZN Ferrite

(3ym g.s.)

"B" Si.N.

3 4

Norton Si„N.

3 4

"a" Si 0 N.

3 4

Lucalox

AVCO A1

2

0

3

Coors A1

203

B

4

C

025 042 12 74 17 72

.21 25 J153I -) 40 165

.42 42 220 59 236

.025 .042 6 79 X11X 93

. 21 24 208 —51' X 26 5 42 44 74/ '-r 31_J JL 2

.025 .042 6 14 6 14

.21 11 31~> X 1 6 38

.42 15 49 23 52

.025 .042 8 27 9 30

.21 14 66 20 74

.42 16 88 27 101

.025 .21 306 895 183 551

.012 .042 31 106 34 106

.21 87 311 84 271

. 42 128 429 125 399

.025 .21 574 1883 312 1046

.012 .042 56 361 100 3 70

. 21 84 49 3 1 26 505

.42 113 562 153 659

.025 .21 409 1 362 254 9 34

.012 .042 26 124 30 103

• £*?1X 80 J " o xU X .: J - .42 114 590 143 408

.025 .042 15 120 25 139

.21 185 654 145 431

.42 176 620 172 498

.025 .042 299 1184 249 775

.21 303 1167 397 1071

.42 293 1159 404 1156

.025 .042 91 413 136 315

.21 132 605 220 762

.42 227 1005 311 1068

.025 .042 73 523 43 154

.21 129 740 386 942

.42 152 838 415 1257

91
Re

R.

Pr

in'

KOI

ffl

H'

ini

lai

iki

sti

lOi

fai

in

pre

•«

sti

1*1

n

isi

p'i

tlf

Us

sti
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

GLASS GRINDING MECHANISMS AND INFLUENCE

ON THE FINISHED SURFACE AND STRENGTH

S. Mai kin

Faculty of Mechanical Engineering

Technion-Israel Institute of Technoloqy

Haifa, Israel

M. Huerta

Sandia Laboratories

Albuquerque, New Mexico

Mechanisms for grindinq of glass were investigated and re- lated to the finished surface morphology and fracture strength.

The abrasive-glass interaction was found to involve flow into

chiDS for grindinq with silicon carbide abrasive, and flow followed by concoidal fracture for grinding with diamond abrasive.

Much less energy was reouired for grindinq with diamond, since

much less of the material being removed is deformed. Grinding

conditions causing more flow were generally found to give less

surface damage and higher fracture strengths after grinding.

Key Words: Glass, grindinq; strength; surface damage.

1. Introduction

Machining of glass is generally performed by grinding. Grinding is a complex process

involving the interaction of numerous abrasive cutting points on the wheel surface with the

workpiece material. For metals, material removal by qrinding is known to occur by shear deformation into chips. For qlass the situation appears to be more complex due to the Possibility of both irreversible flow of the glass and brittle fracture. Although a number of

investigations imply glass removal during grindinq by fracture [1-6], flow can also occur at

large hydrostatic compressive stress conditions [7] such as developed in the vicinity of the

abrasive cutting edges. Such hydrostatical ly induced flow is obtained by indentation and by

scratching of glass surfaces, the flow followed by localized fracture at sufficiently large

loads [8-12].

An important conseouence of the grinding process for glass is the condition of the surface which is produced. The strength of glass is sensitive to the surface condition which,

in turn, is a direct consequence of the mechanisms of abrasive-glass interactions during

processing. One illustration of the influence of processing conditions on fracture strength

was provided by Shand [13] who found for sandblasting of glass rods a decrease in fracture

strength for higher blasting pressure. The parameters of the grinding process can also be

expected to have a significant influence on the strength of glass.

The present paper reviews the results of an investiqation into the mechanisms of material

removal for grinding of glass and their influence on the finished materials [14, 15]. Mechanisms of the qrindinq process were assessed from measurements of the grindinq forces and energies for a number of glasses over a wide range of operating conditions, from the nature of

the particles removed by grinding, and from examination of individual grinding scratches.

These findings were subsequently related to the ground surface morphology and fracture

strength after grinding.

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

93
2. Experimental

Plunge grinding was performed on a straight surface grinder as illustrated in figure 1.

During grinding, the force components F. and F were measured using a semiconductor strain

gaae dynamometer. A fundamental parameter obtained from the force measurements is the sped

fic grinding energy u, which is the energy expended per unit volume of material removed:

bvd (1)

Figure 1. Illustration of plunge surface grindi

Any proposed mechanism of material removal must be consistent with the magnitude of the

cific grinding energy and its dependence on the grinding parameters.

Table 1. Glasses and their Compositions

Weight Percents

speGLASS* Si0

2

A1 2

0

3

B2°3 CaO BaO BaN0

3 MgO K

20 Na

20 PbO ZnO Ti0

2 LiO a T

s

8429 Borosilicate 72 1 14 13 60 834

8390 Lead 30 2 65 3 77 583

1723 Aluminosilicate 57 15 5 10 3 3 7 46 908

7906 High Silica 96 7.5 1500

7913 High Silica 96 7.5 1500

9616 Aluminosilicate 67 21 2 1 5 4 36 919

8871 Potash Lead 35 7 58 102 525

0329 Aluminosilicate 65 20 2 Plus4Pb0- Pb

304

2 2 5 4 39 880

8463 High Lead 5 3 10 78 83 377

Soda Lime 70 10 4 13 92 710

9617 Pyroceram Same as parent glass 9616 12

0330 Pyroceram Same as parent glass 0329 12

Numbers refer to Corning glass code. - , a = coefficient of thermal expansion, x 10 C

T = sotening temperature, °C

The numerous workpiece materials used in this investigation are listed in table 1 together with the manufacturer's identification code. The specimens were typically 50mm long and

6mm wide. Grinding experiments were performed with one type of vitreous bonded silicon car94
bide wheel (37C 801-K8V) of Sin (200 mm) nominal diameter and four different metal bonded

diamond wheels (MD100-N50M, MD100-N100M, MD180-N50M, and MD180-N100M) of 7in (175mm) diameter giving two grain sizes (100 and 180 mesh) and two diamond concentrations (50 and 100).

The silicon carbide wheel was dressed prior to each test with a pyramidal diamond tool taking a radial depth per pass of 0.001 in. (25.4vjm) and a dressing lead across the wheel of

70ym per wheel revolution. The diamond wheels trued using a silicon carbide dressing wheel

(37C60MV) mounted in a brake-controlled device at a radial depth of from 25ym to 38ym and a crossfeed of 80ym per wheel revolution. Grinding fluid applied during grinding consisted of

a commercial soluble oil (Cimcool S2) mixed with water in a concentration of 1:25.

Scratch tests were run using the method developed by Thompson and Malkin [16] to observe

the grinding action by individual cutting points against a glass workpiece. This method consists of taking a single grinding pass across the workpiece with a relatively slow wheelspeed

(V = 4.1 m/s) and fast workspeed (v = 46 cm/s) so that the scratches generated on the workpiece do not overlap.

Fracture strength after grinding was measured for 1723 aluminosil icate and soda lime

glass specimens using the four-point loading arrangement illustrated in figure 2 with the

ground edge of the specimen in tension. Prior to testing, ten grinding passes were taken

along each specimen at the given test condition except for the largest depth of cut (d = 51

urn) in which case only six passes were taken.

P

1 68 mm

_P

2

0 GROUND SURFACE

^ 31 mm WJ

2

_P_

2

E

to

6 mm

Figure 2. Four-point loading arrangement for measuring fracture strength.

3. Grinding Tests

Numerous grinding tests were performed on the various glasses and a few glass-ceramics

with the different grinding wheels over a wide range of operating conditions. The results

to be presented here illustrate the main effects which were found.

A comparison of the specific grinding energies for the various glasses and a few glassceramics is shown as a plot of specific energy versus softening temperature for a silicon

carbide wheel in figure 3a and for a diamond wheel in figure 3b. Both series of tests using

the different wheels were run at the same wheel velocity, workpiece velocity, and wheel

depth of cut. Similar behavior was found for both wheels with the specific energy tending

to be bigger for higher softening temperatures. However, the specific energy for grinding

with the diamond wheel was approximately an order of magnitude less than for grinding with

the silicon carbide wheel, as the ordinate in figure 3a is a factor of ten bigger than in

figure 3b.

The effects of wheel depth of cut and workpiece velocity on specific grinding energy is

shown in figures 4 and 5. Increasing the removal rate by raising either of these parameters

generally reduces the specific energy for both the silicon carbide and diamond wheels in a similar way, although this effect is much more sensitive to wheel velocity.

95
80

60

40

EE\->

>-o

<EUJz

111 y 2

U.o

UJ

0.w

WHEEL:

•J 1 VOWI l\OV

• 9617 (coon* grain) •9616

V = 18.6 m/s • 1723 -

v = 7.6 cm/s

d = 12.7 fxm • 0329 • 0330

7906 O 8360 •

• 9617 (tint grain) SODA LIME • • 8429 — • 8871

• 8463

1 1 1

1 1 1 1 1 0 200 400 600 800 1000 1200 1400 1600

GLASS SOFTENING TEMPERATURE, T

g

(°C)

(a)

n

E 6

E

> 4

o

or

9 2

a.v>

WHEEL'. MDI80-NI00M

V«I8.6 m/»

w«7.6 cm/s

d »l2.7^tm

8390 , SODA LIME

09617 (fine grain)

•9617 (coarse grain) •>96I6 S •0329 •8S7I • 8429

7906

•8463

± _L

200 400 600 800 1000 1200 1400 1600

GLASS SOFTENING TEMPERATURE , T$ CC)

(b)

Figure 3. Specific grinding energy versus softening temperature for (a) silicon carbide

wheel and (b) diamond wheel.

20-

I '5

>-o

pq 10

o

u.o

u

0-

8463 HIGH LEAD GLASS

WHEEL". 37C80I-K8V

V = I58 m/s

v=7.6 cm/s

-O20 40

DEPTH OF CUT.d (/im)

(a)

60

UJz

oUJ

a.01

O SODA LIME GLASS

O 7913 HIGH SILICA GLASS

- A 8463 HIGH LEAD GLASS

V=I8.6 m/s

v= 7.6 cm/s

^°^~o o -a—

-o—

WHEEL: MDI80-NI00M

^ . A

a

20 40 60 80

DEPTH OF CUT, d (p)

(b)

100

Figure 4. Snecific grinding energy versus wheel depth of cut for (a) silicon carbide wheel

and (b) diamond wheel

.

96
20-

to£

E 15-

UJz

UJo

u.oUJ

0.V)

10

8463 HIGH LEAD GLASS

WHEEL' 37C80I-K8V

V = 29 m/s

d=000l in (25 4 ^.m)

-O-

_L

EE

^3

o

cr.

o

u.a

LU

Ql

10 15 20 25 30 35

7906 HIGH SILICA GLASS

WHEEL: MDI80- NI00M

A d = .001 in. (25.4/j.m)

O d = .002 in. (50.8/a m)

D d = .003 in. (76 2/j.m)

W0RKPIECE VELOCITY, v (cm/s)

(a)

5 10 15 20 25 30

W0RKPIECE VELOCITY, v (cm/s)

(b)

35

Figure 5. Specific grinding energy versus workpiece velocity for (a) silicon carbide wheel

and (b) diamond wheel.

The influence of diamond wheel composition is shown in figure 6 where the specific en- engy is plotted versus the number of diamond grains per unit area of wheel surface measured

by counting exposed grains with the aid of a microscoDe. These results imply that the in- crease in specific energy observed with finer diamond grains and higher concentration can be

attributed to having more grains on the diamond wheel surface. The relative influence of

diamond grain size and concentration can be seen in terms of their influence on the number

of grains per unit area of a cube plane for diamond grains arranged with simple cubic packing. Approximating the diamond grains as spheres of diameter a, and noting that the concentration number M in the wheel specification is four times the volume percentage of diamond,

the number of grains per unit area of cube face can be shown to be [141:

C

= 0.022 M2/3

(2)

E

E

o

crUJ

± I

SODA LIME GLASS

V = 18.6 m/s

v = 7.6 cm/s

d = 51 /i.m

Figure 6. Specific grinding energy versus number

of diamond grains per unit area on the

wheel surface.

10 15 20 25

GRAINS /mm, C

97
Non-cubic packing or non-spherical diamond particles would alter the value of the constant

in eq. (2). The relative influence of grain size and concentration on the measured values

of C was found to correlate closely with eq. (2) although the calculated values were approx

imately twice the measured ones with a = 210pm for the wheels with 100 mesh diamond and a = 80ym for those with 180 mesh diamond.

In order to provide some additional information as to the grinding mechanisms, particle

removed during grinding were collected and examined by scanning electron microscopy (SEM).

Two examples of what was found are shown in figure 7. When grinding with a silicon carbide

wheel, the grinding particles consisted mostly of flow-type chips as in figure 7a whereas

the diamond grinding tended to give mainly brittle fracture particles as in figure 7b. The

striking difference in the nature of these particles would suggest that the much higher en- ergies required for grinding with silicon carbide than with diamond (see fig. 3) is due to

material removal by flow rather than by fracture.

Figure 7. SEM micrographs of grinding detritus: (a) silicon carbide wheel, 8463 glass,

V = 18.6 m/s, v = 7.6 cm/s, d = 63.5ym; (b) diamond grinding wheel MD180-N50M,

soda lime glass, V = 18.6 m/s, v = 7.6 cm/s, d = 12.7ym.

4. Scratch Tests

Grinding scratches were generated on soda lime glass specimens by taking a single pass

with a slow wheel velocity and fast workpiece velocity, as described above. Optical micrographs of typical scratch patterns produced in this way are shown in figure 8a for a silicon

carbide wheel and in figure 8b for a diamond wheel. Scratch lengths were found to be more uniformly distributed for the silicon carbide than for diamond wheels, which indicates more uniform radial distribution of cutting points on the silicon carbide wheel surface. With

both types of abrasive, the scratches appear to have been formed by flow of the glass, which

is followed by concoidal fracture for the diamond wheel mainly in the longer scratches made

by grains protruding further out from the wheel and penetrating deeper into the glass. This

would imply that initial deformation of the glass by diamond grains is followed by concoidal

fracture if the penetration is sufficiently deep. No concoidal fractures were observed with

the silicon carbide wheel. Concoidal fractures aopear to be caused by the formation of lateral vent cracks similar to those observed by Lawn and Swain [12] for indentation. Numerous measurements were made on individual scratches to determine their cross-sectional shape

using the method previously developed [16]. Both silicon carbide and diamond wheels were

found to produce very similar trapezoidal scratch cross sections with the width at the bottom of the scratch typically about lym and included angle between the scratch sides of about

150 deg. SEM examination of scratches for silicon carbide and diamond grinding in figure 9

also reveal regularly spaced crevices or cracks extending down below the scratches normal

to the grinding direction.

98
0 100 Mm 0 100

(a) (b)

Figure 8. Optical micrographs of scratches for (a) silicon carbide wheel and (b) diamond

wheel MD180-N50M. The arrows indicate grinding direction.

0

i 50 0 50

_J

(a) (b)

Figure 9. SEM micrographs of scratches for (a) silicon carbide wheel and (b) diamond

wheel MD180-N50M.

5. Surface Morphology

SEM microaraDhs of finished surfaces are shown in figure 10 for silicon carbide grinding

of soda lime and 7906 hiah silica qlasses using a slow workpiece velocity and fine wheel

depth of cut. In these microoranhs and the others to follow, the grinding direction is from

too to bottom. The surfaces for these two different glasses aooear very similar, with striations running along the grinding direction indicative of extensive plastic flow. This flow

is evident in figure 10b which shows a oortion of the high silica glass surface at higher

maanification. Grinding of other glasses under the same conditions gave similar surfaces.

99
(a) (b)

Figure 10. SEM micrographs of surfaces ground with silicon carbide wheel 6 with V = 18.6

m/s, v = 7.6 cm/s, and d = 12.7pm: (a) soda lime glass; (b) 7906 high silica

glass.

Additional SEM observations of surfaces ground with the silicon carbide wheel were made

in order to observe the effect of wheel deDth of cut and workpiece velocity. Increasing the

wheel depth of cut was found to have little noticeable effect on ground surface morphology,

exceot for SDalling which was observed beyond a certain point. For 8390 leaded glass and

soda lime glass, the limiting wheel deDth of cut without soalling was 0.0015 in. (38ym). Increasing the workoiece velocity tended to qive a much rougher surface with less flow. This

can be seen in figure 11 where a 7913 high silica glass surface is shown for a much faster

workpiece velocity than in figure 10. Roth microcracks and localized regions of flow can be

seen on the surface.

Figure 11. SEM micrograph of 7913 high silica glass ground with silicon carbide wheel with

fast workpiece velocity: V = 18.6 m/s, v = 30.5 cm/s, and d = 12.7ym. Arrows

show microcracks on the surface.

100
Surfaces produced by diamond grinding appear markedly different from those ground with

silicon carbide. Comparing the diamond ground surfaces in figures 12 with those in figure 10

for silicon carbide grinding at the same operating parameters, it can be seen that grinding

with diamond results in a rougher fragmented surface with much less flow. As with silicon

carbide grinding, increasing the wheel depth of cut had little effect on the surface appearance, but increasing the workpiece velocity causes a much rougher surface as shown in figure

13. A few striations can be seen in figure 13, although most of the surface is fragmented.

(a) (b)

Figure 12. SEM micrographs of surfaces ground with MD180-M100M diamond wheel at V = 18.6 m/s

v = 7.6 ciii/'s and d = 12.7ym: (a) 1723 aluminosil icate glass; (b) soda lime glass.

Figure 13. SEM micrograph of 1723 aluminosilicate glass ground at fast workoiece velocity

with MD180-N10QM diamond wheel: V = 13.6 m/s, v = 41 cm/s, and d = 51ym.

101
The SEM micrographs in figures 12 and 13 were all for the same diamond wheel with the

finer grain size and higher concentration. Reducing the number of diamond grains on the

wheel surface by using the coarser grained wheel produced a more highly fractured surface

with little evidence of plastic flow as seen in figure 14. With the coarser grained wheel,

the fractured areas are more distinct, and some brittle fracture hackle markings can be seen

at the higher magnification in figure 14.

(a) (b)

Figure 14. SEM micrographs of soda lime glass ground with coarser grained diamond wheel:

V = 18.6m/s, V = 7.6 cm/s, d = 12.7ym, MD100-N100M wheel.

6. Fracture Strength

Fracture tests were carried out on 1723 aluminosil icate and soda lime glass specimens

for various grinding conditions. The results are summarized in table 2 which gives the mean strength a and standard deviation based on a minimum of five data points obtained for

each test condition. Relatively little scatter was found in fracture strengths for each

test condition, the standard deviation being typically about 10% of the mean.

Table 2. Fracture Strengths

Test

Condition Wheel v (cm/s) d (vim)

o (N/mm2

) 1723 aluminosil icate soda 1 ime

1 37C801-KV 7.6 13 77.2 + 7.9 91.8 + 8.8

2 MD180-N100M 7.6 13 71.0 + 8.4 65.1 + 5.0

3 MD180-N50M 7.6 13 63.9 + 4.5 70.1 + 4.8

4 MD100-N100M 7.6 13 44.2 + 2.4 46.4 + 3.7

5 MD180-N100M 7.6 51 69.3 + 6.3 76.1 + 2.3

6 MD180-N50M 7.6 51 71.6 + 3.6

7 MD100-N100M 7.6 51 42.1 + 3.2 46.0 + 4.8

8 MD180-N100M 40.6 13 64.0 + 6.9

9 MQ130-N1D0M 40.6 51 46.7 + 6.6 69.8 + 5.5

102
The fracture strength results follow a similar pattern for both glasses. The highest

fracture strengths for each material were obtained with the silicon carbide wheel (test

condition 1). The strengths for diamond grinding with the same operating parameters (test

conditions 2-4) were significantly lower. With both glasses, the lowest strengths with

these operating parameters were obtained with the coarser grained wheel, and comparable

strength reduction with the coarser grained wheel was also found with tests carried out at

bigger wheel depths of cut (test conditions 4 and 7), Raising the removal rate by taking a bigger depth of cut can be seen to have had almost no effect on fracture strength, but the

effect of the increasing removal rate by taking a faster workpiece velocity is not clear.

For the soda lime glass, the faster workpiece velocity had relatively little effect on fracture strength as seen by comparing test conditions 8 and 9 with test conditions 2 and 5,

respectively. For the aluminosil icate glass, however, a comparison of test condition 9 with

test condition 5 shows a big drop in fracture strength with a faster workpiece velocity.

7. Discussion

Grinding of glass is done with either silicon carbide or diamond wheels. It is apoarent

that the grinding process and the nature of the ground surfaces depend on which type of a- brasive is used. The grinding energy requirements for silicon carbide are much greater than

for diamond. SEM observations of the grinding detritus (fig. 7) suggest that these differences

in grinding energy can be attributed to material removal by flow-type chip formation with silicon carbide as opposed to concoidal fracture for diamond. These same differences appear

again in scratch test micrographs (fig. 8), although the observations here also show

evidence of irreversible flow with diamond prior to fracture. The SEM micrographs of the

ground surfaces showing extensive flow with silicon carbide and fracture and localized flow

with diamond are consistent with these observations.

It can therefore be concluded that grinding of glass occurs by flow-type chip formation

for silicon carbide, and a combination of flow and fracture with diamond. Although most of

the actual material removal with diamond occurs by fracture, the energy required for particle

removal by brittle fracture can be shown to be only a negligible portion of the total grinding energy [14]. Therefore, virtually all the energy for nrinding of glass, whether with

silicon carbide or diamond wheels, is expended by irreversible flow. Much more material is

deformed with silicon carbide than with diamond, which accounts for the differences in magnitude of the specific energy.

When comparing the grinding behavior of various glasses, a general trend was found of increasing specific energy for higher softening temperature both for silicon carbide and diamond grinding. This behavior would appear to be a direct consequence of the adiabatic con- strained flow situation at the abrasive grain cutting edges. At the rapid strain rates imposed during grinding, the deformation takes olace under adiabatic conditions so that all

irreversible work of deformation heats up the material being deformed without time for heat

to be conducted away. Very large strains are required during the deformation in order to

accomodate the severe flow constraint arising from the highly negative rake angles at the

abrasive grain cutting edges. Under these conditions, the energy input during the deformation process can become a limiting factor in the process. When grinding metals, for ex- ample, the shearing energy of the metal can be approximated as the energy reauired to

take the metal from ambient conditions to the melting state [17]. As the limiting energy is

approached during adiabatic deformation, the shear resistance of the material beina deformed

decreases rapidly, with a corresponding drop in the energy input rate with continued deformation. For glass, the softening temperature may approximate the relative energy limit for

adiabatic deformation, insofar as it also provides an indication of the energy input to re- duce the shear resistance (viscosity) to some very low value.

The influence of grinding conditions was observed by varying the wheel composition and

operating parameters. For diamond grinding, higher diamond concentration and finer grain

size were found to provide more grains ner unit area wheel surface. The linear increase of

specific grinding energy with the number of grains per unit area (fig. 6) suggests that each

grain, on the average, expends a certain amount of deformation energy prior to concoidal

fracture. With fewer grains on the surface, the amount of material removed by fracture

would have to be bigger, thus accounting for the more highly fractured apoearance of the

ground surface with a coarser grained wheel (fig. 14). The diamond concentration has much

less effect on grain density than grain size (see eq.2), so that its effect on ground surface

morphology is also much less.

103
The decrease in specific energy with faster workpiece velocities and bigger wheel depths

of cut is similar to what occurs when grinding metals, although the effect of depth of cut

is much less sensitive for glass. With metals, this behavior has been attributed to increased plowing and sliding energy as the removal rate is decreased [17]. This same explanation may be applicable to glass grinding with silicon carbide which gives flow-type chips

as in the case of metals. For diamond grinding, one possible explanation to account for the

variation of specific energy with v and d can be developed by returning to the idea that each

diamond grain exoends a given amount of energy per encounter with the workpiece. When either

the workpiece velocity or wheel depth or cut is increased by itself, there are proportionally

fewer diamond gain encounters with the workpiece for removing a given volume of material.

On this basis, the specific energy should vary inversely with both v and d, which is a strong

er dependence than observed especially for wheel depth of cut.

Additional insight into the effect of operating parameters for diamond grinding can be

gained by referring to the cutting geometry in the grinding zone as illustrated in figure 15.

As an abrasive grain enters the grinding zone, it initially cuts into the workpiece at an interference angle $. After initial deformation, material removal occurs by concoidal fracture. Now the size of the particles removed by concoidal fracture are typically much smaller than the length of the grinding zone, so the abrasive grain repeats this cutting event a number of times during a single traverse through the grinding zone. In figure 15 it can be seen that a faster workpiece velocity increases the initial interference angle 6 , whereas a bigger wheel depth of cut extends the length of the grinding zone but has no effect on this

angle. Assuming that the amount of material removed by each concoidal fracture increases

with g, fewer cutting events each consisting of flow followed by fracture would occur with a larger interference angle, resulting in less flow. This could account for the decrease in

specific energy and the more highly fractured surfaces obtained with a faster workpiece velocity. The slight decrease in specific energy with increased depth of cut may be due to the

progressive increase in interference angle for the successive cutting events during a single traverse of a grain through the grinding zone. Because of the longer grinding zone with bigger depths of cut, the interference angle averaged over all the cutting events be- comes slightly larger. Since the initial interference angle is independent of wheel depth

of cut, the ground surface morphology also does not depend on this parameter.

Figure 15. Illustration of influence of workpiece velocity and depth of cut on grinding

geometry.

104
The fracture strength after grinding apnears to be a direct consequence of the relative

amounts of flow and fracture of the grinding process. The highest strengths were obtained

by silicon carbide grinding where extensive surface flow during grinding produces a relatively smooth surface. The more fragmented surfaces obtained by diamond grinding apparently

provide more favorable sites for unstable crack initiation at lower stresses. Coarser grained wheels lead to more highly fragmented surfaces with less flow, and much lower strength.

Increasing the workpiece velocity had a similar effect on surface morphology, giving a lower

strength as expected for the 1723 al uminosil icate glass, but not with the soda lime glass.

The diamond concentration had little effect on surface morphology and insignificant influence on fracture strength. Both the surface morphology and fracture strength were also un- affected by wheel depth of cut.

These results have some immediate practical implications for grinding of glass. The use of silicon carbide wheels would appear to be preferrable to diamond wheels for obtaining a higher strength. However, the removal rate with silicon carbide will be limited by spalling

of the glass at larger depths of cut, and increasing the removal rate by taking a faster

workpiece velocity can be expected to have an adverse effect on fracture strengths. Silicon

carbide wheels also wear more than diamond wheels, which can make it more difficult to maintain any required geometrical form. Diamond grinding offers the possibility of taking much

larger wheel depths of cut without degradation in strength. The fracture strength is en- hanced with finer grain size and much less sensitive to concentration. Since the diamond

wheel cost deoends primarily on the diamond content, it should be possible to obtain good

fracture strength with inexpensive wheels having a fine grain size and low concentration.

Just how far one can go with large depths of cut, fine grain size, and low concentrations

for grinding with diamond remains to be seen.

The research was supported by a grant from Corning Glass Works.

References

[1] Newton, I, Opticks , 1695, from Holland, L., The Properties of Glass Surfaces , p. 16

(Chapman and Hall, London, 1964).

[2] Hooke, R., Micrographia , from Holland, L., The Properties of Glass Surfaces , p. 16

(Chapman and Hall, London, 1964).

[3] French, J.W., More notes on glass grinding, Opt. Soc. Trans. 8_, 8-48 (1917).

[4] Fielden, J.H. and Rubenstein, C, The grinding of glass bv a fixed abrasive Glass

Technology, 10, [3], 73-83 (1969).

[5] Vaidyanathan , S., and Finnie, I., Grinding of brittle solids, New Developments in

Grinding , Proceedings of the International Grinding Conference , M.C. Shaw, ed.,

pp. 813-831 (Carnegie Press, Pittsburgh, Pa, 1972).

[6] Pahlitzsch, E., Recent results of research into sawing of glass, Proceedings of the

Industrial Diamond Conference , p. 221-244 (Oxford, England, 1966).

[7] Bridqman, P.W. and Simon, I., Effects of very hi qh pressure on glass, J. Appl . Phys.,

24, [4], 405-413 (1953).

[8] Ryskewitch, E., Die plastisch Deformierbarkeit sproder Korper, Glasstechn. Ber. , 20_,

[6], 166-174 (1942).

[9] Joos, P., Uher die Mikroharte von Glass Oberflachen, Z.F. Angew, Phys. 9, 556-562 (1957)

[10] McClintock, F.A. and Argon, A., Mechanical Behavior of Materials , p. 495 (AddisonWesley, Reading, MA. , 1966). ~~ "

105
[11] Marsh, P.M. , Plastic flow in glass, Proc. Roy. Soc. Lond., 279A, 420-435 (1963).

[12] Lawn, B.R. and Swain, M.V., Microfracture beneath point indentations in brittle solids,

0. Mater. Sci., 10, 113-122 (1975).

[13] Shand, E.B., Experimental study of fracture of glass: 1, the fracture process, J. Amer.

Ceram. Soc, 37, [2], 52-60 (1954).

[14] Huerta, M. , and Malkin, S., Grinding of qlass: the mechanics of the process, 0. Eng.

Ind., Trans. ASME, 98, [2], 459-467 (1976).

[15] Huerta, M. , and Malkin, S., Grinding of glass: surface structure and fracture strength,

Trans. ASME 98, [2], 468-473 (1976)

[16] Thompson, D. , and Malkin, S., Grinding wheel topography and undeformed chip shape,

Proc . International Conf . on Production Engineering , pp. 727-732 (Japan Society of

Precision Eng., Tokyo, 1974)"

[17] Malkin, S., Specific energy and mechanisms in abrasive processes, Proceedings Third

North American Metal working Research Conf . , M.C. Shaw, ed., pp. 453-465 (Carnegie

Press, Pittsburgh, Pa., 1975T7"

Discussion

GIELISSE

You showed a difference of about one order of magnitude in specific

grinding energy between vitref ied SiC wheels (about 25 J/p.m3) and metal

banded diamond wheels(about 2.5 J/um^) grinding,, Should not a comparasion between a metal bonded SiC wheel and an metal bonded diamond

wheel with all other factors equal, have been better? In any case, why

do we have this big difference between the two wheels.

MALKIN

The experiments were carried out with vitref ied SiC and metal bonded

diamond wheels because these are the types normally used in practice. We

found significant differences in specific grinding energy, surface morphology, and fracture strength. One possible explanation to account for these

differences is related to the radial distribution of cutting points on the

wheel. With SiC, the distribution is more uniform giving fewer deep

scratches where material removal is by concoidal fracture.

106
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

THE FINE GRINDING OF GLASS AND CERAMIC MATERIAL USING CONVENTIONAL GRINDING WHEELS

E.D. Doyle and S.K. Dean

Materials Research Laboratories, Defence Science and

Technology Organization, Maribyrnong, Victoria 3032,

Australia

The fine grinding of glass and ceramic material, using standard silicon

carbide and aluminium oxide grinding wheels was studied.

Transparent glass surfaces were produced by both grinding wheels using

wheel depths of cut of less than 1 urn. Electron microscopy of the glass

grinding debris produced with the silicon carbide wheel showed that the glass was removed by a chip-formation process, leaving a finely grooved surface. On

the other hand, the aluminium oxide wheel removed glass by a ploughing process,

whereby the appearance of the ground surface on a micro-scale was relatively

rough. Surface quality of fine ground glass was found to be dependent on the

type of cutting fluid used.

The silicon carbide grinding wheel produced a smooth surface finish on an alumina workpiece. The mechanism of material removal was predominantly by chip

formation. Some wear debris was also produced as a result of a chemical reaction

between the wheel and the workpiece in the region of rubbing contact. Alumina

workpieces could not be ground with aluminium oxide grinding wheels due to high

frictional interactions.

Key Words: Abrasive machining; chip formation; fine grinding; glass; ceramics;

material removal.

1. Introduction

Optical glass components are normally ground approximately to shape and finished by a polishing operation to correct the shape and to achieve the required surface quality. Sometimes, particularly with aspheric surfaces, they are ground accurately to shape and then

polished with great care to obtain a high-quality surface without changing the shape. In both cases the polishing is time-consuming and laborious but is necessary as the surface

finish normally obtained in grinding is rough because of the formation and propagation of

cracks in the ground glass surface. The present study was, in the first instance, aimed at

providing a better understanding of the mechanism of fine grinding of glass (i.e., by

keeping the wheel depth of cut down to about 50 nm (2 p in)) with the use of conventional

grinding wheels, and thus introduce the possibility of obtaining high-quality glass

surfaces of controlled geometry.

The success of the results on glass also led to a preliminary study of the fine

grinding of ceramic materials. To facilitate presentation of the findings, the fine

grinding of glass and the fine grinding of a ceramic material are presented separately in

the following two sections.

2. Fine Grinding of Glass

2.1. Experimental

Glass workpieces were prepared by fixing steel end pieces to cylinders of borosilicate

crown glass (30 mm (1.2 in) diameter x 40 mm (1.6 in) long) with epoxy resin adhesive, and

machining centre-holes in the steel end pieces. The glass workpieces were ground on a cylindrical grinding machine (254 mm (10 in) x 610 mm (24 in)) fitted with a stepless,

107
temperature-stable, magnetostriction actuator [l]. The actuator was mounted on the upper

slide of the machine to give small, precise wheel position adjustments in addition to the

normal fine movements available by means of the existing screw mechanism on the lower slide.

To reduce stick-slip on the upper slide, strips of PTFE were cemented to the sliding

surfaces. During grinding, the workpiece was rotated on fixed centres at 7.5 rev/s and was traversed at 1.5 mm/s (0.06 in/s).

Tests were made with aluminium oxide (38A320JVG) and silicon carbide (37C180HVK) wheels

(305 mm (12 in) diameter x 19 mm (3/4 in) wide) using a wheel speed of 29 rev/s. Sharp

cutting edges were produced on the abrasive grits by utilizing a wheel dressing technique

which has previously been described for this purpose {2}. This involved taking a series of

dressing cuts with a cluster diamond dresser, each cut being half the depth of the preceding

cut. For the aluminium oxide wheel the final dressing cut was about 50 nm (2 y in) but, as described previously [2], the grits of freshly dressed wheels had smooth flat tops. These

flats were broken up by taking a 50 nm (2 y in) grinding cut on a workpiece of hardened

steel. The high quality of the surface ground on the steel (12 nm (0.5 y in) CLA)

demonstrated the good cutting condition of the grits. With silicon carbide wheels, however,

the flats could not be satisfactorily broken up by grinding steel and so, to avoid their

formation, the dressing cuts were not reduced below 250 nm (10 y in). Four different

cutting fluids were used, viz, water-based sodium nitrite, distilled water, kerosene (40 ppm

water) and an air jet.

To take an initial grinding cut, the glass workpiece surface was painted with a blue

spirit-based dye and the grinding wheel was advanced until the glass just appeared through

the dye, indicating that the wheel was just touching the workpiece. When the workpiece

arrived at the end of a traversing stroke, the wheel was advanced for a cut. The workpiece

was first trued by taking a series of cuts of diminishing size down to a wheel advance of

about 1 ym (40 y in). The wheel was then redressed as before and from 10 to 20 finishing

cuts on the workpiece were taken with wheel advances of down to 50 nm (2 y in).

When grinding with liquid cutting fluids, glass debris from the grinding process was collected for examination by transmission electron microscopy. The technique used involved

separating the debris from the cutting fluid by centri fuging, taking it up on a cellulose

acetate film softened in acetone and coating it with carbon. The sandwich of debris

between the cellulose acetate and carbon was then cut into 3 mm (1/8 in) squares, placed on copper grids and put in a still with acetone vapour which dissolved away the cellulose

acetate leaving the carbon film holding the glass debris on the copper grids. The grids were inserted into a 100 kV electron microscope for examination of the debris.

2.2. Results

Both aluminium oxide and silicon carbide wheels produced transparent surfaces (see

fig. 1), the degree of transparency varying considerably with the type of wheel and cutting

fluid used.

by wheel fine grinding (silicon carbide wheel with jet of air as cutting fluid).

108
2.2.1. Grinding With The Aluminium Oxide Wheel

Examination of the surface ground, prepared using either the sodium nitrite solution or distilled water as cutting fluid, showed it to have numerous rounded projections resulting

in a rouqh appearance on a micro-scale (fig. 2) with occasional surface pits of the type

shown in figure 3. The glass grinding debris was in the form of individual rods (fig. 4) of

similar dimensions to the rounded projectiles seen on the ground surface in figure 2. A

much smoother surface was produced by using kerosene as cutting fluid, the surface then

being similar to that shown in figure 5. the glass debris, however, was again in the form

of rods. When grinding was carried out with only a jet of air directed into the cutting

zone, the surface quality quickly deteriorated and grinding was stopped. It was apparent

from handling the workpiece that a significant temperature rise had occurred during grinding

Figure 2. Scanning electron micrograph of a ground surface produced with an aluminium

oxide wheel (water as cutting fluid),

showing the micro-scale roughness due to numerous rounded projections.

Figure 4. Transmission electron micrograph

of glass grinding debris produced with an aluminium oxide wheel (water as cutting

fluid), showing the debris to be in the

form of rod- like particles.

109

Figure 3. Scanning electron micrograph of a ground surface produced with a silicon

carbide wheel, showing the nature of the

surface pitting.

Figure 5. Scanning electron micrograph of a ground surface produced with a silicon

carbide wheel (water as cutting fluid),

showing that the surface consists of fine

grinding grooves with very small rounded

projections.
2.2.2. Grinding With The Silicon Carbide Wheel

In comparison with the surfaces obtained for the aluminium oxide wheel, surfaces ground

with the silicon carbide wheel using the same cutting fluids were of higher quality (fig. 5).

When using water or water-based cutting fluids, the surfaces consisted of fine grinding

grooves with small rounded projections and occasional pits. The glass debris was of

particular interest because the particles were in the form of segmented chips similar to

those formed in some metal removal processes (fig. 6). When kerosene was used as cutting

fluid an even better surface was obtained (fig. 7); very fine grooves were present and no rounded projections were evident. The debris particles were again in the form of segmented

chips. A similar high-quality surface was produced when an air jet was the only cutting

fluid and, in this case, no temperature rise was detected.

Figure 6. Transmission electron micrograph

of glass grinding debris produced with a silicon carbide wheel, showing a segmented

chip.

Figure 7. Scanning electron micrograph of a ground surface produced with a silicon

carbide wheel (kerosene as cutting fluid),

showing that the surface consists of very

fine grooves without any evidence of rounde

projections (cf. figs. 2 and 5).

2.2.3. Discussion

The most striking feature emanating from the present study is that, by fine grinding

with a comparatively coarse-grit grinding wheel, glass can be ground to give a clear surface

free from pits. This means that, in the manufacturing of optical components, the wheel

grinding stage can incorporate the polishing stage, thereby eliminating much of the timeconsuming effort of polishing.

The results indicate that a critical aspect of glass removal in an abrasive finishing

operation is that the normal grey surface finish, which is associated with a series of pits

formed by conchoidal fractures, gives way to a clear transparent surface as the grit depth

of cut is reduced. In the case of grinding with a silicon carbide wheel, the glass removal

process occurs by chip formation, while for the aluminium oxide wheel it seems to be

associated with a ploughing process, as evidenced by the grinding debris being in the form

of rod-like particles.

The importance of the depth of interaction between a grinding grit and a brittle solid

is highlighted by comparing the present findings with those of Huerta and Malkin [3]. They

investigated the comparatively coarse-grinding of glass, operating in a range of wheel depths

of cut from 15 ym (0.0005 in) to 65 pm (0.0025 in) using a silicon carbide wheel. Their

ground surface produced by grinding at a wheel depth of cut of 15 urn (0.0005 in) exhibited

surface fracture and was much rougher than that shown in figure 7.

The occurrence of both a chip formation and ploughing process when grinding with small

depths of cut is consistent with the recent observations made by Lawn and Swain [4] on the

micro- fracturing of brittle solids beneath point indentations. The essential features of

the indentation process may be described as follows: Initially, the point indenter induces

110
a zone of irreversible deformation about the contact point, the size of the zone increasing

with load. At some critical indenter load a crack, termed the median vent, initiates below

the contact point where the stress concentration is greatest. On initial unloading the

median vent begins to close, but not heal, and relaxation of the deformed material within

the contact zone just prior to removal of the indenter superimposes intense residual tensile

stresses upon the applied field. Sideways-extending cracks, termed lateral vents, begin to

appear. These lateral vents continue to extend and may cause chipping. It is important to

note that micro- fracturing depends on the sharpness of the indenter [5] and upon the

magnitude of the load.

Brose Van Goenou et al. [6] also showed that the load was important by demonstrating

that, when lead glass was scratched with the point of a diamond pyramid (apex angle 136°), a smooth groove was produced with a normal load of 0.1 N but considerable chipping occurred

with a normal load of 1.0 N. Grinding is more complex, but it may be described as an indentation process followed by a short translation in the direction of grinding. Clearly

the exact depth of interaction of any one grit on the wheel is unknown because of the

uncertainty of the number of grits within the cutting zone and of the elasticity of the

machine-workpiece system [2]. In this regard, fine grinding in the present study refers to

cylindrical traverse grinding with wheel in-feeds in the range of 50 nm (2 y in) to a maximum of 250 nm (10 p in).

Referring to the description of point indentation of brittle solids by Lawn and Swain

{4}, one would expect that the grinding grit indentations are within the zone of irreversible

deformation and that the short translation in the direction of grinding leads to glass

removal by ploughing or segmented chip formation. The generally inferior surface finish

obtained when using an aluminium oxide wheel can be explained by the higher friction

conditions due to the greater chemical affinity between the glass and the aluminium oxide

than between the glass and silicon carbide [7]. This explanation is further supported by

the fact that the glass workpiece suffered a marked increase in temperature when ground in

air using the aluminium oxide wheel, but no temperature rise was detected when the silicon

carbide wheel was used. The fact that silicon carbide is relatively inert to glass means

that, as each chip segment is formed, it is free to move up the rake face of the grit. However, with aluminium oxide the glass adheres to the grit preventing a chip segment from

forming. The glass is then pushed to one side, and this process can be considered similar

to a ploughing mechanism. Adhesion could explain why, with both wheels, the presence of

water as cutting fluid in place of kerosene led to a deterioration in the surface finish.

3. Fine Grinding Of A Ceramic Material

3.1. Experimental

The ceramic material chosen for the present investigation was recrystal lied alumina

because it was readily available. The workpieces were prepared by fixing hollow cylinders

of recrystallied alumina to steel rods with epoxy resin adhesive, centre holes being machined

in the ends of the steel. The workpiece dimensions were similar to the glass workpieces

(fig. 1). The grinding conditions were the same as those outlined previously for the fine

grinding of glass. For the fine grinding of alumina, however, only distilled water was used

as a grinding fluid, but the method of extraction and examination of the grinding debris was identical to that outlined above for glass.

3.2. Results

3.2.1. Fine Grinding With An Aluminium Oxide Wheel

It was evident from the dull appearance of the alumina workpiece that the aluminium

oxide wheel had failed to produce a smooth surface finish. Examination of the surface using

the SEM revealed that the fine-ground areas had a rough torn appearance, as shown in figure

8. The intercrystalline facets evident in figure 8(a) were present on the unground workpiece.

Examination of the grinding debris particles in the transmission electron microscope was uninformati ve as they were too thick to allow transmission of the electrons.

Ill
(a) (b)

Figure 8. Scanning electron micrographs of a ground surface of alumina produced with an aluminium oxide wheel showing

(a) that the surface has a rough torn appearance (arrowed). The

intercrystalline facets were present on the unground surface.

(b) detail of the rough area in (a) above at higher magnification.

3.2.2. Fine Grinding With A Silicon Carbide Wheel

The appearance of the alumina workpiece, fine-ground with a silicon carbide wheel, was much brighter and smoother than that produced with the aluminium oxide wheel. Figure 9(a)

is a scanning electron micrograph showing that the alumina grains had smooth flat tops. The

latter is shown in figure 9(b) at high magnification showing these smooth areas to be

essentially featureless at this level of magnification. The intercrystalline facets again

evident in figure 9(a) were present on the unground workpiece. Examination of the fine

grinding debris by transmission electron microscopy revealed that two distinct types of

grinding debris were produced. The first type of grinding debris, shown in figure 10(a),

Figure 9. Scanning electron micrographs of a ground surface of alumina produced with a silicon carbide wheel showing

(a) that the alumina grains have smooth flat tops (arrowed). The

intercrystalline facets were present on the unground surface.

(b) detail of the smooth area in (a) above at higher magnification.

112
(a) (b)

Figure 10. Transmission electron micrographs of alumina grinding debris produced with a silicon carbide wheel showing

(a) an alunina chip characterized by a fine lamellar structure.

The continuous ring electron diffraction pattern is indicative

of the chip having a very small grain size.

(b) wear debris the electron diffraction pattern of which is again

indicative of wear particles having a small grain size. The

diffraction pattern most closely indexed to aluminium silicon

carbide.

was in the form of a chip with a fine lamellar structure. The lamellar structure was much

finer than that observed in the glass chips, and more closely resembled the lamellar

structure of chips produced from the fine grinding of a hardened steel [2]. The diffraction

patterns obtained from the alumina chips were composed of continuous rings, the d-spacings

of which corresponded to a -A^t^. This indicates that the grainsizeof the alumina chips

was extremely small. The second type of grinding debris, shown in figure 10(b), was quite

different, there being no characteristic shape or lamellar structure. These particles also

gave fine-ringed electron diffraction patterns, the d-spacings of which indexed most closely

to aluminium silicon carbide.

3.2.3. Discussion

As with glass, it is possible to produce a smooth surface finish on alumina with a silicon carbide wheel, providing that the wheel depth of cut, and therefore the grit depth

of cut, is kept small. Again the observations of Lawn and Swain [4] appear to be relevant,

that is, that indentations of small depth produce zones of irreversible deformation rather

than fracture and, from the present study, it seems that translation in the direction of

grinding of the grinding grits at small depths of indentation can lead to the removal of

material by chip formation.

Further important aspects of the present results relate to the grinding debris produced

by fine grinding with the silicon carbide wheel. For example, the chip-like debris shown in

figure 10(a) had an extremely fine grain size, as evidenced by the continuous ring electron

diffraction pattern. Aghan and McPherson {8} also observed chip-like material of fine grain

size produced from slow speed fine abrasion of rutile. The large difference in cutting

speed between abrasion and grinding appears to have little effect on the mi crostructural

changes and the mechanism of material removal; this feature is also common to the abrasion

and fine grinding of hard metals [9]. The question on the origin of the fine-grain

structure has not been resolved beyond the suggestion {9} that it may be a recrystalli zation

process, resulting from the very high stresses encountered when the deformation is confined

to very small volumes of material.

Another important aspect of the results is the evidence of a chemical reaction between

the wheel and the workpiece. The diffraction pattern of the grinding debris shown in figure

10(b) most closely indexed to aluminium silicon carbide, which presumably is formed by

the rubbing contact between the wheel and the workpiece. This must then constitute a form

113
of grinding wheel wear. When grinding with the aluminium oxide wheel the frictional interaction would appear to be so strong that it prevents chip formation and results in a rough

torn surface finish as indicated in figure 8(b).

4. Conclusions

4.1. Fine Grinding Of Glass

(a) Clear transparent ground glass surfaces were produced by fine grinding with

standard silicon carbide and aluminium oxide grinding wheels with wheel depths of cut less

than 1 ym. The best surfaces were obtained by grinding with a silicon carbide wheel using

air or kerosene as cutting fluid.

(b) When grinding with a silicon carbide wheel the surface finish on a micro-scale

consisted of multiple fine grooves and the mechanism of glass removal was by segmented chip

formation

.

(c) When grinding with an aluminium oxide wheel the surface finish on a micro-scale

was very rough, and the presence of glass debris in the form of rod-like particles suggested

the operation of a ploughing mechanism for glass removal.

(d) With both wheels the quality of the glass surface finish was found to be critically

dependent on the type of cutting fluid used.

4.2. Fine Grinding Of Ceramic Material

(a) A smooth surface finish can be produced on alumina workpieces by fine grinding

with a conventional silicon carbide grinding wheel. The mechanism of material removal is

predominantly by chip formation. Some wear debris is also formed as a result of a chemical

reaction between the wheel and workpiece in the region of rubbing contact.

(b) Alumina workpieces could not be ground with aluminium oxide grinding wheels. It

is concluded that the high frictional interaction between the two materials impedes chip

formation and produces a rough surface finish.

5. References

{1} Dean, S.K., Cylindrical grinding to very accurate dimensions, Microtecnic. 20_, 198-200

(1966).

{2} Dean, S.K. and Doyle, E.D., Mechanisms in fine grinding, Proc. Int. Conf. Production

Engineering, Japan, Part 1, 655-658 (1974).

{3} Huerta, M. and Malkin, S., Grinding of glass: surface structure and fracture strength

Proc. General Engineering Conference, Baltimore, U.S. 1975 (AS ME Paper No 75/Prod-2,

1975).

{4} Lawn, B.R. and Swain, M.V., Microfracture beneath point indentations in brittle solids

J. Materials Science 10, 113-122 (1975).

{5} Lawn, B.R. andWilshaw, T.R., Indentation and fracture: principles and applications

ibid 10, 1049-1081 (1975).

{6} Brose Van Goenou, A., Naan, N. and Veldkamp, J.D.B., Scratch experiments on various

ceramic materials Philips Research Reports 30, 320-359 (1975).

{7} Mairlot, H., Texture of ground glass surfaces, Industrial Diamond Review, 61-66 (1972),

{8} Aghan, R.L. and McPherson, R., Mechanism of material removal during abrasion of rutile

American Ceramic Society 56, 46-47 (1973).

{9} Doyle, E.D. and Aghan, R.L., Mechanism of metal removal in the polishing and fine

grinding of hard metals, Metall. Trans. 6B_, 143-147 (1975).

114
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

W. R. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SOME OBSERVATIONS ON ENVIRONMENTAL EFFECTS WHEN DIAMOND DRILLING

G. A. Cooper

Institut CERAC SA

CH-1021+ Ecublens, Switzerland

During the forties, the Russian scientist P. A. Rehbinder presented evidence

to show that when drilling in rock, the addition of certain chemicals to the

flushing liquid had the effect of markedly increasing the rate of penetration.

He explained the effect by supposing that the rock was softened or made more brittle by the additive, and hence was more easily drilled.

Results which have been published subsequently have been contradictory and

have shown that, inter alia, the rock may be either hardened or softened by the

additive. It is significant that even though the original work dates from thirty

years ago, there has as yet been no commercial exploitation of the so-called Rehbinder effect.

The present report describes an investigation of the effects of some additives recommended in the literature on the speed of diamond drilling in marble and

granite. It is found that at the speeds of penetration common in commercial practice, none of the additives has any effect on the rate of penetration. By contrast, however, they have a very significant effect on the rate of wear of the

diamonds. It is this effect which is now thought to be at the root of Rehbinder'

observations, and it may well explain the conflicting results obtained by different groups of workers

Having established this explanation, one then asks how the environment influences the wear of the diamonds. The process is almost certainly chemical,

rather than mechanical, and related to frictional heating. Three possible effects

of the environment are then possible: l) the heat generated during cutting is

reduced by the environment, 2) the efficiency by which it is carried away is improved, or 3) the chemical nature of the additive can influence the wear of the

diamond. Work described shows that any or all of these influences can be significant .

Key Words: Rehbinder effect; drilling; diamond drilling; drilling additives; wear;

surface chemistry; rock fracture.

1. Introduction

Interest in the effect of the environment on drilling in rocks dates from the publication in the forties of work by Rehbinder and co-workers [1] showing substantial increases

in the rate of penetration of hard rocks by the addition of small quantities of common

inorganic chemicals. The results were interpreted by supposing that the additive aided the

formation of cracks by reducing the fracture surface energy of the material being drilled.

An alternative hypothesis, that the additive acts to reduce the wear of the drill bit, was

never examined, possibly because of Rehbinder' s interest in the action of liquids on metals

[2], where the phenomenon of stress-corrosion is of great importance.

There is considerable evidence that a liquid environment can influence the elasticplastic properties of a variety of rocks. Boozer et al [3], in a comprehensive series of

tests, showed that fluids which were strongly absorbed onto the rock surface tend to reduce

the fracture strength of sandstone and limestone under conditions where the confining pres115
sure was sufficient to induce plastic deformation, but not when the failure was intrinsically brittle. These results were confirmed by the work of Robinson [k] who again found

decreases in the plastic yield stress, but no effect on the failure under "brittle" conditions, on exposure of the rock to an active chemical environment. Robinson's work included

drilling tests with cone and drag bits, in which it was found that those additives which

produced increases in the drilling rate with one of the bits produced decreases with the

other, and the differences were most marked in the case of the drag bit. This tool removes rock by a shearing action, and thus can clearly be aided by a decrease in the plastic yield

stress of the rock. In the case of the cone bit, however, where the rock removal mechanism

is essentially one of crushing and chipping, one can argue that an increase in the plasticity of the rock will be unfavourable. These observations will be seen to have considerable

significance to the work of Westwood and co-workers (discussed below). In any event, it is

clear that, as Robinson himself notes, these observations are in direct conflict with Rehbinder's hypothesis, which would have predicted little influence of the environment on the

plastic yield stress, but a large effect on the brittle failure behaviour.

Selim etal. [5] investigated the effect of ethylene glycol, glycerine and an anionic

detergent on the rate of diamond drilling in quartzite, following the suggestion of Joris

and McLaren [6] that additives which lower the surface tension of the flushing fluid will

improve the cooling of the diamonds, and hence decrease the rate of wear. In addition,

such compounds will also improve the efficiency of removal of the cuttings. The results

of Selim etal. have several interesting features. A typical result from their work is shown

in Figure 1, which shows, above, the variation of penetration rate, and, below, the apparent

friction coefficient for the drilling of quartzite in water or 12% glycerine solution re- spectively. After an initial settling-in period the penetration curves show a steady de- crease in penetration rate with time, as the diamonds wear, and this is accompanied by a reduction in the "friction coefficient". Selim et al. were looking for evidence of the

"Rehbinder effect" and so, to reduce their data, they compared the instantaneous rates of

penetration after drilling to a certain depth (that at which the penetration rate in water

was half its peak value). This gave, in the case being discussed, an improvement of c/b

(see Figure l). Such an improvement would have led the authors to conclude in favour of the

"Rehbinder effect" but for the fact that the "friction coefficient" was greater in the presence of the addititive than without, and in fact calculations showed more energy to have

been consumed per drill meter in the presence of the additive than without. Unfortunately,

Selim et al. did not make the critical experiment, which was to use the same bit and drill

alternately in the presence of the additive and in pure water. This would have shown

whether the improved drilling performance was caused by softening the rock, or by the bit

retaining its sharpness better.

Similar results to those of Selim et al. were obtained by Engelmann, Unger and

Snowden [7] on surface set bits drilling in granite. Rates of penetration up to 35%

greater were obtained in some circumstances, coupled with reduced rates of wear. Again,

however, the critical experiment of taking the same bit and changing back and forth between

the additive and pure water, to distinguish between the phenomena of rock softening and

diamond wear, was not done.

(a)

3

2

en

'

T T T T T T ———12 pet Glycerine — Water

600 lb thrust

V 1,000 rpm __ . r "v 27.0 in.

"-fx*. \

J_

3 4 5 6 7 DISTANCE DRILLED, in

10

116
0.48

w .36

o

(b) S.24

oi—

i—oI—

Ll_

i i r i

1 1 r i r

12 pet Glycerine

Water

600 lb thrust

1,000 rpm

4 6 8 10

DISTANCE DRILLED, in

12 14

Figure 1. (a), penetration rate in quartzite as a function of distance drilled, and

(b) , "coefficient of friction". From Selim et al. [5]. Note: the

"coefficient of friction" was derived directly from the torque on the

shaft of the drill.

At this point, it is worth reviewing what is known of the influence of the environment

on the behaviour of ceramic substances on the atomic scale. The subject has been fully

reviewed by Westwood and Macmillan [8], and will not be discussed exhaustively here.

It is, however, beyond doubt that a surrounding environment can have a significant

effect upon the mechanical behaviour of the near surface regions of a variety of materials.

Not only are chemical environments effective, but also light and electric fields, under

certain circumstances. Of the various explanations put forward, the original surface energy argument of Rehbinder has been shown to be suspect, since in some circumstances, the

environment can cause increases in the strength. This would require that, following Rehbinder, a species which is spontaneously absorbed onto the solid surface can cause an in- crease in its surface energy. Such a proposition is thermodynamically unreasonable, and in

addition, it has been shown that for certain materials such as MgO [9, 10] and soda-lime

glass [11] certain environments which markedly influence the hardness have little or no effect on the cleavage fracture energy, at least for reasonably rapid crack growth.

An alternative explanation, due to Hanneman and Westbrook [12] that the additive acts

in some way as a lubricant (to aid in the penetration of an indenter) has been shown to be

untenable, at least in the case of calcium fluoride, since not only was the penetration of

a diamond into the material aided by the environment, but also the recovery of the deformation upon removal of the load [13]

The final explanation is that the environment directly affects the near surface flow

behaviour of the solid it contacts. This implies that for crystalline materials at least,

chemisorbtion on the crystal surface influences dislocation mobility in the crystal to a depth of a few microns or so. This explantation has been championed by Westwood and co- workers, and much data has been collected to support it. The evidence is particularly clear

in the case of non-metallic ionic solids, and includes such materials as silver bromide and

iodide, alumina, calcite, and magnesia, in a variety of environments [Ik]

It has further been shown that many of these effects may be correlated with the

£ -potential, which is a measure of the state of electric charge of the solid surface. In

particular, it is observed that dislocation mobility is a minimum when the £ -potential i s zero (fig. 2), and this minimum can be shown to coincide with a maximum in the measured

hardness of the crystal.

117
o

Q_

100

50

0

-50

10

s 5

<

- (a) ^

—

MgO

i

i i i i

(b)

- f

Edge

\

p Screw

i 1 Oi 1 i

10 12 14

PH

Figure 2. Variation with pH of (a) ^-potential and (b) dislocation mobility parameter

-80

-60

si -40

1 -20

t—UJ

l

. 0

o

Q_

i 20

40

60

80

1200

i/i

i

-z.

i

CO 1100

cn

<t

t—\ 1000

cn<c

inQ 900

a.

[akno3

) 3

] mol/l

Al(1000) for MgO in 10~ 2 aqueous NaOH. Dislocation mobility is least, i.e., the crystal

is hardest, when C = 0. From Westwood and

Macmillan [8]

.

Figure 3. Variation of ^-potential and pendulum hardness respectively for (0001) surface

of single crystal quartz in buffered and un- buffered aqueous A1(NC>3)3 solutions. From

Westwood and Macmillan [8],

Why this should be so is not, however, known in detail, since there is no theoretical

correlation between dislocation mobility and the state of surface charge of the material

being deformed. More puzzling is why the same correlation exists with soda-lime glass [8]

and more significantly in the present case, with quartz (fig. 3), since dislocations as such are not known in glass, and, to the author's knowledge (with the possible exception

of work by Payne [A. Ball, private communication]), dislocation movement in quartz has never been observed at room temperature.

Whatever the connection between the phenomena, however, there is considerable evidence

to connect the ^-potential with drilling rates. Figure k, for example, shows this correlation for alumina, and similar data have been obtained for magnesium oxide, calcite, and

glass [8, Ik] . One of the most puzzling aspects of this work is, however, the observation

that, in the three latter cases quoted above, the same material when drilled in the same environment with a tungsten-carbide-tipped twist drill gave the opposite behaviour to that

obtained when diamond drilling; an example is shown in figure 5, and the reader's attention

is drawn to the similarity of these results to those of Robinson [k] discussed earlier.

Various explanations have been put forward to account for this behaviour, generally

based upon supposed differences in the cutting mechanism for the two types of bit [k, lh , IT]. Specifically, it is claimed that the carbide spade bit operates by a ploughing action,

and hence works best when the material is soft, whereas the diamond bit operates either by

an impacting mechanism [8] or by opening tensile cracks in the wake of the diamonds as they

118
oQ_

-30

-20

-10

0

10

20

(a) Alumina

( Speculative)

•s. y —

\ / V /

Water

+67

. X 1 1 1

1 1 1 1

20

16

§ 12

0£

£ 4 0£a

lb) (1102 )

JL

12

10

Si 6

Polycrystalline MgO in n-Alcohols

o 0(200)- -Diamond Bit D(600)-Spade Bit

Toluene

Water 10 12

Figure 4. (a) ^-potential, and (b) rate of core bit drilling for alumina using toluene,

water and n-alcohol environments [15]. Note

the maximum hardness and drilling rate at zero ^-potential

atoms in the alcohol molecule

Nq is the number of carbon

Figure 5. Influence of n-alcohol environments and type of bit on efficiency of drilling polycrystalline MgO. After Westwood and

Latansion [16]

pass over the material being drilled [lU] . The diamond bit thus works at its best when the

material to be drilled is hard and brittle. Why the two tools should function in a different manner is, however, not at all clear. Even more confusing is the observation that for

those conditions where the diamond drilling of alumina is a maximum, the rate of cutting by

a diamond- impregnated wheel is a minimum [18] . Here again, it is difficult to understand

how the cutting mechanism can be so different as to generate results which are diametrically

opposed.

Westwood and his group have also investigated the effects of surfactants on the

drilling of rocks. Figure 6 shows results obtained with a diamond drill on Westerly

Granite. In this earlier work, it was consistently shown that the maximum rate of penetration was obtained at zero C -potential, and that this coincided with the conditions of

maximum hardness. At first sight, the improvements in drilling rate are very spectacular - greater than a factor of four in the present case -, but it should be emphasised that the

absolute rates of penetration obtained in these laboratory scale experiments were very low.

Taking the data of figure 6, for example, the best rate obtained was some k.Q mm/minute,

compared with that in water of 1 mm/minute. These values may be compared with the peak

values reported by Selim et al. [5] in quartzite (a harder rock) of 90 and 80 mm/minute

with 12% glycerine and pure water, respectively (see fig. l). Results obtained at CERAC

using pure water were typically in the range of 50 mm/minute at 1000 rpm, and up to 120

mm/minute at 3000 rpm (O.k'J and l.k2 metres per sec. linear cutting speed, respectively).

119
The subsequent results of Westwood and co-workers, at larger scale, will be discussed below,

and compared to data obtained by the present author.

Westerly Granite

[ DTAB ] .MOL/l

Figure 6. Variation of D(l50) (rate of drilling after 150 seconds) for Westerley granite

in aqueous DTAB (dodecy trimethyl ammonium bromide). Diamond core bit, room temperature.

After Macmillan et al. [19]

.

2. Results obtained at CERAC

Work at CERAC was confined to the area of diamond drilling, using impregnated core

bits. These drills contain diamonds distributed throughout the volume of a bronze matrix,

and in commercial operation, conditions are adjusted so that the bronze wears away at

approximately the same rate as the diamonds are blunted. Hence, due to periodic loss of the

worn diamonds, and exposure of new ones, a steady drilling performance can be obtained. In

the laboratory, however, conditions can be adjusted to produce steady-state drilling, as above, or, if the wear of the bronze matrix is reduced, conditions similar to those ex- perienced with "surface-set" bits may be obtained. ("Surface-set" bits have one layer of

larger diamonds cast into the surface of the bit. When these are worn, the bit must be

returned to the maker for re-setting the stones.)

Experiments were made at small scale, but great attention was paid to the scaling so

as to reproduce similar feed pressure per unit area of bit, and similar cutting speeds to

those used in commercial practice. Accordingly, experiments were carried out using

diamond-impregnated core drills of 10 mm O.D., x 8 mm i.d. The diamonds were of type

de Beers MDA-S, of 70-80 mesh (U.S.), present at a concentration of 19 volume percent in a bronze matrix. The drills were used in a bench-mounted drill press having a co-axial water

flushing system provided with flow and pressure-measurement systems, and a continuouslyvariable rotation speed in the range 500 to 3500 rpm. At 3000 rpm, the mean cutting speed

of the diamond in our drill was 1.U2 m/sec, which should be compared with the following

values common in large scale drilling*: surface-set bits: 1-2 m/sec, impregnated bits:

Anders Oden, Craelius Diabor, private communication.

120
2-5 m/sec, saws, 30-50 m/sec. It will be seen that our cutting speeds were somewhat lower

than those used in the field. The feed pressures recommended for impregnated bits range

from 2 to 10 N/mm2 of bit surface, depending upon whether the material is soft and abrasive

(sandstone, concrete) or hard and non-abrasive (taconite, granite). As will be seen below, we have ourselves established the optimum feed pressure required for our drill in different

rocks, and the values are found to lie within these limits.

Experiments were made on two rocks, a pure white marble from Lasa, Ticino, Switzerland,

and the well-known pink granite from Bohus, Sweden. The marble was chosen since it is

well-known that calcite can deform plastically at room temperature, and hence if dislocation effects are important to the mechanism of the Rehbinder effect, this is a likely

material in which to observe them. The granite was chosen for the opposite reason, since

dislocation movement is not known to occur in quartz at room temperature, and because

granite is itself a very widespread and commercially significant rock. The experimental

programme carried out was first to investigate the performance of the core drill in each

of these rocks in the presence of pure water alone, in order to establish the basic performance envelope of the drill, and then, having found certain "standard conditions", to

investigate the effect of a number of additives recommended in the literature.

2.1. Drilling in marble: establishment of drill performance

The variable parameters when diamond drilling are: the rotation speed, the feed force,

and the water flushing pressure. Penetration speed increases with the first two provided

that the loads do not become so great that the drill is destroyed, and provided that

adequate flushing is maintained to remove the rock debris and prevent overheating of the

drill.

When drilling in marble, it was observed that the rate of wear of both diamonds and

matrix was very low, and it therefore proved possible to carry out a large number of

experiments with the cutting capacity of the drill essentially constant.

To establish the drilling performance envelope, therefore, tests were made at rotation

speeds of 500, 1000, 2000 and 3000 rpm, and the drilling rate was measured as a function of

water flushing pressure for different feed pressures. Typical results, for a rotation

speed of 500 rpm are shown in figure J. In this figure, it will be seen that, for a given

feed force, the penetration rate decreases slowly with flushing water pressure. This is

simply because of the upthrust on the drill column exerted by the water pressure. More

interestingly, it will be seen that there is a region at low water pressures where

drilling is not possible due to choking. This occurs, particularly at high feed forces

(when there are large quantities of cuttings produced and the drill is forced deeply into

the rock), when the liquid flow is not sufficient to remove the cuttings as fast as they

are produced. Under these conditions, the viscosity of the slurry of cuttings increases,

the flow rate further decreases, and eventually the drill is completely choked. Then,

unless the feed force is immediately released and the liquid flow re-established, the heat

generated will be sufficient to melt the bit within a few seconds . The required flushing

water pressure to prevent choking increases with the feed force, as is evident from figure

7. A simple theoretical description of the choking process [20] indicates that the line

separating choked from free-flushing conditions should be of the form: P a R

3

, where P

is the minimum flushing pressure, and R is the rate of advance. The dashed line in

figure 7 has the form

P, x s = 2 • 10

_5 R

3

, . -n

(at) (mm* mm )

which appears to describe the observed data quite well.

Using the data obtained in these and other tests, it was decided that a feed force of

150 N and a flushing pressure of 3 bars would give satisfactory drilling performance in the

speed range investigated, and these were therefore adopted as the "standard conditions" for

121
0 1 2 3 4

WATER PRESSURE BAR

Figure 7. Marble: penetration rate as a function of flushing water pressure for different

feed loads at 500 rpm.

the tests with different environments.

2.2. Drilling in marble: effect of additives

Using the standard conditions at different rotation speeds in the range 1000 to 3300

rpm, tests were made with a variety of additives in aqueous solution in the flushing fluid.

These were:

- Al CI 3 0.5, 0. 75 and 1%

- Na2 C0 3 0.2, 0.33, 0.5 and 1%

- Na OH 0.025, 0.05, 0.1 and 0.25$

- "Quat 3l6" (dodecyl ethyl dimethyl ammonium bromide) 10" 3 and 10

-2 mole/

litre

- CWS liquid soap, containing 21$ salts of fatty acids and 5$ sulphonated

compounds as its active constituants , diluted to 0.25 and 0.5%.

All of these compounds have been recommended by various authors as improving the rate of

drilling in calcareous rocks.

The results obtained are shown in figure 8: mean results only are given for the additives, while a scatter band indicates the 95% confidence limits for the rates obtained with

pure water. An examination of the data shows, as might be expected, that the penetration

122
rate increases approximately linearly with drill rotation speed. Apart from this, there is no systematic trend in the data which separates one environment from any other. One must

therefore conclude that, at least for marble under our standard conditions, none of the

recommended additives causes a measurable increase in the drilling speed. In the face of

this surprisingly consistent negative result, experiments on drilling in marble were

abandoned

100

£50

Marble Lasa

Load 150 N

T

X mln.

with water 3 bars

o Ouat 10 /I • i. 10'3 O Soap .5 %

m » .25 V Na2 C03

1 V .. .5 A << .33 A .2 f NaOH .25 • • .1 O " .05

x » .025 • AICI3

1 • 11 .75 A » .5

4

4

t

A

RPM = 10 3

Figure 8. Marble: penetration rate as a function of drill rpm for various additives in

aqueous solution. "Standard conditions" of 15 kg. feed load and 3 bars flushing pressure.

From ref . [21]

2.3. Drilling in granite: establishment of drill performance

When diamond drilling in quartz- containing rocks including granite, it is observed

that there is a continuous wear of the diamonds. This effect may be more or less masked

by the simultaneous wear of the matrix material, leading to the periodic loss of worn

diamonds and the exposure of new ones on the drill surface. These three interacting wear

mechanisms are shown schematically in figure 9. The drill bit we have used for our studies

had a bronze matrix which would be considered somewhat too hard for successful commercial

drilling in granite, since when drilling with pure water as the flushing medium, the abrasive wear of the matrix did not, in general, keep up with the wear of the diamonds. Thus,

if one measured the penetration distance as a function of time over a wide range of conditions, an approximately exponential curve was obtained, with the rate of penetration of the

bit decreasing continuously to zero. Under commercial conditions, such a performance would

not be considered to be satisfactory, since the bit would have to be periodically removed

from the hole for re-sharpening (for example, by drilling for a short distance in a weak

but very abrasive material: fire brick is commonly used for this purpose). On the other

hand, such a bit was ideal for our purposes since it allowed us to separate out to some extent the phenomena of the wear of the diamonds and the wear of the matrix.

123
Figure 9- Schematic view of the drilling process, showing the three possible wear mechanisms (after Cooper, ref. [22]).

Although the penetration-time curve for our hit in granite was nearly always exponential, the asymptote (i.e. the total distance drilled) varied very greatly. Examples are shown in figure 10 for two feed forces at 1000 rpm and 3 bars flushing pressure, and these

are by no means extreme values. Depending principally upon the feed force, we have

observed total lives varying from less than five mm to more than three metres. The range

of drilling conditions we have investigated lies between the following limits: flushing

pressure 3 or 6 bars, feed forces 100-600 N, and rotation speeds 500-3000 rpm.

Figure 10. Granite: distance drilled as a function of time for two different load

conditions (from ref. [21]).

Time (s)

124
Under most circumstances, provided that the drill is not operating in "self-sharpening"

mode, the wear process produces a decrease in performance which is exponential, and a measure of drill performance may then he obtained by calculating the diamond half-lives,

and, by multiplying by the rotation speed, obtaining the total distance travelled by the

diamond over the rock in this period. Such data are presented for a variety of conditions

in figure 11, plotted against the feed force. It will be seen that the life of the diamond is strongly dependent upon the feed force, but not, apparently, upon either the rotation speed or the flushing water pressure. It would be tempting to conclude that the

relationship was approximately linear if it were not for a group of data lying in the upper

right-hand corner of the figure. These values, however, come almost exclusively from tests

in which the steady degradation of cutting performance is interspersed with sudden, irregular increases. These are almost certainly associated with the loss of worn diamonds, by

erosion of the supporting matrix, and the exposure of new, sharp diamonds.

700 i

1 1 1 . 1

I

,

600

500

£ 400

S 300

200

-rasa ft HT*

100 —»o-tr<

GRANITE B0HUS 6 bar

500 RPM

1000 " 3000 " *

3 bar

*

100 200 300 400

DISTANCE TRAVELLED BY THE DIAMONDS [M]

500

T, -RPM-tt-9

60-1000

600

Figure 11. Granite: relationship between the feed force and the distance travelled

by the diamond over the rock up to its half life.

At first sight, the observation that the life of the diamonds is less under low feed

forces is surprising. The following hypothesis is, however, proposed to explain this

phenomenon. Firstly, we note that the diamond is very much harder than the hardest phase

present in the rock being drilled (the quartz), and therefore even if the latter is supported by its surroundings, it is extremely unlikely that the diamond can be caused to flow

plastically by the stresses imposed on it by the rock.

In addition, we observe that the half life of the diamond (in distance travelled)

appears to be independent of the traversing speed, and hence of the strain rate applied to

the rock. This would argue in favour of a brittle failure mechanism in the rock which

does not depend upon thermal activation of, for example, dislocations for deformation to

occur. In consequence, it is reasonable to believe that an increase in the feed force

applied to the drill will not result in increased contact stresses at the diamond-rock

interface, since the latter is already at its failure stress and this is strain-rate independent. Instead, it is rather to be expected that the diamonds will be pressed more

deeply into the rock and that the total contact surface will increase rather than the contact stress. This indeed appears to be the case, as it has been shown that under conditions

of drilling with a sharp bit, the initial penetration distance per rotation is linearly proportional to the feed load and essentially independent of rotation speed or flushing water

125
pressure [20]. Since the total load carried by each diamond must necessarily increase with

the feed force, the internal, rather than the surface stresses will increase. In the limit,

the diamonds may be cracked, broken or torn out of the matrix at very high feed loads, but

we did not attain such conditions in our experiments because choking of the drill occurred

beforehand. In commercial practise, however, such damage can easily be obtained (Anders

Oden, Craelius Diabor, private communication).

Microscopic observations of the drill bit tend to support these arguments.

Figure 12 shows two S.E.M. photographs of diamonds in a worn bit. The wear is normally

characterised by a rounding of the sharp crystal edges, and various patterns of striae or grooves on the crystal surface, suggestive of an ablative wear process, can be seen.

Figure 12. Scanning electron micrographs of diamonds in a worn drill bit.

a) mag UoOx, b) mag TOOx.

If, as seems likely, mechanical deformation is not responsible for the wear of the

diamonds, some chemical transformation may be involved. Two possible candidates are oxidation or graphitisation, both occurring as a consequence of frictional heating of the diamonds. Oxidation would be expected to be a strictly surface phenomenon, resulting in a steady ablation of the diamond surface, and with the production of rather rounded wear

forms. In the case of graphitisation, however, it is possible that, in view of the high

thermal conductivity of diamond, graphitisation will occur not only on the surface, but

also in the interior of the diamond. This would result in a general weakening of the diamond throughout its volume (although of course most markedly in the hotter regions near the

cutting edge), and it is possible that this would lead to failure by progressive fracture

of the weakened regions. At present, it is not possible to say whether this reasoning is

correct, to distinguish whether both mechanisms are operative, or to say if one is dominant.

Some further evidence will, however, be presented below in the section on environmental

effects

.

Whatever the mechanism of the wear process, however, the microscopic evidence shows

that wear occurs principally by a progressive ablation and rounding-off of the cutting

edges of the diamond. In addition, we have seen that changing the feed force probably

results in the diamonds being forced more or less deeply into the rock, rather than changing

the contact stresses. The wear process under high or low feed forces is then envisaged to

126
occur as follows: Starting from an initial condition with a sharp diamond, and applying a high feed force, the diamond will "be forced deeply into the rock. The stress-field associated with the indentation will be large, and a large volume of rock will he fractured as the

diamond sweeps over the rock surface. As wear occurs, the cutting edge of the diamond will

he progressively worn away, until a large wear flat has been produced and the diamond

ceases to cut. Contrast the case of a bit in which the feed force is low. Here, the diamond points penetrate only a little way into the rock, and, since the feed force is low,

only a small quantity of diamond must be worn away before the surface of the wear flat is

sufficiently large to transmit the feed force to the rock without it fracturing under the

diamond. We now make use of our assumption that since, during cutting, the rock is at its

fracture stress, the surface stresses on the diamond where it contacts the rock will be

the same in the two cases. Since, in addition, we presume the wear to occur as a consequence of frictional heating at the rock-diamond interface, we suppose that the rate of

ablation of the cutting edges will also be the same in the two cases. It is thus clear

that since a greater depth of diamond must be worn away in the case of the high feed force,

the life expectancy of the drill will be correspondingly greater in these conditions

We have made one simple experiment to check whether the wear mechanism proposed is

plausible. We started with a newly-sharpened drill, and ran it at 1000 rpm, with three

bars flushing pressure and a feed force of kOO N. At this feed force, the drill had a long

life and cut well. Reference to fig . 13 shows an initial drilling rate of 60 to 90 mm per

minute in the first 200 seconds. After 6000 seconds, the rate of penetration had fallen

to 2 mm per minute. At this point, we imagined the diamonds to have been worn down very

Figure 13. Granite: experiment on drill

wear: speed of penetration as a function of

time for various drilling conditions.

60 65

TIME OF DRILLING f\0

2 tj

considerably, and to have developed large wear flats. When the feed force was subsequently

reduced to 100 N, therefore, it was not surprising to find that the rate of penetration was

only 1 mm per minute, and that this dropped rapidly to 0.02 mm per minute after a further

150 seconds. The drill was then removed from the rock and re-sharpened. It was thereafter

returned to the rock again and made to drill under a feed force of 100 N, as had been used

127
previously. Under these conditions the initial rate of penetration was 6.3 mm per minute,

which, as expected, was "better than the performance after drilling for 6000 seconds at

hOO N. On the other hand, the rate of penetration dropped below this value after only

60 seconds, and after a total of 350 seconds it had fallen to 0.2 mm per minute, at which

time the experiment was stopped. The feed force was then increased to hOO N, and drilling

was re-started without re-sharpening the bit. As will be seen from figure 13, the drilling

rate at this higher load lay between 60 and 90 mm per minute, which is almost the same as had been measured at the beginning of the experiment with the newly- sharpened drill. To

summarise this series of tests, therefore, we find that a drill which has been worn at high

feed loads has a low initial rate of penetration and thereafter blunts very rapidly when

worked at a lower load, but that a drill which has been worn at a lower load is still

capable of cutting as well as a newly-sharpened drill when the feed load is increased.

These observations are completely consistent with the wear model we have proposed.

In view of what had been learned about the behaviour of our drill in granite, we chose

as "standard conditions" to work at 1000 rpm, 3 bars flushing pressure, and a feed force of

kOO N. Because of the continued wear of the drill, it was not possible to make straightforward comparative tests of the different flushing fluids as had been done with marble, but

it was necessary to change back and forth between water and the environment in order to

allow for the effects of wear. We therefore present the results in this section as the

speed of drilling as a function of distance drilled, taking a measurement of the average

rate of penetration during each successive advance of 25 mm.

The first group of tests was concerned with the series of pure straight-chain aliphatic

alcohols belonging to the homologous series CH 3 (CH 2 ) n OH between methanol and decanol inclusive, since these have been reported by Westwood and co-workers as improving the rate of

penetration in quartz-containing rocks [19] . The first results obtained are shown in figure

Ik . Here we have taken an initially sharp drill, and have drilled with each successive

alcohol, alternating with pure water. There are several striking features in these results.

Firstly, as would be expected, we see a generally decreasing rate of penetration with time,

which presumably results from the progressive wear of the diamonds. Secondly, it is quite

2.U. Drilling in granite: effect of additives

Granite Bonus

Load 400 N

1000 RPM

1 50 n - Alcohols

3 bars \

o water
clear that, with the possible exception of methanol, the rate of wear of the diamonds is

much greater when water is used as the flushing medium, than when any of the alcohols is

used. Thirdly, there is in general no sharp change in the instantaneous rate of penetration

when changing from alcohol to water or vice-versa. Together, these observations throw an interesting light on the "Rehbinder effect" in drilling. In essence, they suggest that the

principal effect of the additive is not to increase the instantaneous rate of penetration

of the drill, but to decrease the rate of wear of the diamonds. This conclusion may go a long way towards explaining the confusing and often conflicting results which have been

previously reported in the literature. If, for example, we had conducted our experiments

by taking a sharp drill each time, and measured the time taken to drill, say, 500 mm, we would have obtained striking differences in these times, and might have been led to conclude

that, since the average penetration rates were different, then the rock had been made harder

or softer by the additive. Indeed, with hindsight, one can look back at the results of

other workers (eg. Selim et al., figure l) and find exactly the same behaviour as we have

found in the present work.

At certain points in figure lU, one can find evidence of an increase in the rate of

penetration (see, for example, results obtained with propanol, pentanol and octanol). It

is thought that these are evidence of a spontaneous re-sharpening of the drill. As was mentioned above, we believe that our "standard conditions" for drilling are not far from

those necessary for steady-state re-sharpening to occur. Admitting our hypothesis that the

additive influences the rate of wear of the diamonds rather than the resistance of the rock,

and having in mind the three wear processes active in the drill (see figure 9), we suggest

that with water, the rate of wear of the diamond is dominant, whereas with the alcohols it

is much reduced. In the former case, as the diamonds wear, the rate of production of rock

cuttings decreases, and hence, therefore, so does the abrasive wear of the matrix. This

has the consequence that a decreasing number of worn diamonds will be lost from the face of

the drill, until, finally, a steady-state condition is reached where all the diamonds in

the drill face are blunt, and both the penetration rate of the drill and the rate of re- placement of the diamonds drop to zero. When one of the alcohols is used as the flushing

agent, however, the diamonds are blunted more slowly and the rate of production of rock

cuttings is maintained. Thus the abrasive wear of the matrix is maintained, with consequent loss and renewal of the diamonds on the working face, and the rate of penetration of

the drill is maintained or can even increase.

To be sure that there was no undetected systematic error in our previous set of observations, we made another series of tests where we changed at random between the different

alcohols. The results obtained were consistent with those reported above: in no case did we find instantaneous changes in the rate of drilling when changing between water or any of

the alcohols, but a clear difference in the rates of wear of the drill [20].

We next tested a range of different chemicals which have variously been recommended as drilling additives in the literature. These included:

Sodium Hydroxide, NaOH, 0.1%, 0.3%

Sodium Chloride , NaCl, 0.1%, 0.5%

Sodium Carbonate, Na2 C0 3 , 0.1%, 0.3%

Aluminium Chloride, A1C1 3 , 0.02%, 1%

CWS liquid soap (see above), 0.5%

"Quat 316" (see above), 10-3 Mol/litre,

all in aqueous solution.

Of these various solutions, the inorganic salts showed no influence on either the

instantaneous rate of penetration of the drill, or on its rate of wear, while the soap and

the quaternary ammonium compound showed no influence on the rate of penetration, but a re- duction in the rate of wear. Results from the latter group of tests are shown in figures

15 and 16, respectively. Inspection of figure 15 shows that of the three dilutions of soap

used, a concentration of 0.2% does not appear to be very effective, but at concentrations

of either 0.5 or 2%, the rate of wear is definitely reduced. In the case of the "Quat 3l6"

129
The tests with a concentration of 10 4 Mol/

3 Mol/litre show

(fig. 16), the evidence is less clear cut.

litre show a rate of wear comparable with that of water, while those at 10" a high rate of wear in the first test, hut a rate of penetration oscillating irregularly

about TO mm/minute in the second. At a concentration of 10~ 2 Mol/litre, only one group of

tests was made, at the beginning of the experiment, immediately after re-sharpening. There

is a fairly rapid fall from 85 to near 65 mm/minute at first, but the rate of penetration

100

S 50

en

• H

2 0 SOAP CWS . 2t

• " " .5t cP

B " " 2\ •

> —ft» RESHARPEN I NG ^

o oD

cfti

o <

•• •••D •

..•'»

i .

GRANITE BOHUS

1000 RPM

<<00 N

3 bar

.5 1 1.5

DISTANCE DRILLED [m]

100

•a: * 50h •

en

• SOAP CWS .5*

GRAN IT BOHUS

1000 RPM

<«00 N

3 bar

* 0 .5 1 1.5

DISTANCE DRILLED [m]

Figure 15. Granite: rate of penetration as a function of distance drilled using water

or various concentrations of CWS liquid soap as flushing liquids.

appears to stabilise near this value before falling away again as the concentration of the

additive was reduced.

Our conclusion from the tests using either the aliphatic alcohols or various additives in aqueous solution must therefore be that certain environments indeed affect the

drilling process. On the other hand, all the evidence points to the conclusion that the

environment affects the rate of wear of the diamonds, but does not modify the properties

of the rock.

There is recent evidence from other workers to support this view. Figure IT, taken

from the work of Mills [23] shows very similar results to our own, but comparing the performance of the quaternary ammonium compound D.T.A.B. and water. Other data obtained by

the same authors are shown in figure 18, and should be compared with those of figure 1:

the similarity is quite striking, and they are clearly both evidence of the same effect.

130
1 00

c

-He

<

atz

o

<

0Lz

UJ

Q.

50 -

GRANITE BOHUS

1000 RPM

«»00 N

* 3 bar

CP

x in

« I

CD

e H

2 0

• QUAT 10 Mol/1

i 1 0

3 „

10* « t RESHARPEN INGS

1

t

. 5 1

DISTANCE DRILLED [wi]

o - Figure Id. Granite: rate of penetration

as a function of distance drilled using

different concentrations of "Quat 3l6"

in water as the flushing fluid.

1 . 5

10

£

E

<C

Westerly Granite T

(a)

l(f

3 mol/ liter X

DTAB /

i 5 10

(b)

(a)

Figure 17. (a) Penetration vs time curves for

freshly redressed, 0.6 cm o.d. diamondimpregnated bits drilling_into Westerly

granite under water or 10 3 M D.T.A.B.

2000 rpm. 40 N thrust. (b) Similar data for one freshly dressed bit. Note, however,

that on changing from water to the 10 3 M D.T.A.B. environment, no significent

change in slope occurs. After Mills (Ref. 21)

(b)

TIME ! sec )

131
Westerly Granite

PENETRATION (cm)

Figure 18. Penetration rate vs penetration distance data for Westerly granite drilled with

a 0.95 cm o.d. diamond-impregnated core bit under a) 10

-3 M aqueous D.T.A.B. and

b) water. The thrust was 60 N, and the "bit rotation speed 5000 rpm. From Mills [21].

2.5. The wear process and the effect of the additive

In view of the evidence presented above, it can be accepted that, at least for rates

of penetration which are of interest commercially (some tens of millimetres per minute),

the dominant influence of the additive is in reducing the rate of wear of the diamonds.

Why this should be so, however, remains to be answered.

We have discussed above the different possible wear mechanisms of the diamond, and

have arrived at the tentative conclusion that wear occurs by physico-chemical rather than

purely mechanical means. Of the different possibilities, either graphitisation or oxidation

are likely candidates, and either will be aided by the heat generated during drilling.

Mills and Westwood [2U] have argued that the quantity of heat generated, and hence the wear of the diamonds, depends upon the coefficient of friction between the two materials,

and that this in turn is influenced by the environment, since the latter acts to harden or soften the rock by the mechanisms discussed in section 1. An essential question then must

be answered before this hypothesis can be accepted: If the rate of wear is affected

principally by changes in the hardness of the rock, why then does the instantaneous rate

of penetration not change also? In none of our experiments have we found differences in

the rate of penetration which could be attributable to changes in the hardness of the rock:

this was particularly disappointing in the case of marble, since dislocation motion in

calcite is easily possible at room temperature, and the influence of the environment on

its plasticity has been documented previously [3, ^, 15].

In the present state of uncertainty, however, it would appear unwise to exclude a

priori-, any possible mechanism, and the following would seem to be worth consideration:

l) The quantity of heat generated at the cutting points during drilling is affected by

the environment. The cause could be related to the chemomechanical effect, lubrication,

or changes in the flow characteristics of the slurry of rock debris in the flushing

fluid.

132
2) The efficiency of heat removal from the diamonds is changed; the effectiveness of heat

transfer from hot diamond to the flushing fluid could be affected, either directly,

or through the rheological properties of the rock debris slurry.

3) The environment has a direct chemical effect (corrosive or protective) upon the

degradation of the diamond.

10 •

0

0.1 0.2 0.3

Flow 1/min

Figure 19. Granite: apparent power dissipated during drilling at different rates, in

the presence or absence of an additive.

We have made some experiments to try to clarify the situation, as follows: In the

first experiment, we have used a differential thermocouple to measure the heat generated

during drilling in the presence or absence of the soap used in previous tests (fig. 19). In

such work, it is important to eliminate spurious effects such as the heat generated from

the turbulent stirring of the flushing liquid [25], and these deductions have been made in

preparing figure 19. Several observations may be made from this data:

1) More power is dissipated at higher rates of penetration than at low. This is reasonable,

in view of the greater quantity of rock which is broken.

2) The addition of soap produces a lower power dissipation than water alone under all

drilling conditions. This must be seen as direct confirmation of the first hypothesis

above, since if less power is dissipated, the diamonds will in all probability run

cooler. Whether the effect is due to the hardening or softening of the rock, or is

attributable to a lubricative effect, however, cannot be deduced from the present data.

3) The power levels lie in the range 20-50 watts, for a 10 mm O.D. drill bit, having a projected cutting surface area of about 28 mm

2

. The concentration of diamonds in the

bit was 19 V/ Q

. If we suppose the heat to have been dissipated into the liquid over

133
the exposed surface of the diamonds, we obtain values from 20/(28 x 0.19) = 3-8 MW m

-2

to 50/(28 x 0.19) = 9.^ MW m

-2

. This is not an exact calculation: some of the diamond

will he in intimate contact with the rock, and will not be well cooled, and some heat

will be conducted away through the supporting metal of the bit. Nonetheless these are

extremely, high power levels to be carried over a solid-liquid interface. We suggest,

therefore, that under these conditions, the temperature of the diamonds may easily

exceed 100°C, and thus any discussion of the cooling process must include the possibility

of boiling heat transfer. (indeed, it is commonly observed that if one drills within

a few millimetres of the free surface of a translucent rock such as quartzite or granite,

red and orange streaks can be seen which are thought to be individual diamonds at red

heat .

)

We would agree with the opinion of Mills and Westwood [2h~\ that the presence of small

concentrations of additives will not significantly affect thermal transport properties such

as the specific heat or conductivity of the flushing fluid, but a possible role could

easily be for the additive to influence the transition from nucleate to film boiling. Such

an effect would not require the additive to be present in large quantities: Rhodes and

Bridges [26] boiled water on the surface of a heated tube. They found that film boiling

can be aided by making the surface less wettable, as for example by coating it with wax or by adding very small quantities of oleic acid or mineral oil to the water. On the other

hand, inorganic salts such as sodium carbonate or chloride tended to promote nucleate

boiling with correspondingly increased heat transfer. In our experiments, we have found

soaps to be effective as drilling additives, and these are, of course, designed to have

good wetting properties. Rehbinder and other workers have claimed that simple inorganic

salts are effective also [l], and the quaternary ammonium compounds used by ourselves and

Westwood and co-workers are well-known mineral-dressing additives. It would seem, therefore, that this hypothesis is worth testing further. Finally, the influence of the additive on the rheological properties of the suspension of rock cuttings should not be forgotten, and it could be that if the surfactant significantly changes the viscosity of the

slurry around the bit, the cooling could be much improved. There is some evidence to

support this view; Clarke [27] has investigated the rheology of suspensions of glass and

quartz particles of various shapes in the size range of 20 to 200 u, following the work of

Schack et al. [28] and Moreland [29], and has demonstrated that at concentrations as low as 0.6% a common anionic detergent can reduce the effective viscosity of a 30% solids quartz

suspension by 50%. Further indirect evidence is available in the work of Haydon [30] who

investigated the effects of several additives recommended by Rehbinder and co-workers on the efficiency of grinding quartz in a ball mill. Here it was found that the major part of

the change in milling efficiency could be attributed to changes in the viscosity of the

milling medium caused by the additive, and that when this was subtracted out, the residual

change in efficiency was at best less than 3% (using sodium sulphonapthenate ) . This is

very much less than the 80% or so claimed by Rehbinder as the increase of drilling rate

obtainable with this additive.

100

SO

GRAN1TF BOHUS

1000 RPM

1(00 N

3 bar

• H

2 0

• H

2 °2

" 151 30* A

*RE SHARPEN I NO

v.*

4

to ft 1

Figure 20. Granite: rate of penetration as a function of distance drilled using water, or

10, 15 or 30% hydrogen peroxide solution.

1.5

134
1 00

en

50

UJUJ

a.

, 1 I

>

GRANITE BOHUS

1000 RPM

ow HoO 1 1 p V/

• HYDRAZINE 5%

400 N

C-r /o m —

3 bar MINERAL WATER

-RESHARPENING

•

•

•• ••

•

'-GO

••

•

• 0

m •

•

•

•

'••"••o.o,

•

a • •• • *

* mm

•

0

t

- - ' 1 1 0 1 2

DISTANCE DRILLED (m)

Figure 21. Granite: rate of penetration as a function of distance drilled using water,

soda water or hydrazine solutions

To investigate the third parameter of possible importance, namely that the additive

can have a direct chemical effect on the wear of the diamonds, we have drilled in the presence of hydrogen peroxide. The results of these tests are shown in figure 20, where the

damaging effects of this powerful oxidising agent are immediately obvious. Unfortunately,

the complementary experiment, to drill with a reducing agent such as hydrazine (fig. 21

)

proved disappointing for the drill performance was not perceptibly improved. The figure

also shows results of one final experiment connected with the work on hydrogen peroxide.

It had been imagined that a possible cause for the rapid degradation of the drill was that

the oxygen bubbles liberated during drilling could interfere with the heat-transfer process.

The author is grateful to Professor Nabarro for a suggestion to drill with soda-water , which

shows this not to be the case (fig. 21 )

3. Conclusions

The conclusions to this study may be summarised as follows:

1) When diamond drilling in marble, no additive has been found which increases the instantaneous rate of penetration of the drill. Because the rate of wear of the drill is

very low, however, it was not possible to determine whether this parameter was affected

or not

2) In granite, again, no additive has been found which increases the instantaneous rate

of penetration of the drill. On the other hand, there are numerous additives which

can either increase or decrease the rate of wear of the diamonds.

3) The wear of the diamond appears to be caused by physico-chemical rather than mechanical

means. Oxidation or graphitisation, both thermally activated, are possible mechanisms,

and we have obtained evidence of the former when drilling with hydrogen peroxide.

135
k) How the additive affects the wear of the diamond in general is not clear. We have

suggested three possible mechanisms, as follows:

A) The additive affects the quantity of heat generated during drilling. There is

direct evidence for this, although we have not determined how it is achieved.

B) The additive affects the efficiency of heat removal from the diamond. It is

suggested that the most probable way for the additive to act in small concentrations is to promote wetting and hence nucleate rather than film boiling heat

transfer at the coolant-diamond interface.

C) The additive has a direct chemical effect on the diamond. This has been shown to

be the case for a strong oxidising agent, which is deleterious, but improvements

have not been obtained with reducing agents

.

Finally, the relative importance of these different mechanisms remains to be established in any given situation, without excluding the possibility of additional unexpected

influences

.

References

[1] Rehbinder, P. A., Schreiner, L.A. and Zhigach, S.F., Hardness reducers in rock

drilling (Academy of Science USSR, Moscow, 19hh) (CSIRO Translation, Melbourne,

19hd).

[2] Rehbinder, P. A. and Schchukin, E.D., Surface phenomena in solids during deformation

and fracture processes, in Progress in Surface Science 3, [2] (Pergamon Press,

Oxford, 1972).

[3] Boozer, G.D., Miller, K.H. & Serdengecti, S., Effects of pore fluids on the deformation

behaviour of rock subject to triaxial compression (Fifth Symposium on Rock Mechanics,

University of Minnesota, May 1962).

[U] Robinson, L.H., Effect of hardness reducers on failure characteristics of rock, Soc

.

Pet. Eng. J., 295-300 (Sept. 1967).

[5] Selim, A. A., Schultz, C.W. & Strebig, K.C., The effect of additives on impregnated

diamond bit performance, Soc. Pet. Eng. J., 9., ^25 (1969).

[6] Joris, A.C.T. and McLaren, G. , Mining and Mineral Engineering, 3., [5], 190 (May 1967).

[7] Engelmann, W.H., Unger, H.F. and Snowden, B.S., Diamond drilling with surfactants in

upper Michigan amygdaloidal basalts using surface-set bits, Trans. S.M.E.-AIME, 285 , 185 (Sept. 1975).

[8] Westwood, A.R.C. and Macmillan, N.H., Environment- sensitive hardness of non-metals

(5th Technical Report to ONR. Contract No. NOOOlU-TO-C-0330 . Nr. 032-52U, April 1972).

[9] Shockey, D.A. and Grooves, G.W., J. Am. Ceram. Soc, 51, 299 (1968).

[10] Shockey, D.A. and Grooves, G.W., J. Am. Ceram. Soc, j>2, 82 (1969).

[11] Westwood, A.R.C, Parr, G.H., Jr. and Latanision, R.M. in Amorphous Materials , R.W. Douglas and B. Ellis, eds . , pp. 533- (Wiley, London, 1972).

[12] Hanneman, R.E. and Westbrook, J.H., Phil. Mag., 18, 73 (1968).

[13] Westwood, A.R.C. and Goldheim, D.L., J. Appl. Phys . , 39, 3^01 (1968).

[14] Westwood, A.R.C, J. Mat. Sci . , 9, [11], 1971 (197*0.

136
[15] West-wood, A.R.C, Macmillan, N.H. and Kolyoncu, R.S., J. Am. Ceram. Soc, 5_6, 258

(1973).

[16] Westwood, A.R.C. and Latanision, R.M. in Science of Ceramic Machining and Surface

Fininshing , NBS monograph, S.P. 3^8, p. ihl (NBS, Washington, 1972).

[17] Westwood, A.R.C. and Latanision, R.M., Surface and environmentally sensitive mechanical behaviour - some current issues (Proc. Uth Int. Conf. on the Strength of Metals

and Alloys, Nancy, 1976, to be published).

[18] Swain, M.V., Latanision, R.M. and Westwood, A.R.C, J. Am. Ceram. Soc, 58, 372 (1975).

[19] Macmillan, N.H., Jackson, R.E. and Westwood, A.R.C, Proc. 15th IRSM Symp. on Rock

Mechanics (Sept. 1973 - to be published).

[20] Cooper, G.A., CERAC Report No. U03:5 (May 1977).

[21] Cooper, G.A. and Berlie, J., J. Mats. Sci . , 11, 1771 (1976).

[22] Cooper, G.A. , The Civil Engineer in South Africa, 19, [10], 238 (Oct. 1977).

[23] Mills, J.J., Final report on N.S.F. Contract APR 73-07787 (Feb. 1978).

[2k] Mills, J.J. and Westwood, A.R.C, J. Mats. Sci., in press.

[25] Cooper, G.A. and Berlie, J., J. Mats. Sci., in press.

[26] Rhodes, F.H. and Bridges, C.H., Trans. Amer. Inst. Chem. Eng., 35, 73 (1939).

[27] Clarke, B. , Trans. Inst. Chem. Engrs . , _^5, T. 251 (1967).

[28] Schack, CM., Dean, K.C and Molloy, S.M., U.S.Bureau of Mines, Report of Investigation No. 533^ (1957)

[29] Moreland, C, Can. J. Chem. Eng., kl, 2h (1963).

[30] Haydon, B.C., Report on tests for the Rehbinder effect (imperial College Internal

Report, July-November 1966). Kindly communicated to the author by Dr. J. A.

Kitchener, Royal School of Mines.

Discussion

FIRESTONE

Does the additive reduce the blunting of the diamonds or does it

soften or weaken the bond so the diamonds fall out more easily and

fresh diamonds are more rapidly exposed?

COOPER

As far as I know, the additive has no effect on the strength of the

diamond - matrix bond.

WILCOX

When the diamonds wear , do they get rounded or just fracture and

chip off?

137
COOPER

They are normally rounded, often with striae and ripple patterns reminiscent of ablative erosion. See, however, the text of the paper. There are

two micrographs included.

CHOUDRY

The environment acts on all the three components; namely, the workpiece (rock in this case), tool (diamond) and the matrix (bronze). Effects

on the rock (classical Rehbinder effect) could have been independently

checked by scratch or single cut experiments on rocks soaked and not soaked

in the liquid. Was the 'hardness' of the rock checked as a function of the

liquid by methods other than the drilling tests?

COOPER

We made no measurements of the hardness of our rocks, by scratch, indendation, or other tests. Such experiments have been carried out by many other

workers, and have shown the hardness to be affected at the low strain rates

used in such work. Whether the environment affects the hardness at high

strain rates, we do not know; what we observe is that the instantaneous rate

of penetration is not affected by the presence or absence of the environment

(which is difficult to square with supposed differences in the hardness of

the rock), and that the rate of wear of the diamonds is changed. The main

contention of our paper is not that the Rehbinder effect does not exist, but

that the most important effect of the enviornment is on the rate of wear of

the diamond.

KIRCHNER

We have shown* that greater damage penetration occurs in environments

with zeta potential close to zero compared with similar AI2O3 specimens

scratched in environments with higher zeta potentials. However, low zeta

potential implies clean surfaces so that cracks can heal. Therefore, the

specimens scratched in low zeta potential environments were stronger instead

of weaker than the others showing that these experiments do not always have

the expected results. It seems likely that Westwood's results are correct

but that, at the high drilling rates reported in the present paper, the high

temperatures and other factors have altered the conditions so that comparable

results were not observed.

R. M. Gruver and H. P. Kirchner, "Effect of Environment on Penetration of Surface Damage and Remaining Strength of A1 2 0 3 , M J. Amer. Ceram. Soc. 57

(5) 220-223 (1974).

138
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

THE EFFECT OF CHEMICAL ENVIRONMENTS ON THE

FRACTURE OF CERAMIC SURFACES

Robert E. Cuthrell

2 Sandia Laboratories

Albuquerque, New Mexico 8?185

Environmental effects on the deformation and fracture of a variety of

materials during drilling/grinding and four point bending to failure have been

traced to embrittlement by hydrogen. The extent of embrittlement is determined

by the availability of hydrogen in the drilling fluid or in the gaseous environment in an ultrahigh vacuum chamber (10 Torr). The hydrogen availability is

in turn dependent on its binding in liquids (ionic, covalent, ion paired in media

of low dielectric constant) or by the blocking effect of adsorbed carbon monoxide

in the vacuum chamber. These hydrogen effects were shown in drilling rates, maximum strain to fracture, acoustic emission amplitude (energy) distributions accompanying fracture, and scanning electron microscope photomicrographs of the fractured surfaces. The embrittling effect appears to be universal, affecting in

varying degree the behavior of ceramics, glasses, single crystal minerals, polycrystalline materials, semiconductors, refractory coatings, rocks, and metals

containing brittle phases (organic solids may be exceptions).

Key Words: Acoustic emission; chemomechanical effects; environmental effects;

fracture of brittle materials; hydrogen embrittlement; Rebinder effects, Westwood

effects.

1 . Introduction

Rebinder demonstrated that adsorbed surface- active species could affect the microhardness of certain minerals [i]. These and other manifestations of the effects of chemical en- vironments on mechanical behavior are often termed Rebinder effects.

The influence of several chemical environments on the microhardness of a large number

of ionic and ceramic crystals was investigated by Westbrook and Jorgensen [2], following

earlier work by Mitsche and Onitsch [3] and Walker and Demer [4]. The anelastic behavior of

glass fibers [5»6], the Vickers microhardness and indentation creep of fused silica and alkali

glass [7], the pendulum hardness of soda-lime glass [8,9] and quartz flO] , and the drilling

rates in calcium fluoride [ll] , magnesium oxide [ll] , Pyrex glass [l2j, soda-lime glass [9],

aluminum oxide [l3]t gray granite [13] 1 and. Westerly granite [13] are all known to be strongly

dependent on the chemical environment. Possibly the most significant contribution to understanding these Rebinder effects was made by Westwood, et. al. [l3]» on interpreting the corre- lation between the maximum in the hardness of materials and the zero in the surface charge

(zeta potential) which is produced by ions adsorbed from solution environments.

2. Environmental Effects on Drilling/Grinding

Indicated by Acoustic Emission

In brittle polycrystalline materials, crack propagation occurs at an average velocity,

^This work was supported by the Division of Basic Energy Sciences, U.S. Department of Energy

(DOE), under Contract AT (29-1)- 789.

2

A U.S. DOE facility.

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

139
v, determined by the applied stress intensity factor, K, [14]; usually,

V = AK

n

(1)

where A and n are constants for a specified environment and temperature [15]. Most acoustic

emission count rate studies performed on polycrystalline ceramics [15-I8J have shown an approximate proportionality between the acoustic emission count rate and the average crack ve- locity and the dependence of these on the stress intensity factor. It was shown^recently by

Guthrell thfct the acoustic emission count rates and the peak emission amplitudes were proportional to the low speed drilling/grinding rates'' of brittle materials such as Pyrex glass

[19] and soda- lime glass [20] in various liquid environments. Minimum acoustic emission count

rates (wear rates) were associated with the zero zeta potential (fig. 1) [19] where the maxi- mum hardness would be expected (as shown by Westwood [13] for a wide variety of materials in

various liquid environments). In addition, minimum low speed drilling/grinding rates were found in drilling fluids having maximum binding of monovalent ions as neutral ion aggregates

(fig. 2) [19] (a condition which was mathematically predicted [19] using the theories of

Bjerrum [21,22] and Fuoss and Kraus [21,23]). It was proposed (l) that the higher drilling/

grinding rates were primarily results of the embrittling effects of hydrogen ions from liquid or gaseous environments and (2) that the minima in drilling/grinding rates were results of

reduced availability of hydrogen through its aggregation in ion pairs [24].

xlO

Pyrex, Diomond Bit, 1500g, lOrpm

160-

O 140-

120-

100-

60-

Hi

/MCI3 f\t

/

1 u

\ AI(N03

) 3

y

1 1 1 1 1 1 1 1 1 1

160

80 >

E

0 £

--40

xlO 3

180-

Pyrex, Diamond Bit, 1500g, lOrpm

160-

140-

= 120-

100-

80-

< 60-

1 . Benzene

2. Water

3. Acetone

4. n-Oclane

5. Methanol

6. Glycerol

7. Ethanol

8. n-Butanol

9. Ethylene Glycol

10. Formamide

1 1 . Formic Acid

12. Acetic Acid

13. Toluene

14. Oleic Acid

15. 50»Methonol-Ethanol

Water r -i Log AIXj

L Jmole/liter

Figure 1 . The acoustic emission rate during

drilling (upper curve) and the zeta potential

(lower curve) of glass as functions of the

concentration of trivalent aluminum ions in

aqueous solution [19].

-

r

-

20 40

I 60 —r80 -r—

100

Dielectric Constant

Figure 2.

drilling

constant

The acoustic emission rate during

glass as a function of the dielectric

of the drilling fluid [19].

Acoustic emission was detected over the frequency range 0.125-2 MHz.

'Drilling was performed using a spherical- ended diamond studded bit rotated at 10 rpm under

a 1500 g load in various drilling fluid environments. Drilling at this very low speed may

be more analogous to grinding than to conventional high speed drilling.

ult:

BOX.

tic

0CC1

alsi

anpl

Appa

prod

liqu

low

littt

anpl

[20]

tmies

tonsi >s

fract

fttat

soiirc

140
3. Embrittling Effects of Hydrogen

In order to test this theory, several drilling experiments were performed

under very carefully controlled conditions in an ultrahigh vacuum chamber. Figure 3 [2^] shows

the acoustic emission rate as a function of drilling time in ultrahigh vacuum (5 x 10 Torr),

in carbon monoxide (10 Torr), in a mixture of carbon monoxide (10~' Torr) and hydrogen (680

Torr), and in hydrogen alone (680 Torr). For Pyrex glass the presence of at least 10~' Torr

of carbon monoxide appears to inhibit the diffusion of hydrogen into the solid. There is a remarkable change in behavior when the carbon monoxide is differentially pumped out such that

the glass is exposed to hydrogen alone. The effect is reversed when the hydrogen is pumped

out and the return to the original condition is accelerated by heating to 100°C in vacuo

Minutes Days

Time

Figure J. Acoustic emission count rate as a function of drilling time for Pyrex glass in an ultrahigh vacuum chamber. Note that there is no effect of gaseous environment in carbon mo- noxide nor in hydrogen when the carbon monoxide is present. The time required for the acoustic emission rate to reach a maximum in hydrogen indicates that Rebinder-Westwood effects occur in the near-surface region rather than solely at the surface. The time dependence is

also consistent with a hydrogen permeation mechanism for the chemomechanical effect, [24] .

Figure 4 [24] shows the shift to higher amplitudes (greater energy release on fracture)

of the acoustic emission peak on exposure to hydrogen and reversal of this trend on pumping

out the hydrogen. Drilling in air or under water environments results in acoustic emission

amplitude distributions very similar to the lower right, figure k, as shown in figure 5 [24]

Apparently sufficient hydrogen is available from water vapor in air or the liquid water to

produce the embrittling effect. As shown in figure 5» the results obtained for drilling under

liquids follow the same trend observed for hydrogen in the gaseous state , i.e., a shift to

low amplitude for drilling in a hydrogen free liquid (carbon tetrachloride) and a higher amplitude distribution for drilling in a hydrogen bearing liquid (water). Peaks of intermediate

amplitude were produced for drilling under hydrocarbons in which hydrogen is covalently bound

and is thus less available than in a partially ionized aqueous environment [24],

Figures 6 and 7 show similar results for polycrystalline alumina [25] and for silicon

[20] (the intermediate amplitude for drilling alumina under carbon tetrachloride may be a re- sult of moisture trapped in cracks, crevices, and pores in the sintered solid). The changes

from peaks at higher amplitudes in hydrogen bearing environments to progressively lower amplitudes in drilling fluids containing less hydrogen were found to be time dependent and thus

consistent with a hydrogen diffusion mechanism for embrittlement. Figure 8 shows that drilling silicon under carbon tetrachloride, a hydrogen free liquid, produces primarily ductile

fracture in the near surface layer (ductile plowing in smooth grooves which follow the drill

rotation) while primarily brittle fracture (with some ductile plowing) occurs under water, a source of hydrogen, for the same sample surface.

141
Drilling Pyrex Glass Dril 1 ing Pyrex Glass

Amplitude, dB

Figure Acoustic emission amplitude distributions for the drilling of Pyrex glass. The

upper left curve is typical of each of the

three conditions listed. The drilling times

correspond to the abscissa in figure 3

30 40 50

Amplitude (A), dB

60

Figure 5. Acoustic emission amplitude distributions for the drilling of Pyrex glass under

carbon tetrachloride (a hydrogen free liquid)

and under water (a source of hydrogen) \_2k].

U

ci

ill

kr:

fri

sti

ta

dai

as;

Drilling Polycrystalline Alumina

Drilling Silicon

Amplitude (A), dB

Figure 6. Acoustic emission amplitude distributions for the drilling of polycrystalline

alumina in vacuo , under carbon tetrachloride,

and under water. The intermediate amplitude

for drilling alumina under carbon tetrachloride may be a result of moisture trapped in

cracks, crevices, and pores in the sintered

solid [25].

T—i

1 r

10 20 30 40 50

Amplitude (A), dB

60

Figure 7. Acoustic emission amplitude distributions for the drilling of the (ill) plane

of single crystal silicon under carbon tetrachloride and under water [20]

.

142

fra

for

con

int

tai M

He

stn
on

c-hydk. Ductile, Brittle, More Brittle

The assignment of "ductile" or "brittle " designations for acoustic emission amplitude

distributions was based on the observation that curves decreasing asymptotically to the abscissa were obtained for the drilling of a wide variety of metals which were known to be ductile, whereas peaks at higher amplitudes were obtained without exception for the drilling of

brittle materials. The effect on the acoustic emission amplitude distributions of changing

from a ductile to a brittle condition in the same sample (Ferrovac E iron) without changing

structure or composition was shown (fig. 9) by driving above and below the ductile- to- brittle

transition temperature (about -73°C) in vacuo (10 Torr, to avoid ice, condensation, or oxidation) [20], These data, obtained under very carefully controlled conditions, confirm the

assignment of "ductile/brittle" designations for acoustic emission amplitude distributions

and show that acoustic emission measurements may be used to distinguish between these two

fracture modes. Figures k-7 and 9 show "brittle" peaks which occur at different amplitudes

for different materials under the same conditions and for the same material under different

conditions. In order to facilitate comparison, the concept of "degrees of brittleness" is

introduced. Thus a material for which an acoustic emission peak at higher amplitude is obtained on fracture may be said to be more brittle under the particular experimental conditions.

Initial results of a study of the environmental dependence of slow crack growth in glass indidicate that higher acoustic emission amplitudes during drilling may be correlated with lower

stress intensity factors for crack propagation in the same environment [2?]

Drilling Ferrovac E* (99.90%Fe)

N'AE

23'C

(Ductile)

(Brittle)

0 10 20 30 40 50 60

Amplitude (A), dB

Figure 9. The drilling of Ferrovac E iron at

23°C produces an acoustic emission amplitude

distribution which is typical of ductile materials. Drilling the same sample at -115°C

(well below its ductile-to-brittle transition

temperature of about -73°C) produces a peaked acoustic emission distribution at higher

amplitudes which is typical of brittle materials [20].

143
5. Conclusions

The results presented indicate that the embrittling effect of hydrogen is one of the

origins of the environment-dependence of the deformation and fracture of the near surface re- gion of solids (the Rebinder-Westwood effect). It was indicated that the results for drilling

solids under hydrogen bearing liquids were the same as those for drilling in a well-controlled

gaseous hydrogen environment except that the availability of the hydrogen in the liquid case depends on its bonding (covalent, partially ionized, or bound as ion aggregates). The time

dependence of the chemomechanical effect is consistent with a hydrogen permeation process.

Although silicon exhibits one of the largest observed hydrogen-dependent changes in me- chanical properties, similar effects have now been observed in many quite different materials i (Pyrex glass [19,2*4-], soda- lime glass [20] , silicon J20] , fused silica [25], quartz [25], polycrystalline alumina [25] , calcite [25], fluorite [25J, sierra white granite [25], Texas pink

granite [25], hot pressed boron carbide (B^C) [25] . hot pressed titanium diboride [25], type

02 tool steel [25], 1080 steel [25]. and *440C stainless steel [26]). It is proposed that em- 1 brittlement by hydrogen is a universal effect (organic solids may be exceptions) which would

be expected to affect other materials and other forms of wear in addition to low speed drilling (such as high speed drilling, erosion, grinding, polishing, comminution, machining, and

cutting). It is also proposed (l) that the environmental effects on slow crack growth or fa- | tigue in brittle materials, which have been previously attributed to water, may be attributable to hydrogen from the water, and (2) that embrittlement by hydrogen may be the origin of

the Rebinder-Westwood chemomechanical effects.

The author wishes to thank D. M. Mattox for suggesting acoustic emission techniques for

this investigation, E. Randich for fruitful discussions and metallurgy contributions, and

V. D. Nogle for sample preparation.

References

[l] Rebinder, P. A., Proc. 6th Physics Conf . , Moscow, 1928 (unpublished).

[2] Westbrook, J. H. and Jorgensen, P. J., Trans. Metall. Soc. AIME 233, ^25 (1965).

[3] Mitsche, R. and Onitsch, E. M. , Mikroskopie 2* 257 (1957).

[*»-] Walker, W. W. and Demer, L. J., Trans. Metall. Soc. AIME 230, 6I3 (196*4-).

[5] Aslanova, M. S., G. R. (Dokl.) Acad. Sci. URSS 1215 (195*0.

[6] Aslanova, M. S. and Rebinder, P. A., C. R. (Dokl.) Acad. Sci. URSS 96, 299 (195*4-).

[7] Hanneman, R. E. and Westbrook, J. H. , Philos. Mag. 18, 73 (1968).

[8] Westwood, A. R. C. , Parr, G. H. , Jr., and Latanision, R. M. , in Amorphous Materials , edited by R. W. Douglas and B. Ellis (Wiley, London, 1972), p. 533.

i [9] Westwood, A. R. G. and Huntington, R. D. , in Mechanical Behavior of Materials (Society

for Material Science, Japan, 1972), Vol. IV, p. 383.

[lo] Westwood, A. R. G. and Huntington, R. D. , (unpublished).

[ll] Westwood, A. R. G. and Goldheim, D. L. , J. Am. Geram. Soc. 5J, 1*4-2 (1970).

[12] Westwood, A. R. G. and Latanision, R. M. , Natl. Bur. Stand. (U.S.) Spec. Publ. 3*4-8 , 1*41 (1972).

[13] Westwood, A. R. G. and Macmillan, N. H. , in The Science of Hardness Testing and Its

Research Applications , edited by J. H. Westbrook and H. Conrad (American Society for

Metals, Metals Park, Ohio, 1973), P. 377.

144
Wiederhorn, S. M. , in Fracture Mechanics of Ceramics , edited by R, G. Bradt, F. F. Lange,

and D. P. H. Hasselman (Plenum, New York, 1974), p. 613.

[15] Evans, A. G. and Linzer, M. , Ann. Rev. Mater. Sci. £i 179 (1977).

[l6] Evans, A. G. and Linzer, M. , J. Am. Geram. Soc. J56, 575 (1973).

[17] Evans, A. G., Linzer, M. , and Russell, L. R., Mater. Sci. Eng. 1J5, 253 0-974).

[18] McKinney, K. R. , Ghaskelis, H. H. , and Freiman, S. W., Bull. Am, Ceram. Soc. 3j, 618

(1974).

[19] Cuthrell, R. E. , J. Appl. Phys. 49., 432 (1978).

[20] Cuthrell, R. E. , The Influence of Hydrogen on the Deformation and Fracture of the Near

Surface Region of Solids — Proposed Origin of the Rebinder-Westwood Effect, J. Mater.

Sci., in press.

[2l] Harned, H. S. and Owen, B. B. , The Physical Chemistry of Electrolytic Solutions (Reinhold, New York, 1958), p. 70.

[22] Bjerrum, N., K. Dan. Vidensk. Selsk. Mat.-Fys. Medd. £, (No. 9), (1926).

[23] Fuoss, R. M. and Kraus, C. A., J. Am. Chem. Soc. j>5_, 2387 (1933); j>7, 1 (1935).

[24] Cuthrell, R. E., Near Surface Embrittlement of Solids by Hydrogen — The Origin of the

Rebinder-Westwood Ghemomechanical Effects, J. Appl. Phys., in press.

[25] Cuthrell, R. E. and Randich, E. , (unpublished).

[26] Cuthrell, R. E, , Surface Effects in Crystal Plasticity , edited by R. M. Latanision and

J. T. Fourie (Noordhoff, Leyden, 1977), p. 773.

[27] Marion, R. H. and Cuthrell, R. E. , (unpublished).

145

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SOME RESULTS ON THE WEAR OF A BRONZE-BONDED GRINDING WHEEL

A. Broese van Groenou

Philips Research Laboratories

5600 MD Eindhoven, The Netherlands

Abstract

Small parts of a diamond grinding tool have been investigated in a grinding

apparatus where the tool is fixed whereas the workpiece material, here a MnZn

ferrite, moves at 20 m/s . The wear of the bronze binder and of the diamonds can be measured on plastic replicas. Apart from an estimate of the average G value

it is also possible to study locally the erosion of the binder and the wear of

the diamonds. Two areas of erosion are distinguished: the region between the

diamonds and the zone in front of a diamond. For a dozen diamonds it is found

that the average erosion rate is proportional to the removal rate of the workpiece, whereas the local rate of increase of zone depth is about twice as fast.

For this type of wheel the conditions for self-sharpening can be formulated.

For a nickel-plated wheel the wear was too small to be measureable

Key words: Bronze binder; diamonds; erosion; grinding; wear.

1. Introduction

As part of a program on the grinding of ceramics and glass some aspects of the wear of a grinding wheel have been studied. Usually the wear is characterised by the G value, G being

defined as the ratio of the volume of material removed by grinding to the loss of grinding

wheel volume. For grinding steel with alumina wheels [ 1] ^ and carbide tools with diamond

wheels [2] G is between 100 and 1000. Much higher values (10^ - 105) are found in the grinding of ceramics and glass with diamond wheels. The correspondingly long life of the wheels

allows the grinding of a cubic metre of ceramic. If one is not satisfied with this performance, one may ask: why not a longer life? Which wheel should one choose ? Can dressing

be avoided? In view of the large number of wheel variables the answers are not readily

available. Rate of workpiece removal Z and rate of wheel wear are related. For high Z

values the grinding force increases and so does the force F3 exerted on a dimaond. Introducing the grinding hardness Hg one has for a given workpiece material

F, = H Z/v N, ( 1 ) d g s d

where vg is the wheel speed and N^ the number of diamonds in action. The depth of cut of a single diamond is of the order of a micron or less. Moreover, the height of the diamond

tips above the binder varies over half the size of the particles, e.?. 50 to 100 ,um. Not even on this sale is the radius of a grinding wheel constant, which implies that even

a slight unroundness has a pronounced influence on N^. In order to study wear it is therefore necessary to locate the site of action with some precision. The apparatus to be

described has been designed especially with this requirement in mind.

Wheel wear involves the diamonds and the binder. Diamond wear may not contribute much

to the loss of volume because of the small particle volume, but it does influence binder

wear if it is true that "the diamond protects the matrix"[ 3] . A detailed study of the

wear of diamonds is not possible on a grinding wheel, because the unequal protrusion of the

particles above the binder, even when it is flat, means that the force on each particle

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

147
is very uncertain. Single diamond studies are more appropriate [ k] . In the present paper

the wear of the binder has "been measured quantitatively, enabling the G values to be determined. It was also possible to determine the number of diamonds in action, especially the

changes in this number due to diamond loss and "birth". This is of interest, because it

gives information on the "sharpness" of a wheel. It is important that a wheel remains sharp,

since the continuity of the grinding operation is essential in manufacturing and corrections

by dressing are very costly. Sharpness is therefore more important than the G value, for

only after a continuous operation has been realised, can G be used to calculate the wheel

contribution to the cost of grinding.

This paper describes a wear apparatus where the grinding "wheel" is fixed and the workpiece moves at a relative velocity of 22 m/s. The grinding tool is moved into the samples

at a speed vw . Normally, every abrasive particle acts once during a wheel revolution. Such

intermittent action was imitated here by interrupting the contact with the workpiece: twelve

samples were used. During the interruption the diamond tool is cleaned and cooled by a mist beam of water and air. The force on the tool and the volume removed from the workpiece are measured. The detailed behaviour of the diamonds and the binder is determined by

observing replicas of the tool. The apparatus is described in the next section. Data

on the combination of a MnZn ferrite ceramic and a bronze-bonded diamond "Wool are given

in section 3. Other binders were investigated, especially a nickel-plated tool, but the

small amount of wear could not be measured. Wear details are given in section h, and in

section 5 we return to the question of the wheel sharpness in the discussion.

2 . Experimental

a. Workpiece wheel

The apparatus consists of a wheel containing twelve workpiece samples and a mechanism

feeding the diamond tool. The wheel is connected via a flexible coupling to a 1.5 kW

motor running at 1^00 r.p.m. The wheel is supported by air bearings (2k0 mm i.d., 300 mm

o.d., TO mm high) fed by an air supply at 7 atm. Two-stage air filtering is used. In

the axial and radial directions the running inaccuracy is less than 0.5 /Urn. The workpiece samples are fastened on the face of the wheel (figs. 1 and 2) and not on the circumference because a flat wear plane was preferred to a curved one. The sample speed is 22 m/s,

which makes the contact time 0.^5 ms for a sample width of 10 mm. The sample blocks,

9x10x11 mm3, are glued in steel blocks in a h mm deep racess. The apparatus was carefully

balanced. For this purpose the wheel was provided with balancing blocks.

Figures 1 and 2: Photograph and schematic

cross-section of the grinding and wear

apparatus

.

148
b. Tool motion

The diamond tool is a block, 20x10 mm, cut from a grinding wheel. It is held in a brass holder, together with the dynamometer2 for the measurement of up to three force com- ponents. The holder is subject to a horizontal and a vertical (fig. 2) motion. The first

one is parallel to the axis of the wheel: it consists of a coarse motion by hand over 20mm

and a fine motion by motor over 1 mm. In order to increase the stiffness of the feeding

construction, a wedge is used (held in cylinder bearings) which is moved by a spindle and

motor. The feed rate vw

was given the values vw= 19, ^3 and 56 /um/s , which correspond

to a depth of cut of a = vwT

g =0.76, 1.71 and 2.22 /um, where Ts is the motor period.

The rate of material removal Z, is given by Z = abvs with b = sample width; introducing

the equivalent chip thickness heq by Z p he(, b vs

, one has a = h „ for this apparatus.

The Z values follow from v : Z-| = 166, Zp = 376 and Zo = I+89 mmJ/s.

The second motion of £he tool is in the vertical direction, in order to avoid the

diamonds repeating the same grooves on the ferrite samples. By means a second d.c. motor

and an eccentric a vertical amplitude of 0.8 mm is obtained with a period of O.kh sec. This motor moves the wedge, its bearings, the tool and the first motor against the conteraction of two horizontal springs.

c . Samples

The ceramic samples were cut from blocks of MnZn ferrite (commercial grade, porosity

Q%) , taking care that all twelve samples were the same height when fastened on the wheel.

The diamond tool was cut from an L-bronze wheel3, containing nominally 80 /um diamonds

in a concentration ho (10% volume concentration). The Vickers hardness of the binder was about 300 kg/mm2 ( 3. 109 N/m2

).

The experiments were performed as follows: for each Z value several runs were made;

in each run about half of the heights of the ferrite blocks, i.e. about 6 cm3 in total, was ground away. After a run the tool was inspected, the blocks were replaced and a new

run was started. Five runs were made in the first stage (Z^), three in the second

(Zp) and third (Z^). The data will be characterized by the oumulative volume of ferrite

removed at the end of the run and by the removal rate (Z) used in the run.

Before every run the tool was observed under a microscope and scanned by a rolling

wheel as used in a roughness measurement. This procedure was found to be less satisfactory than making a permanent replica and measuring the wear there . Such replicas were made

by pouring a mixture of powder and liquid^ on the diamond sample, which had first been

surrounded by a plasticine dike. With some care air bubbles can be avoided. On the various

replicas the individual diamonds could be followed as well as the details of the binder.

3. Results

a. Grinding force

The output of the dynamometer showed twelve pulses per revolution of the wheel, as expected for the same number of samples. Superposed is a slower modulation (period O.kk s)

due to the vertical motion. The force values, averaged over the modulation, showed no significant change with time due to wheel wear (with a much softer wheel we did observe such

changes). The average normal and tangential force components increase with increased feed vw , from which the grinding hardness (20. 10^ N/m2 ) and the specific energy (^.109 J/m3)

can be deduced. These values correspond to a = he ^

= 1 um. At higher he^ the specific

energy decreases strongly, as found before at wheel speeds of 60 to 100 m/s in a horizontal

axis surface grinder [ 5] •

2. Kistler, 925 1 A.

3. Diamant Boart, Brussels

k. Technovite 5071.

149
b . Average wear

On the successive replicas the amount of wear was measured with a microscope, using a conspicuous diamond as a reference. This particular diamond, a rectangle, was present in

all replicas. Because of its uneven surface it was possible to measure the wear of the top - about 5 ym over the whole range of experiments - but this was negligible compared with the

amount of wear of the binder as will be shown below. By scanning the replicas over a fixed

line through this diamond, the profile was measured at 21 equally spaced points (fig. 3).

The steeply rising parts at the ends of the profile are caused by the parts of the tool

which did not show wear because the sample is less wide than the tool. The fluctuating data

in figure 3 have been averaged over the active area. As can be seen in figure k, the averaged results are linear in the amount of ceramic that was removed. From these data the G

ratio can be calculated, taking into account that wear occurred on about half of the tool

surface. The G values range from 29000 at low Z (h = O.76 ym) to 11000 at high Z (heq

= 2.2 ym).

Fig. 3. Profile depths measured over the

length of the replicas; the arrow points to

the fixed reference diamond (k) . Replica

depth increases with tool wear. Data for

various amounts of ferrite ground away (V^)

.

Fig. 4. Average profile height deduced from

figure 3 versus ground volume. The three

regions of different Z values are indicated.

The zero point of the vertical axis has been

shifted with respect to that of figure 3.

On the replicas the number of visible diamonds per unit area can be determined. At low Z the number is about 7 diamonds per mm2

, independent of the ferrite volume. Loss of diamonds

is apparently balanced by the emergence of new ones. At the highest Z value a decrease is

observed to 5 per mm2 (fig. 5). The details of this process will now be considered.

100%

80%

0 10 20 30 40 50 60

ferrite volume (cm 3

)

Figure 5. Number of diamonds per unit area, visible on replicas of the active part of the

tool (full line), and the fraction of diamonds with a groove (broken line, righthand scale), both as a function of the volume removed by grinding.

h. Details of wear

Comparing the same area on consecutive replicas one can observe the life and times of

the various diamonds. Figure 6 shows a number of particles where all wear phenomena can be

150

an;
seen. The diamond on the right (to be labelled j) is already active after 6 cur5 has been

ground; it shows wear flats, a widening groove in front and on both sides, and a hill behind it, which disappears as soon as the particle itself is lost (for this particle after

grinding h2 cm3). A second diamond (to be labelled e) appears on the left side of j and

quite near to a square diamond. No wear is visible on e , and it is uncertain whether

this particle contributed to the removal of ferrite. In the centre two diamonds become

visible after 12 cm3, the upper one (to be labelled a) shows a nice groove after k2 cm^,

whereas the binder in front of the lower one seems to be protected, although this particle

is subject to some wear.

Figure 6. Diamonds a, e and j after grinding various amounts of ferrite. Photo (a): V

p

= 6cm 3 diamond j on the right, a and e just emerging; (b) : V

p

= 12cm 3

, a's edge above the middle, ju

right of a lost diamond, e is next to the white square on the right, both on the left of

j; (c): 24cm 3

, all diamonds quite well visible, a groove is forming around a and the white

square; (d) : 42cm 3

, a has a deep groove and e's two neighbours have vanished, j's groove is

nearly gone.

The fraction of diamonds with a groove on three of its sides increases from 30% at

low Z to 10% at Z^, where it flattens off (fig. 5). In contrast to the G value, which is

constant for each Z, the number of diamonds per unit area and the fraction with a groove

still varies in the constant Z ranges.

We noted that a groove appears when the diamond emerges above the binder over a distance of

10 to 15 um. The groove in front is narrow, where the diamond has an edge, and wide where

it has a flat plane. The groove on the side becomes longer with increasing grinding time.

151
The hill behind the particle seems to be a passive zone, since it appears only when the

side groove makes it visible. A schematic view of the groove and hill is shown in figure 7.

Ed

D DB DG /

m

.M l! D-DG

Figure 7. Schematic cross-section of diamond

in binder, indicating the distances DG, DB

and BG, the forces F^ (due to grinding) and

Fp(ploughing the binder near the groove).

D is the diamond size.

-£70

360

8 50

I AO m^ 30

| 20

10

DG

diamond j BG

"db

^>— ,

80-

E 70-

3 60-

(Uo 50-

co 40-

Vio

30-

A 20-

10-

10 20 30 i.0 50

ferrite volume (cm 3

]

60

diamond a

^-"'A

DG

*>—""db

' -o BG

"t

i 10 20 30 UQ 50 60

ferrite volume ( cm 3

)

Figure 8. Measured values of the distance between top of diamond and binder (DB), diamond

and groove-bottom (DG) and binder to groove (BG) for diamonds j and a. The three ranges

of Z are shown. Along the abscissa is the volume ground away. After loss of diamond j the

groove erodes away rather quickly (BG).

Table I. Data on selected diamonds of various shapes and with or without neighbours.

Name Shape Start Loss DG(um) Erosion slopes (um/cm ) Neighbour Wear Remarks

run( stage

)

max. DB BG DG width

a 3(1) 73 0.96 0.52 1.U8 1.76 yes no

b 1(1) 8(11) 81 0.86 1.36 2.22 1 .76 no yes fractured

flat

c Ml) 9(IH) 61 0.70 0.U6 1 . 16 2.31 no yes chipped

d 5(D 75 0.72 1.01 1.73 0-1.7 yes(2) yes flat

e 3(1) 6h O.65 0.1+5 1 . 10 2. 12 yes( j ,a.o) no

f 1(1) 9(IH) 66 0.53 1 .01 1.51+ 1.79 no yes flat

g 7(IH) 59 0.98 2. 17 3.15 2.08 no

h 5(1) 10(111) 65 1.37 0.72 2.08 2.33 yes(i) yes chip

i 2(1) 10(111) U5 0.38 0.20 0.58 0.60 yes(h) yes chip

j KD 8(11) 65 0.3U 0.7^ 1.08 yes(e) yes flat

k 1(1) lh 0.57 0.69 1.25 1.53 no yes flat

average erosion slope 0.73 0.85 1.58

standard deviation 0.30 0.5^

Table I. Data on selected diamonds of various shapes and with or without neighbours. The run of first appearance is given under Start, the last run before loss is given under Loss.

The runs are labelled 1-5 (i), 6-8 (II) and 9-11 (ill), where the roman numerals refer to

the Z values. The erosion slopes were determined from plots such as in figure 8. The definition of DG etc. is given in figure 7.

152

I
Quantitative data were obtained on eleven diamonds during the grinding of 5*+ cm of

ferrite. The rate of erosion of the binder far from the diamonds was measured from the

distance between the top of the selected diamond and the level of the binder: diamondbinder data (DB) are shown in figure 8 for the diamonds e and j of the previous figure.

Also shown are the distance between the average binder level and the deepest part of the

grooves (BG) and the sum, that is the distance diamond-groove (DG). The geometric meaning

of DB, BG and DG is shown in figure 7 • All distances increase linearly with the volume

of ferrite removed Vf. The slopes of the lines are given in table I, together with information about the shape of the particles, the presence of wear, the stage where the particle

became visible and where it had disappeared, the corresponding maximum DG value and the

presence of other particles nearby. No systematic variation is seen with shape, amount

of wear or the presence of neighbours. After the particle has disappeared, e.g. diamond

j in figure 6, the groove gradually disappears (fig. 8), where the hill vanishes within

6 cm3 of ferrite volume. The DG value at loss is 65 to 80 urn, the BG value is about half

this value.

5. Discussion

a. Binder erosion

In the previous sections three modes of wear where reported: diamond wear, loss of

diamonds and binder erosion. The first two are unimportant in terms of volume and therefore of no direct consequences for the G ratio. The grinding action, of course, takes place

where the diamonds are, which reduces the relevance of G as a grinding characteristic

to a quantity that determines cost, once the wheel grinds, but does it? The capabilities

of a wheel must be described by the details of the interplay of diamonds and binder together. The life of the particles in the case of the bronze-bonded wheel which we used

consists of the following stages:

1. The diamonds emerge because the binder wears away; the rate of change of the number of

diamonds visible on a unit area N is related to the change in the binder level and thus

to the change in the distance DB by

g = n[^b]/L = NAZ/L (2)

where L is the interparticle distance, A is the slope of the DB- ferrite volume line in

figure 8 and Z = dV^/dt. From a count of the number of diamonds (Fig. 5) L is calculated

to be O.k mm, which is a factor two more than expected on the basis of diamond size D

(80-100 urn) and concentration C ( 10$ by volume).

2. A groove is formed when the top of the particle is about 15 to 20 urn above the binder.

The groove first forms in front of the particle, then also on the sides, leaving a passive

area behind it; this grows into a hill, because of the groove getting deeper. Where the

particle has an edge in front, the width of the groove is smaller than where it shows a plane face. Groove depth and width increase with time, just as the length and height of the

hill.

3. At a given moment the diamond is lost. It may be that this happens because the distance

DG between the top of the particle and the bottom of the groove has reached a critical

value, between 50 and 80 urn. The hill disappears very quickly, the grooves take longer to

disappear. In this stage the change in binder-groove distance, BG, is faster than the

binder erosion itself, given by DB (fig. 8).

From these observations one can state precisely what is meant by "the diamond protects

the matrix "

[ 3] : If there is no diamond, the matrix (i.e the binder) wears away, whereas

if there is a diamond, part of the matrix wears away much quicker (the groove), part does

not wear at all (the hill), until the diamond is lost. The latter process does not contribute much to the G value, at least for the Z values used here, because diamond loss is a relatively rare event. A "background" value of binder loss is given by the change in the

height DB, taken far from a diamond; as in eq.(2) the rate of change equals A* Z« from table

I, g_ (DB) = 0.73 Z, which gives a G value of 55000 independent of Z. The G values reported here are smaller and do depend on Z. The reas.on is the formation of grooves near

the diamond: their width, depth and length all increase with Z, their volume thus increases stronger than proportional to Z.

153
The presence and the shape of the grooves near the place where the grinding debris

is formed suggest that the hinder wear is an erosion process, caused by the hard ceramic

particles, probably in combination with the cooling water. The wear on the flat binder is

less than near the diamonds. This resembles the results of Thomas and Brunton [6] in their

study on erosion of metals by water jets at speeds up to 125 m/s. They found that on a flat surface hardly any wear occurs, whereas a disturbance, e.g. a groove , has a considerable effect. This may also be relevant for the faster binder wear near the groove when

the diamond has disappeared. The rate of change of the binder height can be discussed in

terms of the well-known relation between rate and pressure,

t = k P V

s /H

b <*J

where is the binder hardness, k a constant and P the unknown pressure exerted by the

particle responsible for the material removal. From the average values of the rate of

change of DB and DG in table I and the Z values, kP is calculated from eg. (3) to be of the

order of 100 N/m^. For a free particle erosion [ 7>8] P is expected to be of the order of

\ p v^, where p is the particle density (here about 5.103 kg/m3) and v the particle velocity. If v equals vs

, the P value would be 10^ N/m^, which makes k about 10~^.

One reason for this low value may be that the horizontal motion of the water in the

gap between workpiece and binder disturbs the vertical motion of the debris that is generated at the top of the diamond and moves along the leading plane downwards to the foot

of the diamond. The water moves near the top at vs = 22 m/s and is presumably turbulent

(this raises the question of whether more viscous liquids would show less binder wear).

For smaller diamonds the high water speed would be more effective. Another reason for low

wear rate near the diamond might be that perpendicular erosion [ 7 5 8] on metals is much

less effective than when the particles impinge under an angle of 10-20°. The latter case may apply to the wear of the flat part of the binder, as measured by DB in table I.

The experimental data are limited to one combination of binder (%), wheel speed

(vg

) and diamond size (D). In the erosion model one would expect higher wear rates for

softer binders , higher wheel speeds and smaller diamonds . The present results are too

limited to check this.

b. Diamond loss

The lifetime of a diamond, from its emergence to its loss, plays an important part in the

choice of a binder for a given work-piece. The short duration of the present experiment

did not allow the quantitative determination of the life for a single Z value. In the

following model a crude estimate of the lifetime t

c can be made from the grinding force

on the one hand and on the other from the "holding force" of the binder at a given groove,

DG. One finds (see appendix)

= BZ BZ [ VVvDvJ '

'

s d

where B is the slope of the DG - ferrite volume line (Fig. 8), H^. is the grinding hardness

of the workpiece and N, the number of diamond in contact.

When the lifetime of the diamonds is too long and the particles become blunt, the expressions (1) and {h) can be used to determine which factor (H^, D, vg

, Z, C) should be changed

Although a number of rules found in the literature agree with this equation [3, 9» 10] > it

is obvious that the limited extent of the present experiments is not sufficient to prove

their validity.

6. Conclusions

An apparatus has been described which allows the measurement of the wear of small

parts of a grinding tool, as well as the grinding force. With the aid of replicas of the

tool made at regular intervals , the details of the diamond and binder wear could be followed. Apart from some diamond blunting ,the formation of grooves near the diamond was observed. Loss of diamond occurred when a certain groove depth (here 50-80 urn for 80-100 urn diamonds) had formed. The life of the diamonds is determined by the erosion of metal

154
binder, which is presumably caused by the action of the debris particles. Although detailed

results were recorded, which allowed some conjectures about diamond loss and the need for

dressing, the experiment was limited to one binder, one workpiece material, one wheel speed

and one diamond size and concentration. Obviously more data are necessary for an accurate

prediction of the behaviour of the combination of workpiece, diamond and binder.

The author is indebted to a number of colleagues for discussions about the principle

of the investigated wear tester, to P. Merkelbach for the design of the apparatus and to

M.A. Immerix for his dedication in testing the apparatus and performing the experiments.

References

[ 1] Malkin S. and Cook N.H. , The wear of grinding wheels, J. Eng. Ind. (Trans. ASME) 93,

Bk, 1120-1133, (Nov. 1971).

[2] Freeman J. P., Some aspects of the evaluation of diamond wheels in dry grinding (Proc.

Int. Ind. Diamond Conf . , Oxford 1966), Book, Science and Technology of Industrial

Diamonds II , J. Burls, ed. , pp. 263-271 (ind. Diam. Inf. Bur., London, 1967).

[3] Thomas D.M., Some aspects of wheel formulation in ceramic grinding, Ind. Diam. Rev.,

pp. 50U-507 (Dec. 1971 )

[ k] Gielisse P.J., Kim T.J., Choudry A., Short J.F. and Turker E.J., Force and wear analysis

in ceramic processing, Report, pp. 35 ff. (Univ. Rhode Island), Kingston R.I. (1972).

[5] A. Broese van Groenou, Grinding experiments on ceramics, (Proc. 9th Int. Conf. on Science of Ceramics, Noordwi jkerhout , November 13-16, 197T ) » Book, Science of Ceramics

Vol . 9

,

K.J. de Vries ed. , pp. U86-U93 (Nederlandse Keramische Vereniging, 1977).

£6] Thomas G.P. and Brunton, J.H. Drop impingement erosion of metals, Proc. Roy. Soc

London, A3jji, 5^9-565 ( 1970).

[7] Oh, H.L., Oh K.P.L., Vaidyanathan , S. and Finnie , I., On the shaping of brittle solids

by erosion and ultrasonic cutting, (Proc. Symp. Science of Ceramic Machining and Surface

Finishing, Gaitersburg, Md, 2-h Nov. 1970) pp. 119-132 (NBS Spec. Publ. 3^8).

[8] Engel, P. A. , Impact wear of materials , pp. 106-110 Elsevier Scient. Publ. Cy, Amsterdam,

(1976).

[9] Reed, J.R. , Applying metal bond diamond blades - art or science? Ind. Diamond Rev.,

pp. 291-296 (Aug. 1973).

[ 10]Decneut, A. And Aerents R., Etude theorique et pratique du comportement des meules

diamantees et a base de nitrure de bore cubique, (Preprints Intergrind 76, Nov. 1976,

Vastervik , Sweden )

Appendix

The lifetime of a diamond

The following model takes into account the presence of a groove in front of the

diamond and the grinding force which tries to rotate the particle against the resistance

offered by the binder. The binder is assumed to yield when the stress near the groove

exceeds the yield stress of the metal. The shaded area" in figure 7 is removed by a "ploughing" force

,

F

p

= b (D-DG) (A1)

where b - D is the width of the diamond and the stress necessary for the ploughing.

The grinding force F^ is given in eq.(l). As soon as F<j exceeds Fp the particle rotates

and is lost. This occurs when at time tc the distance DG has reached its maximum value

DGmax =BZtc^ , , F = H D(D-BZ t ) = F, = H Z/v N^ (A2)

p t> c a w s d

This leads to eq.(U). The ratio H /H_ may be about 10; this is multiplied by a term of the

order of d

2 /D2 . Since in heavy grinding the depth of cut for one diamond, d, may be an appreciable fraction of its size D, the second term in t

c is by no means negligible.

155
Discussion

CHOUDRY

Could not the formation of the 'hill' and the 'groove'; behind the diamond grain; be interpreted as the process in which the diamond grain acts on

both the workpiece and the bronze matrix. The latter is plastically deformed forward in the 'plowing' mode forming the 'hill' whereas the 'groove'

is just cavitation due to the the motion of the grain., In this manner, one

does not have to invoke the erosion of the bronze matrix by the workpiece

debris, and the process would thus be dominated by the plastic yield strength

of the matrix.

VAN GROENOU

Would this mechanism not depend on the force exerted on the diamond - or

on the workpiece? The force would increase from zero at no contact until a maximum value when the diamond is lost. We have found that the grooves start

rather early in the life of a diamond, when it protrudes about 15 to 20 p,m above the binder. I am not sure that such a particle is taking part in the

grinding. Moreover, the rate of change of depth of the groove seems to be

linear in ferrite volume rate, i.e. force, and not be increasing the more

the particle cuts into the workpiece. The proposed mechanism also raises the

question, why the groove extends behind the particle, sometimes even in a

symetric warp. Would not the distance between a not-yet active diamond and

an active diamond increase, if the latter is pushed into the binder? We did

not notice this.

GIELISSE

Might there be anything unique about your results on ferrites in as

much as there are many examples to quote in which diamonds in wheelapplications do_ show considerable wear before pull out? In any case,

thank you for a very nice piece of work.

VAN GROENOU

The apparatus can be used for many other combinations of binder - diamond-workpiece . The wear that was found here for a ferrite is only moderate,

for harder workpiece materials different results are expected. As regards to

self -sharpening of the wheel, the present combination seems to be a good one.

156
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

NEW CONCEPTS ON SURFACE FINISHING AND ITS APPLICATION TO CERAMICS

RECENT PROGRESS IN ULTRA-FINE FINISHING IN JAPAN

Osamu Imanaka

Department of Precision Engineering, Kanazawa University

Kodatsuno, Kanazawa, Ishikawa 920, Japan

Mamoru Okutomi

Electrotechnical Laboratory, MITI

Mukodaimachi, Tanashi, Tokyo 188, Japan

This paper outlines the ultra-fine finishing techniques developed in Japan.

First, mechnical behavior of a crystal is considered in connection with the "working unit", that is, a bit of stock removal or the range of stress field induced in

the working zone. Since the defects pre-existing in the stressed part affect the

microscopic deformation or fracture of the material, importance of the working

unit as a basic concept in machining or fine finishing should be emphasized. As

an attempt to realize ultra-fine finishing by minute elastic fracture of the lattice size order, the study on "Elastic Emission Machining" is described. Chemical phenomena in abrasive machining process are also dealt with; these have led

the development of mechanochemical finishing techniques, e.g., "mechanochemical

polishing of a hard crystal with a softer powder", and "hydration polishing".

Polishing characteristics for the latter method are somewhat detailed. Additionally, fine finishing techniques of ceramics by electron beams and ion beams are reviewed.

Key Words: Damaged layer; "elastic emission machining"; environmental effect;

hydrate; hydration polishing; mechanochemical polishing with soft powder; preexisting defect; sapphire; solid state reaction; "working unit"; water vapor.

1. Introduction

Material processing could be defined as the procedure which makes a material "memorize" or fix certain desired information by supplying the working energy to it. The

desired information to be fixed is usually the dimensional accuracy or surface roughness.

In some cases, e.g. in manufacturing electronic and opto-electronic devices, physical

perfectness of the finished surface is increasingly needed. As to integrated circuits

in the electronics field, the surface of silicon wafers, for example, have to be finished

to a high grade of flatness, onto which very thin films of precise thickness are coated

and some impurities are doped to precise depths. This is the reason that the crystallographically undisturbed surface, free from chemical contamination is necessitated. This

kind of ultra-fine finishing or machining is required in processing of laser elements,

crystal and ceramic oscillators, and so on. Moreover, in manufacturing mechanical parts

of high precision machines, endeavor is being made to obtain not only geometrical accuracies but also high grade surfaces with the minimum deteriorated layers.

It would be a new trend in precision machining to attain surfaces of higher dimensional accuracy without vitiating the original function of the material. From the emergent needs based on these industrial requirements, new techniques have been investigated

and developed in Japan with substantial results. In this paper, some of these are outlined. Their background, conception and development are also considered.

2. Mechanical behavior of material in connection with the range of stress field

In finishing or machining processes of materials, mechanical behavior of the materials

157
is an important factor that affects the surface texture generated. In abrasive machining,

for instance, the material removal and the surface finish are brought out by a sum of

microscopic deformation or fracture at the point where an abrasive grain acts at the working instant. The mechanism of machining would vary according to the size of stress field

induced in the working zone. The size of the stress field is hereafter called "working

unit".

Yoshikawa [1]1 theoretically studied the effects of the working unit upon the material

behavior, and gave the basis to select the finishing method which is optional for the surface

property required. He classified the working unit into four domains as shown in figure 1 [1]

Each domain is characterized by the lattice defects which preexist or have been generated by

processing. Investigation of the operating predominance of lattice defects for each domain

revealed the following distinct behavior of crystals when impressed by an indenter.

I! I nr w

i

Atom

i

|— Crack

h* Dislocation

Crystal lattice

[Domain]

[Factor affecting

deformation and

fracture]

10" 8 10' 7 10" 6 10" 5 10

_1*10" 3 10" 2 10" 1

Working uni.t, cm

Figure 1. Factor affecting deformation and fracture.

In domain I, material removal of the order of atoms or molecules has been done

generally by chemical decomposition, electrolytic decomposition, evaporation, and ion

sputtering processes. Material removal of this order could not occur by purely mechanical

means. Chemical action activated by mechanical stress may play a role in material processing in this domain.

In domain II, where no dislocations and no cracks exist, a crystal should behave

mechanically as an ideal crystal. Theoretical consideration proved that the impressing

of an indenter upon a crystal surface in this domain gives rise to generation of dislocations prior to brittle fracture. After dislocations have generated, a crystal is

thought to behave as in domain III.

In domain III with only dislocations, plastic deformation occurs first, some amount of

which allow cracks to nucleate in the deformation zone. In domain IV, the crack defect is

the dominant factor. This can be realized by the indentation onto glass surface using indenters with different tip radii. Contrary to the case where a ball indenter with relatively

large radius produces ring cracks on the glass surface, a sharp indenter at relatively light

load yields a plastic indentation mark. The breaking tensile stress to cause the ring crack

is estimated at about 10^50kg/mm^, and the shearing stress under the plastic indentation at

about 40(H800kg/mm^ . The brittle fracture of glass is caused, as is well known, by preexisting surface cracks, which are intrinsically distributed with a density of the order of

10^/cm^ (mean distribution pitch = 10pm). Therefore, if the indentation area becomes smaller

than about lOum, brittle fracture caused by tensile stresses can no longer occur. Figure 2 also shows an example of a relatively large working unit which causes brittle fracture in

high strength material at relatively low stress. Here, the alumina ceramic specimen was attacked by a water jet of 0.15mm diameter at exit pressure of lOkbar [2].

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

158
Figure 2. Fracture of alumina ceramics caused by

water jet of 0.15mm diameter at lOkbar.

10mm

The concept of the above-mentioned working unit might establish some general rules for

selection of machining methods or conditions. In rough machining of ceramic materials, a comparatively large working unit should be adopted to obtain efficient removal rates.

Rough lapping with coarser grains at higher pressure, for instance, is the case, where the

vertical load acting on each grain is higher and the material removal proceeds by microfracturing. In the aspect of working energy, a large working unit was proved to be efficient [3].

In fine finishing, the working unit smaller than the distribution of crack defects

has been selected from experience. Polishing with finer grains on soft polishers demonstrating elastic or viscoelastic behavior at lighter pressures are the examples. To

achieve ultra-fine finishing mechanically, the working unit should be at least smaller

than the dislocation level.

Tsuwa, Osaka University, and his colleagues [4, 5, 6, 7] have examined the possibility

of mechanically machining on the order of lattice dimensions to obtain finished surfaces

which are not only of mirror quality but are crystallographically and physically undisturbed. The fundamental concept is that the collision of fine powder particles with the

work surface, if the size of stress field becomes smaller than the distance between the

pre-existing defects, may cause the elastic fracture of the order of the lattice size and

may not leave a plastically deformed surface layer. The density of pre-existing dislocations has a certain limit and is of the order of 10^/cm^. Thus, the distance between

them is of the order of 1 um. The crystal structure and the strength of material between

the defects can be assumed as the ideal ones. If the fracture is introduced mechanically

in the minute ideal crystal parts between defects, and if the pre-existing defects do not

act, elastic fracture may occur which destroys the ideal binding of atoms. This conception led the new finishing method named as "Elastic Emission Machining" (EEM). The

principle of EEM is schematically shown in figure 3 [4].

In this case, a doubt may arise on the feasibility of EEM, that is, when powder particles

collide with the ideal crystal parts, new dislocations and cracks may be induced. However, the

emergence of dislocations and cracks in these ideal crystal parts may mean that the tensile or shear frac.ture is overcoming the ideal strength. This suggests that ideal elastic tensile

fracture or shear fracture of the order of the lattice size can be introduced by controlling the

magnitude, shape and orientation of stress field, for example, by shooting the powder particles

almost parallel to the work surface.

3. Ultra-fine finishing by utilizing minute elastic fracture

159
Figure 3. Model of elastic emission machining.

I

The collision of powder particles with the work surface was done in the following way:

vibratory collision (fig. 4 [5]), circulatory collision using rotary disk (fig. 5 [4]), circulatory collision using gas stream (fig. 6 [4]), and charged particle collision using

electrostatic acceleration (fig. 7 [7]). Factors governing the machining characteristics

of EEM are the density of pre-existing defects in the material, their mobility, elastic

properties and binding energy of the material and so on. As for powder particles, the

factors are size, angle of collision, kinetic energy and so on.

In figure 4, powder particles are loaded on the workpiece in a container. Horizontal

vibration of the container (frequency 30 Hz, amplitude 4mm) gives the mutual displacement

of the workpiece and powder particles. The surface of the workpiece is intended to be

finished by the collision with powder particles in the horizontal direction. The vertical

load can be taken as the weight of each particle itself. Mean kinetic energy given to a single

particle is calculated to be 4.3*10~3

, 2.6xl0

-

^ and 4.0xl0-6

erg for alumina particles of 20,

8 and 2um diameter, respectively. Assuming that the kinetic energy is converted to the sur- face energy at the time of fracture, and that the surface energy is of the order of lOerg/cm^,

it becomes possible to fracture the area of 200, 13 and 0.2um"% respectively. It was proven

that the electron diffraction image of a Si single crystal surface finished with 2um particles is composed of clear spots and Kikuchi lines.

Figure 5 shows an apparatus made for ductile materials. In the figure, rotation of

the disk accelerates the powder and H gas in the container. Experiments were done with

0.1 um zirconia at 50 m/s peripheral speed of the disk. Aluminum single crystal was reported to be finished fairly successfully.

In Figure 6, particles are enclosed in a closed circuit and circulated with air by means of a blower. From the electron diffraction patterns of the finished Si surfaces with

zirconia powder of 0.1 and 0.6 um diameter at the collision angle (a) of 20°, 30°, and 45°,

it became evident that the smaller the particle size is and the smaller the collision angle,

the better crystalline property is retained. Namely, the extent of stress field and its

direction at the time of collision are related to the quality of finished surface.

160 j
Figure 5. Schematic diagram of rotary

Saw teeth EEM apparatus.

Mixed fluid of powder

and N 2 gas

^

^Workpiece

Workpiece holder

Mixed fluid of

Figure 6. Schematic diagram of EEM

apparatus using blower.

Mixed fluid of powder and air

To realize EEM using electrostatic acceleration, a new type of powder particle beam

gun was developed 16,7]. The apparatus shown in Figure 7 was reported to accelerate the

tungsten particles of 5 um diameter to the velocity of 570 m/s at the accelerating voltage

of 100 RV. The beam density of 10^ particles/s . cm^ was obtained.

T

Charging

^ coil

Powder gun

Electrostatic

lens

Accelerating

electrode —

Machining

chamber

t

Figure 7. EEM apparatus using electrostatic

acceleration.

10" 5 ~10" 6 Torr

To vacuum pump

161
Recently Suzuki and Watanabe [8], Musashino Electric Communication Laboratory, Nippon

Telegram and Telephone Public Corporation, have developed a surface polishing machine

utilizing the principle shown in figure 5. As shown in figure 8, workpieces are floating on a rotating annular disk. Slurry lying between them collides with the work surface.

Adding some chemical agents in the slurry, chemo-mechanical polishing is also done.

4. Chemical phenomena in abrasive machining

As the background of the development of new finishing methods described in the next

section, this section briefly deals with chemical interactions between the work materials,

grain and machining environments.

When thread-grinding a glass bobbin for a standard inductor to be used in the absolute measurement of electric resistance, one of the authors experienced the importance of

chemical phenomena in abrasive machining [9]. At that time, the wear rate of alumina

grains on the grinding wheel used was recognized to be higher than anticipated, in spite of

the facts that AI2O3 is much harder than the glass. The wear was dependent on grinding

environment. The same tendency was also observed in grinding steel. Thorough investigation confirmed that the chemical phenomena at the interface played a significant role in

the wear of hard materials. That led to the development of the mechanical polishing method,

[10,11]. This novel surface finishing technique will be detailed in the next paper in

this symposium.

For fundamental understanding of the effect of machining environment or of fluid on

the fracture behavior of abrasive grains, the fracture strength of alumina grain was measured by quasi-static loading and by repeated impact loading in various environments [12].

Main results obtained were as follows: The strength of the grain may be lowered with the

increase of surface tension and dipole moment of surrounding liquid medium. This can be ex- plained mainly by the lowering of the surface free energy of the grain by adsorption of sur- rounding fluid on their surface or their surface being wetted by the liquid. As to the effect of superheated water vapor in a stress field, as shown in figure 9 [12] , water vapor

lowers the fracture strength as its temperature is raised. This effect can not be explained

by physical absorption, for, in physical adsorption, the amount adsorbed which must be re- lated to the strength of the specimen, becomes smaller as the temperature is raised. The

prior exposure of a specimen to superheated steam under an unstressed condition (aging)

caused an irreversible change in its fracture strength, which suggests that a structural

change of the specimen surface has occurred. In fact, x-ray analysis of finer powder aged

in superheated steam for a considerable time showed the formation of alumina hydrates.

As an example of the formation of alumina hydrates under a stressed condition for a short time duration, Figure 10 113] is given, where the grinding wheel used was WA46N-nV

and the temperature of water vapor sprayed, 100 °C. Under grinding environments containing

water vapor, various diffraction lines of alumina hydrates (gibbsite, boehmite, and diaspore) are seen. The formation of alumina hydrates can be considered as the reaction between the

activated alumina surfaces and water vapor molecules. On the worn alumina grain, surface

162
deposits and metallic layers were not observed; the smooth wear surface was found. Moreover, the hydrates were also identified from the x-ray analyses of the ground chips

collected. These results inspired the development of the "Hydration Polishing" method to

be mentioned in the next section.

cn

re

S-cn

4-O

T3

(OO

S-

-(-> o

n3

S2.5 r

2.0

1.5 -

1.0

Air

0 100 200 300

Temperature of environment, °C

Figure 9. Effect of superheated water

vapor on fracture load of AI2O3 grains.

Before

Grinding

o

I

Oi

25

After grinding

in air

30

After grinding

in heated

water vapor

a

o

co

i

I I I . ,

I • - 35 40 45 50 55°

UL—\_

. a a s | • a

a a a

a

<

•< <

I L- _l I I 1_ -I I I I 25 30 35 40 45

Diffraction angle (26), Co-Ka

I I I 1 L—l 50 55°

A, A1 20 3 ;

A-H(b) , A1 2 0 3 -H 2 0;

A-H(d), 1/2(A1 2 0 3 .H 2 0);

A-H(g), A1 2 0 3 .3H 2 0;

Sp, FeAl 2 0it;

Wu, FeO;

Ma, FC30I+;

Q» Si0 2 ;

*, Al 20 3 -Fe 2 0 3 s.s.

Figure 10„ X-ray diffraction patterns of A1 2 0 3 wheel surface after grinding mild

steel. Grinding distance was 33.6m.

163
collected. These results inspired the development of the "Hydration Polishing" method to

be mentioned in the next section.

5. Ultra-fine finishing by utilizing mechanochemical phenomena

5.1 Mechanochemical polishing with soft powder

This newly developed finishing technique [10,11]> is based on mechanochemical reactions at real contact points between the work and powder, which is mechanically softer than

the materials to be worked, and can react chemically with them. The mechanism and application of this technique will be detailed in the next paper in this symposium.

In addition, Namba and Tsuwa of Osaka University [14] have recently reported the surface characteristics of sapphire finished with various soft powders on a tin polisher. Ion

microprobe analysis revealed the existence of iron atoms or ions on the sapphire surface

finished with Fe2C>3 powder as shown in figure 11 [14], where ions of carbon and hydrogen are contaminations. Figure 12 [14] gives the depth profile of -^Fe+ . The curve does not show

the exact chemical composition due to the characteristics of this analysis, but shows that

iron exists in the first few monolayers of the finished sapphire surface. In the case of

finishing with LiF powder, the first few atomic layers (less than loA) on the finished surface were proved to contain atoms constituting the used powder. These results may suggest

the occurrence of solid state reaction between the work material and the powder.

ro

CDC

+

O)

inO

toc

CD

Primary ion:

Ar+ 7keV

0.8mm diameter

Electron

bombardment:

80ym

m/e

10 20 30

o Depth from surface (A)

40

Figure 11. Ion microprobe analysis of sapphire surface finished with Fe2C>3 powder.

Figure 12. Depth profile of 56pe+ on sapphire surface finished with Fe2C>3 powder,

obtained by ion microprobe analysis.

164
5.2 Hydration polishing

This surface finishing method has been developed on the basis of hydration phenomenon,

occurring by the action of water vapor molecules on the stressed surface of the hydrophilic

material to be polished [15,16]. Its feature is to use neither powder nor chemical etchants. Polishing is done in the environment containing superheated water vapor. Hydrates

formed on the activated surface are rubbed off by the friction of the polisher. If the hydration occurs only at the true contact point between the work and the polisher, and if only

the hydrates are removed from the work surface, this method may fulfill the ultra-fine finishing of the order of molecular size and yield the surface with the least possible defects.

Figure 13 shows the apparatus used in the experiment. A lapping machine is enclosed

in an electric furnace. To the apparatus, saturated water vapor can be introduced from

autoclave. The temperature of the introduced steam can be equalized by heaters to a cer- tain desired value from 100 to 200°C. The specimens used were mainly synthetic sapphire,

which were lapped with SiC abrasives of #4000 grit size prior to every polishing test.

Inlet of Heater Pressure gage

Figure 13. Schematic diagram of hydration polishing

apparatus

The polishing rate of sapphire is shown in figurg 14, where the polisher materials used were straight-grained cedar wood 100 ^ 210 kg/mm , average width of annual rings,

4 mm) and carbon for electric use (#v , 380 kg/mm2

) . Polisher materials in hydration

polishing are desired to be inert to water vapor and to be undeformable under hydrothermal

conditions. In the figure, the polishing rates with carbon polisher are seen to be

higher in the lower temperature range than those with cedar wood polisher. This would

be due to the difference of the hardness between the polishers.

CDE

<so

CL

CO

V) 4-O

<U

4-> raso

2 r

Polishing pressure, 0.5 kg/cm2 Revolution of polisher, 44rpm

Stroke of specimen,!. 9m/min • = Cedar wood polisher

o=Carbon polisher

Figure 14. Effect of water vapor

temperature on polishing rate of

sapphire

50 100 150

Temperature, °C

200

165
The surfaces obtained with the different polishers are shown in figure 15. The

surface with cedar wood polisher is very smooth and free from scratches. In the case of

carbon polisher, carbon layers were observed to adhere to some sites on the polished

surface in the high temperature range. The adhesion would affect the polishing rate. On

the surface with mild steel and silica glass polisher, are seen small scratches (fig. 15-b)

and many deposits which may be considered as the fine debris removed from the polisher

during processing (f ig. 15-c) . Mechanical polishing with diamond paste leaves the deep

scratches on the surface as shown in figure 15-d.

(c) (d)

Figure 15. Sapphire surface finished by different processes: (a) hydration

polishing with cedar wood polisher, (b) hydration polishing with mild steel

polisher, (c) hydration polishing with silica glass polisher, (d) conventional diamond polishing at room temperature.

Figure 16 shows the sapphire surfaces reacted under a controlled partial pressure

of water vapor. The analysis of the aged specimen and the hydration-polished surface by means of Raman spectroscopy Cwith reflected Ar laser) revealed the following: The

spectra of H-0 is seen in the range of 3000 ^ 4000 cm

--'- on the aged surface, as shown in

figure 17. On the polished surface, the H-0 spectra were not identified, but appearance

of fluorescence scattering was found. The results obtained iwth ESCA proved that the

hydration polished surface of sapphire shows little chemical shift in Al levels. From

the results mentioned above, it might be concluded that the sapphire surface hydration

polished is clean from the hydrates and hydroxyl groups and shows little disorder of

crystal lattice. These phenomena would be the minimal effect of the hydroxyl group

coordinate bond to the Al-O., surface. The present technique could also be applied to

final polishing of hydrophilic crystals.

166
lOym

Figure 16. Reaction products on the sapphire surface aged in water vapor at

135°C, 2.5kg/cm2 for 20hr.

6. Surface finishing by non-traditional methods

6.1 Ion beam sputter-machining

Ion beam sputtering is basically a knocking-out of atoms composing the work material

due to bombardment of inert gas ions with high kinetic energy. That is, the bombarding ion

makes sequential elastic collisions with atoms in the target material and knocks out the

atom from its surface layer. Therefore, the sputtering is essentially a non-thermal

process, and can remove the stock of the work material by the order of atomic size.

Figure 17. Observation of sapphire surface

by Raman spectroscopy.

1 . Polished surface

2. Aged surface with water vapor

3. Water

167
From this point of view, Taniguchi and Kanekama, the Institute of Physical and

Chemical Research, started the study on ion heam sputter-machining 117] and developed an

ion sputter-machining apparatus for aspherical surface lens 1 18] . A duo-plasmatron type

ion gun was used as the ion source. The position and the incident angle of the ion beam

to the work were set with mechanical devices by sliding, rotating, and swinging mechanisms.

The surface roughness of the lens obtained was reported to be of the order of A/100 (A,

wavelength of mercury lamp)

,

6.2 Finishing of sintered ceramics by electron beam.

In mechanical lapping of sintered ceramics, fine finishing is not possible due to the

micro pores existing in the materials. Abrasive grains would be left in the pores. In

order to solve these problems, Taniguchi [19] examined the thermal finishing of ceramics of

high melting point by using an electron beam of a relatively wide diameter. Surfaces of

beryllia ceramics were finished to below lym roughness under the following conditions: the

accelerating voltage, 12kV (max) ; mean beam current, 10mA (max) ; spot diameter of beam,

0 . 2mm.

7. Conclusion

Higher integration density of electronic and opto-electronic devices has characterized

the trends of "electronic age" in the world. This has constantly called for higher precision processing in machining technology. Ultra-fine finishing to attain the surface of

high dimensional accuracy without vitiating the original function of the material, is con- sidered as one of the most important problems in this field.

The techniques introduced in the present paper, being developed very recently, cannot

claim to be firmly established as a generally-accepted technology. However, a higher

grade of surface finishing, unattainable by the conventional methods, is expected to be

accomplished through further investigation on the basis of these novel techniques.

Additionally, new techniques have been often developed in the course to solve some troubles. The proverb "Turn a misfortune into a blessing" is thought to be true.

References

(1) Yoshikawa, H. , Brittle-ductile behavior of crystal surface in finishing, J. Japan

Soc. Precision Engg. 35 (10), 662-667 (Oct. 1967). In Japanese.

(2) Imanaka, 0., Fujino, S., Shinohara, K., and Kawate, Y., Experimental study of

machining characteristics by liquid jets of high power density up to 10 W l.to~2, Proc.

1st Intern. Symp. Jet Cutting Technology, pp. G3/25-35, Univ. Warwick, England,

April 5-7, 1972.

(3) Taniguchi, N., Analysis of mechanism of various material-workings based on the con- cept of working energy, Sci. Papers Inst. Phys. Chem. Res. 61, (3), 79-91 (Sep. 1967).

(4) Tsuwa, H. and Aketa, Y., Elastic emission machining as an ultra-fine finishing method,

Proc. Intern. Conf. Production Engg., Pt. II, pp. 33-38, Tokyo, Aug. 26-29, 1974.

(5) Mori, Y., Tsuwa, H. , and Sugiyama, K., EEM (elastic emission machining) (1st Report)

Concept of EEM and its feasibility , J. Japan Soc. Prec. Engg. 43, (5),

542-548 (May 1977). In Japanese.

(6) Tsuwa, H. , Aketa, Y., and Ide, T. Studies on powder particle beam machining by

electrostatic acceleration (1st Report), ibid. 38, (8), 659-666, (Aug. 1972). In

Japanese

.

2 Now, at Tokyo Science University

168
(7) Ikawa, N., Mori, Y., and Ide, T. , Studies on powder particles beam machining by

electrostatic acceleration (2nd Report), ibid. 43, (3), 542-548 (March 1977). In

Japanese

(8) Suzuki, J. and Watanabe, J., Study on polishing without contacting between work and

polisher, Preprint at Annual Meeting of Japan Soc. Prec. Engg. pp. 213-214 (Oct. 1977).

In Japanese.

(9) Imanaka, 0., et al., Precision machining of standard inductor bobbins (1), Bull.

Electrotechnical Lab. 19, (10), 736-742 (Oct. 1955) . In Japanese.

(10) Yasunaga, N., Obara, A., and Imanaka, 0., Mechanochemical polishing of single crystals with soft powders, Proc. Intern. Conf. Production Engg., Pt. I, pp. 32-37,

Tokyo, Aug. 26-29, 1974.

(11) Yasunaga, N. and Imanaka, 0., A new polishing technique using the mechanochemical

phenomenon, Technocrat 8, (9), 15-21 (Sept. 1975).

(12) Imanaka, 0., Fujino, S., and Shinohara, K. , Effect of environments on fracture

strength of aluminum oxide grain, Bull. Japan Soc. Prec. Engg. 2_, (1), 22-30

(April 1966).

(13) Okutomi, M. and Imanaka, 0., Chemical phenomena in grinding process, Proc. 3rd Intern.

Conf. Prod. Engg., pp. 296-301, Kyoto, July 14-16, 1977.

(14) Namba, Y. and Tsuwa, H. , Ultra-fine finishing of sapphire single crystals, Paper presented at 27th General Sssembly of CIRP (New Delhi, Aug. 1977).

(15) Okutomi, M. and Imanaka, 0., "Hydration polishing" of sapphire, Bull. Japan Soc. Prec.

Engg. 10, (4), 156-158 (Dec. 1976).

(16) Okutomi, M. and Imanaka, 0., A novel finishing technique using hydration phenomenon,

Proc. Intern. Conf. Prod. Engg. Vol. 2. vi 79-86, New Delhi, Aug. 1977.

(17) Taniguchi, N. and Kanekama, N., Ion-beam sputter machining of glass, Annals CIRP 2_2,

(1), 69-70 (1973).

(18) Taniguchi, N. and Kanekama, N., On the ion beam sputter-machining of aspherical sur- face lens, Bull. Japan Soc. Prec. Engg. 7, (1), 29-30 (March 1973).

(19) Taniguchi, N., Surface finishing by electron beam, Preprint at Annual Meeting of

Japan Soc. Prec. Engg., pp. 159-160 (April 1963). In Japanese.

169
j
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

MECHANISM AND APPLICATION OF THE "MECHANO CHEMICAL

POLISHING METHOD USING SOFT POWDER"

Nobuo Yasunaga, Noboru Tarumi, Akira Obara

Electrotechnical Laboratory, MITI , Mukodaimachi,

Tanashi, Tokyo 188, Japan

Osamu Imanaka

University of Kanazawa, Japan

Polishing mechanisms of the mechanochemical polishing method

using soft powders are discussed by estimating the reaction rates

at interfacial contact points from static reaction experiments of

mixed powder. Considering the case of polishing sapphire by Si02

powder, reaction process of a-Al i 03-SiOi powder was examined at

high pressures (36, 50 kbar) and high temperatures (630~1230°~C)

which were supposed to be generated at real contact points in

polishing operation. The following are the main results obtained.

(1) Reaction rates for kyanite formation at high pressures were

much higher than those for mullite formation at atmospheric pressure. (2) The theoretical polishing rate estimated from these

reaction rates approximately coincides with the one measured in

the actual processing of sapphire by Si0 2 powder. (3) Polishing

of sapphire could proceed only by mechanochemical reactions at

contact points.

Mechanochemical polishing of Si single crystals and quartz

crystals were tried in order to extend the applicability of this

surface finishing method to other hard materials besides sapphire. It was disclosed that damage-free mirror surface could

be obtained by mechanochemical polishing using powders such as

BaCOj, CaC03 etc. for Si single crystals and Fe 3 04., MgO etc.

for quartz crystals.

Key words: High pressure; high temperature; mechanochemical

polishing; quartz crystal; sapphire; Si single crystal.

1. Introduction

High precision surface finishing technologies have been required

especially in the electronic industry. The "Mechanochemical polishing

method by soft powders" developed by the authors is ideal for this purpose,

because very flat and damageless mirror surface can be easily obtained by

using this new technique, which can be carried out in the same way as conventional mechanical lapping and polishing methods. The main distinction

of this surface finishing method is to use powders which are mechanically

softer than the materials to be polished and chemically reactable with them,

for example, Si02 , iron oxides, MgO powder etc. for polishing sapphire.

As the soft "abrasive" powders neither mechanically indent nor scratch the

material surface, smooth and scratch-free finished surface of hard crystals

can be obtained, even if powders of large grain sizes and relatively rigid

polishers are used.

171
The conception of this polishing method was derived from the investigation on wear mechanisms of sapphire sliding on softer materials. The

wear experiments of sapphire sliding on a rotating quartz glass ring have

revealed that amorphous mullite was produced at the real contact points

between the sapphire surface and glass debris. In the wear test.' of sapphire against steel ring, solid solutions of FeAl 2 04.-Fe 3 04 and a-Al_203 - Fe 2 0 3 formed. In both cases the worn sapphire surface was so smooth that

the wear of sapphire was understood to proceed fundamentally in consequence

of the exfoliation of very small reacted parts at the real contact points.

The positive usage of these wear mechanisms yielded the new mechanochemical

polishing method.

The evaluations of machining efficiencies and surface observations

have been made in detail about sapphire polishing(l}1

. It is, however, very

difficult to detect reaction products and to determine the reaction

processes in the actual polishing operation. In this report, material

removal processes at the polishing interface are estimated by means of reaction experiments of mixed powder at high pressures and high temperatures

which are supposed to be generated at the real contact points. Application

experiments of this machining technique to other crystals such as Si single

crystals and quartz crystals are also described.

2. Experimental

Polishing experiments were carried out using a lapping apparatus

schematically shown in figure 1. In a round holder A which rotates independently of the polishing plate, 3 specimens are mounted. Using holder B,

unidirectional polishing tests with an oscillation of about 20 mm in the

radial direction are done for a single specimen. Before polishing experiments, every specimen oriented by X-ray back reflection Laue method, was

lapped with fine abrasives and furthermore etched by 36 0" C H3PO4 for

sapphire and CP-6 (HN03 iHF: CH 3 C00H=2»1:6 ) for Si single crystals in

order to eliminate the surface damage produced by the mechanical lapping.

Load

i—1 /Holder 0 Holder (5)

Self- aligning ball bearing

Polishing plate

Strain gauge

20mm^~

Powder

receiver

250mm

Specimen Specimen

Figure 1 Schematic diagram of polishing apparatus.

For reaction experiments of mixed powders, a girdle-type high pressure

apparatus was used, which is capable of generating up to 80 kbar at

room temperature and up to 50 kbar at 1000° C. The schematic arrangement

is shown in figure 2. In consideration of the polishing of sapphire by Si0a powder, mixed powder of a-Al2 03 (average grain size 2 um) and Si0 2 (0.5 um)

was prepared for reaction experiments (the mole ratio was 1«1).

The powder sample was heated in a graphite tube. The temperature of the

every sample was measured by a thermocouple ( Pt-13%PtRh, 0.5 mm^).

The pressure generated in the sample was previously calibrated with the

Bil-2, T15-I, Bal-I and Bi¥-H transitions. After the reaction experiments,

the samples were milled and analyzed by means of X-ray diffraction method.

1 Figures in brackets indicate the literature references at the end of this

paper.

172
(a) High prMMirt •pportut ( b) Sample OM.mbly

Figure 2 High pressure apparatus and sample assembly

3. Mechanism of Mechanochemical Polishing of Sapphire

3.1. Polishing rate of sapphire by Si02 powder.

In wear experiments of sapphire sliding on fused quartz glass ring,

which was conducted in order to investigate what kind of reactions would

occur at the interface and how the products would be removed, the following

results have been obtained £2] . The wear rate of sapphire strongly

depends on crystallographic orientation as shown in figure 3.

Especially in the ( 0001 )+5

6 d irection the wear rate is high, whereas very

low in the opposite (0001)-5° direction of the same slide surface and

other sliding direction. This peculiarity is the same as Duwell has already reported 13 i « The wear debris collected was amorphous to X-ray for

all sliding condition and crystallized to form mullite when heated above

ca.950°C as listed also in figure 3» From X-ray measurements of the grain

size of the crystallized mullite and of the reaction process of the debris

at higher temperatures than that of crystallization, fig 4 was proposed as

the appropriate removal model at the contact points (.2} . In the (0001)+5o high wear direction a small region of sapphire where basal sheets could

easily slip toward the surface is rubbed off, being joined to the thin

reaction layer (fig. 4(a) ). In the other directions, where sapphire could

hardly deform, only 9 the reacted part (of which size was estimated to be

under several tens A) exfoliates as seen in figure 4 (b).

-5°

direction + 5° direction Sliding

direction

Wear rote

( m»/ km )

Products in wear debris after heat treatment (980*0

(OOOU + 5* 5.75 mullite + ff-AhOs

(0001 )-5* 0-55 mullite

( 1 0T2)<I2T0> 0.50 mullite

( 1 2 I0)<I0T0> 0-09 mullite

Sliding condition ; Load 0.5kg , Sliding speed 2 4

m/min Relative humidity 100%

Figure 3 Wear characteristics of sapphire sliding on quartz glass

173
Figure k Removal model of

wear debris

It is technically very difficult to detect the interfacial reaction

products in the actual polishing operation. As an expedient for confirming the possibility of mechanochemical reactions during polishing, a

wear experiment of sapphire sliding on a quartz glass disk (polisher) was

carried out on the same sliding condition as in the actual polishing

operation, i.e. at the load of 0.5 kg and o the sliding speed of 20 m/min

under water lubrication. In the (0001)+5° direction the debris collected

was also amorphous and the same phase transition that had been seen in the

wear experiments mentioned above was observed (fig. 5). This fact means

that even in the actual polishing process mechanochemical reaction could

occur at real contact points.

(a) 970*C, 2hours

-*— 1 1 1 . 1 , / t—' L_ 15 20 25 30 35 40 45 65 70"

2 6 ( CuK. )

Sliding condition : l 000 1) + 5* direction. 0.5 k«/cm\ 20m/min. water supply

Figure 5 X-ray diffraction patterns

of heat-treated wear debris

In this mechanochemical polishing, the machining efficiency is much

higher in dry conditions (without fluid) than in wet conditions, because

chemical and thermal activities at the contact interface are expected to

be_higher in dry conditions. Figure 6 shows polishing rates in (1210)<

1010> direction which is known as the most undeformable direction of

sapphire (holder B was used for this experiment). Higher reactivity at

the contact interface at 20 m/min would have resulted in the higher polishing rate than at 10 m/min. Figure ? shows, the orientation dependence of

polishing rates. The highest polishing rate is obtained on the (0001)

plane which most easily deforms plastically and the lowest rate on the

plane where the machining is supposed to proceed idealy based on the

removal mechanism of figure k (b).

174
0.8

E

4) .1 0.6

-ca

a

a

0.4

O) ~ 0.2

of

Polishing

direction : ^ 1210 )<10T0>

Powder : 0.5//m Si02

Polisher ; Quart glass

0.5 1.0 Load

Figure 6 Polishing rates of

sapphire in dry

cond ition

1.5 (kg

2.0

E~ 2

4)O

w

CO 1

1

Number of rotations

Polisher : 40rpm

Holder A : 3 7rpm

y Load : 1. 0kg

11

0.02jjm Si 02

5pm Si 02

1L

Figure ? Polishing rates of

sapphire polished with holder A

(0001) ( I0T2 ) (1010) (1210)

Sapphire plane to be polished

3.2. Reaction rates of mixed powder at high temperatures

and high pressures.

It is very important to measure the reaction rates at contact points

in polishing operation, in order to verify the contribution of chemical

reactions to this mechanochemical polishing. For this purpose, reaction

experiments of mixed powder were conducted on the same conditions (temperatures and pressures) which would have been generated at the real contact

points between sapphire surface and SiOa powder.

The plastic flow pressure pm at a real contact point would be nearly

equal to the hardness Hv of the softer material at that temperature. The

pressure generated at a contact point between sapphire surface and Si02 powder is therefore estimated from the hardness of quartz glass. The pm

is for example about 38 kbar at 800°C and 50 kbar at 600°C. In this report, reaction experiments were done on some of these "critical conditions"

and on some of higher conditions than the critical ones.

175
Figure 8 shows the X-ray diffraction patterns of a-Al,0^-SiO, mixed

powder after 10 sec reaction at 1030" C and 36 kbar. As seen, a reaction product formed kyanite which is the stable phase of Ala03 -SiOa system in the

high pressure region. The amorphous SiOa changed to coesite in this condition. The relation between the fraction a of kyanite produced and the time of reaction at 36 kbar is shown in figure 9. It is notable that

very high a (0.3-0.5) was obtained even by rapid cooling just after rapid heating to 1230° C (the existing time above 1000° C was about 2 sec).

In comparion with the very low reaction rates of this system at the atmospheric pressure (products were little observed even after 10 hours heating at 1200 C), the rate at a high pressure is extremely high.

o

+ o kyanite a a-AltOs

a.- quartz

coesite 36 kbar

I030°C

10s

oC=0.39

SiOz :<X-Al2 0s= I 36kbar

I230*C

20

Figure 8

25 30 35

2 0 ( Cu Ka )

40 45 50*

X-ray diffraction pattern of

Si02 -ot-Al2 O3 mixed powder after

the reaction at 36 kbar

s 10 15

JT ( t : Reaction time, sec )

20

Figure 9 Fraction a of kyanite

formation related to ,/t

The SEM image of the as-reacted powder reveals that the kyanite

grows to columnar and plate-like structure (fig. 10). As a suitable contact model of the individual spherical powder for such a reaction form,

the partial contact model like figure 11(a) should be adopted instead of

the one of whole surface contact like figure 11(b). It is known that

the thinkness / of the reacted layer and the reaction time t can be connected by the parabolic law
la) Reaction model for

partial contact

(b) Jander's model

j>=K't (1)

where K' is rate contant. If the fraction of the contact area to the

total surface area is represented by f, for small f/rB (rB is the radius

of the larger particle) a is given by

J a=47irg.f.//|iirj=3f//rb (2)

From equations (1) and (2) is given by

a

2 =K M t (3)

where K"=9fz K'/r|. When F=l and a«l, the equation (3) agrees approximately with the Jander's equation { 1- (1-a )*3 }

a =K' t/rg . In figure 10,

the relation of a to Jt can be regarded as linear up to a=0.3~0.i|..

The K'1 in the equation(3) can be obtained from the initial gradient

(the broken line) of the each curve. The values of K" for each temperature at 36 kbar and 50 kbar are listed in figure 12(a). Then, the time

t' necessary for formation of reaction layer of 50 A was calculated

assuming F=l/3 and r8(of a-Al^Os particle)=2 um. The results are shown in

figure 12(b).

(o) k" for kyonite formation ( 0Cl « k"t )

"""----^Pressure

Temperature""""-----^ 3 6 k bar 5 0 kbar

6 30*C 2.9 X 1 0"6

7 80*C 7.2 X 1 0" 4

8 30*C 6.3 X I0"

4 1.7 X 1 0" s

I030*C I.I X 10'* I230*C 9.8 X I0"2

lb) t ' corresponding to f/^ «=0005/2 tin the co»e of f

= l

^^^Pressure

Temperaturi~~~-

3 6 kbar 5 0 k bar

6 3 0*C 2.2 780*C 8.6 X 1 0' S

830*C 1.0 X 1 o"2 3.7 X 1 0*'

1 03CTC 5.7 X 1 0" 4

1 2 30*C 64 X l 0" 8

Figure 12 Reaction rates of SiOj-ct'AliOj

mixed powder at high pressures

and high temperatures

177
Connecting these results with the polishing process, the following

interpretation would be possible. Assuming the temperature generated at

a contact point to be 830°C, the pressure generated there would be about

36 kbar. Assuming further that the reaction layer is rubbed off from

the surface when the thickness/1 reachs 50 A (the actual reaction layer

at the contact points is presumed to be thinner than 50 A), the time

required from the beginning of contact till rubbing off should be about

lxlCf* sec.

3.3* Estimation of the polishing rate from the reaction

experiments.

The actual contact points in a polishing operation are considered to

be fresh and chemically very active, compared with the surface of the

powders contacting in the static reaction experiments. Besides, the

value of f/rB in K" of the equation (3) is supposed to be smaller in an

actual reaction than in the assumption adpoted for the above calculation. Taking these circumstances into consideration, the actual contacting time could be shortened to about one- tenth of the t* in figure 12

(b).

A polishing rate would be now estimated on the assumption that a

temperature of ca.830°C and a pressure of ca.36 kbar arise at the sliding contact points

3and reaction layers 50 A thick are produced and rubbed off within 1x10" sec. The volume v of each exfoliated part is given

by

v=W-h/p m (2+)

Where W is the load, p^ is the plastic flow pressure of SiOz and h is

the thickness of exfoliated layer. % If the volume v is rubbed off in

every time t', the polishing rate w per unit time is

w=d.v/f =d.W.h/pm« f (5)

Where d is the density o of the reaction product. When w=1.0 kg, pm=36

kbar=360 kg/mm2

, h=50 A=50xl0"7 mm, t'=lxl(f5 sec and d=3.7 mg/mm», the w

is calculated at 0.18 mg/h. This value is nearly equal to the average

polishing rate on the least deforamble (1210) plane (fig. 6 and 7), where

polishing would proceed approximately on the basis of the simple model

discussed above. The fact that the polishing rate estimated from the

static reaction experiments would approximate to the actual rate proves

that sapphire can be only polished by solid state reaction.

It is known that kyanite is unstable under the atmospheric pressure

and changes to mullite phase above 1325°C. Kyanite also easily changes

mechanochemically to amorphous and mullite phase by light milling in a

mortar £2 J. It would, therefore, be reasonable to conclude that kyanite

forms first at the sliding contact points, and then it changes to amorphous phase by receiving frictional stresses.

4. Application of the Mechanochemical Polishing Method for Si

Single Crystals and Quartz Crystals.

4.1. Polishing characteristics of Si single crystals.

As it is the substance of this new mechanochemical polishing method to

use powders softer than the material to be worked and reactable with it,

this technique should be applicable for polishing other hard crystals

besides sapphire. So, as an example, the possibility of polishing Si single

crystals by this method was investigated.

It is known that the crystal surface of Si is apt to oxidise easily,

so that "solid state reactions between the softer powder and the oxide film

178
formed on Si surface" would be the foundation of Si polishing.

In this investigation, polishing experiments using BaC03 , CaCO^ and Fe 3 0/j- powders was tried. All of these materials, which are softer than Si and

its oxide film, are expected to have mechanochemical effect on Si polishing.

The holder A in figure 1 was used for polishing experiments. 3 specimens of 12 mm by 12 mm by 2 mm thick were polished at the number of polisher

rotations 20 rpm and at the nominal pressure 120 g/cm* The polishing rates

obtained on the (111) plane are shown in figure 13. It is noticeable xnat

the polishing rate by BaCO^ powder in dry condition (without any fluid) is

high enough, when compared with that of the conventional polishing by SiOa and CeOa. powders. BaCOs is reactable with SiOi even at a relatively low

temperature of about 850°C [4], High reactivity like this would have brought

about the efficient polishing. Figure 13 also shows tnat mechanochemical

polishing by BaC03 and Fe^O* in dry conditions gives much higher polishing

rate than in wet conditions (water supply). This feature which has been

also seen in the case of sapphire polishing indicates that high activity at

the contact points in dry conditions brings high machining efficiency in

the mechanochemical polishing.

Powder Grainsize (pm) Polisher Environment

Polishing rote

( tm/h )

BoCO*

3 ~ 1 0

Cloth

Cloth

Dr y Wet

I.I 6

0.1 7

2 ~4 Bakelite

Bakelite

1

Dry

Wet

2.6 0

0.30

F*s 0* IO~20 Cloth

C loth

Dry

Wet

0.68

0.2 8

Co COs

4 — 7 Cloth Dry 0.5 9

2 ~3 Cloth Dry 0.97

2~ 3 Bakelite Wet 1.30

Si Ob

as

Cloth

Cloth

Dry

Wet

1.00

0.85

O.OI 0.02

Cloth

Cloth

Dry

Wet

1.9 1

1.05

C.Oe 2 — 3 Bakelite Wet 0.9 1 Diamond 1 Cloth Dry 2.2~4.2

Figure 13 Polishing rates on Si (111) plane

4>.2. Observation of polishing surface

Etched surfaces of the polished Si (111) plane were observed by a

differential interference microscope (fig. 14). It is evident that little

worked damage remains on the Si surface polished by coarse BaC03 powder,

while cracks and scratches are seen on the surface mechamically polished

by Ce0 2 and SiO* powders.

179
Polished surface Etched surface

( etching time : 10 sec )

lOpnvBoCCte

Bakelite

:

3pm CeOs

bakelite

v

9

3pm Sj02

Bakelite

' ' ""7"—— ****

Figure 1^ Si surface polished in wet condition

The electron diffraction patterns (fig. 15) show that no detectable

defects are remaining on the surface polished by BaC03, whereas no detectable crystalline structures are observed on the surface polished by

diamond paste. Figure 16 shows ESCA spectra of Si2p (binding energy Eb=99

eV) and 01s(E|,=532 eV) on the polished surfaces. The peak at E b =103 eV

means the chemical shift of Si2p caused by surface oxidation. The figures

lOpmBaCOa. Cloth 2pm Diamond . Cloth

Accelerating voltage . 45 kv

Figure 15 Electron diffraction pattern of Si surface polished

in dry

180
Olt Si2p

—i i i i 1 1 1

534 532 530 104 102 100 98

Binding energy ( ,v )

Figure 16 ESCA spectra on polished surface of Si (111) plane

in ( ) on the right side of figure 16 represent peak hight's ratios of the

shifted Si2p peaks (E,j,=103 ev) to the proper Si2p peaks (E p =99 ev).

It can be qualitatively said that the smaller these figures are, the less

the surface oxidation is. Figure 16 indicates that mechanochemical polishing by BaCOs powder brings less surface oxidation than mechanical polishing by fine diamond and SiO* powders. From the ESCA analysis in the low

binding energy region the worked layer was estimated as under about 20 A.

Polishing of quartz crystals

This mechanochemical polishing method was also applied to surface finishing of quartz crystals. It was found that powders such as Fe 3 04, MgO,

MnOi etc. are useful for high precision polishing of quartz crystals.

Especially, polishing by FeaO^ powder in dry condition brings as high machining efficiency as conventional polishing by Ce02 and Fej.0} powders does.

Furthermore, a damage-free surface with high flathess is obtainable by using

a relatively rigid polisher like one made of copper.

5. Conclusion

Reaction experiments at high pressures and high temperatures have shown

that chemical reactivity at contact points is the fundamental factor which

controls the polishing rate. In order to improve the polishing efficiency,

it should therefore be considered how to increase the interfacial reactivity,

namely} it should be investigated what kind of powders and environments

could give the highest polishing rates. This polishing method is fundamentally applicable to various kinds of hard materials as discussed in k..

Selection of powders and machining conditions suitable for polishing of the

respective materials would be very important.

181
References

( 1 ) Yasunaga, N; , Obara, A. and Imanaka, 0., Mechanochemical polishing

of single crystals with soft powders (Proc. 1st Intern. Conf.

Production Engg., pp. 32-37 Tokyo, August 26-29, 1974).

(2) Yasunaga, N., Obara, A. and Tarumi, N. , in Study of mechanochemical

effect on wear and its application to surface finishing, pp. 73-76

(Researches of the Electrotechnical Laboratory No 776, 1977).

13) Duwell, E. J., Friction and wear of single crystal sapphire sliding

on steel, J. Appl. Phys. 21' 9 , 269I-2697 (1962).

Duw',11, E. J., and H. C. Butzke, The effect of interface composition

on tie wear rate of sapphire, ASLE Trans., 7_, 101-108 (1964).

(4) Yamaguchi, T. , Fujii, H. and Kuno, H. Kinetic studies of the solid

State reaction involving intermediate phases system BaCO -SiO , J. inorg. nucl . J34, 2739-2745 (1972).

Discussion

GIELISSE

In the case of BaCO^ on Si is the actual chemo mechanical mechanism

known? Are we getting a reaction between the two components to give a reaction product or does Ba(X>3 break down followed by subsequent oxidation of

Si? We have noticed that the amount of "breakdown oxygen" available on the

interface influence the reation rate. Are you getting similar results?

YASUNAGA & TARUMI & OBARA

Since the Si surface is apt to oxide easily, the mechanochemical reactions are considered to occur, at the interface between the powder and the Si

oxide film. Yamaguchi et al have showed that BaC03 reacts easily with

Si0

2 to form BaSiOg and Ba2 Si04 above 850° C. So, it can be inferred that

"breakdown" and rubbing off of these reaction products formed at the real

contact points lead to the fine machining of Si surface.

*Ref. 4 in text.

NOVAK

1. If no cooling is used in the experimental apparatus, why doesn't the clot:

or bakelite thermally degrade in a very short time period?

2. What pfl was used during the polishing conditions listed in the slide as

wet?

3. Is the powder continually replenished on the wear surface?

YASUNAGA

1. In these experimental conditions, apparent temperatures generated

on the polishers are not high enough to degrade their function,

although at small real contact points very high temperatures would

be generated in each short contact time.

182
2. In our experiments water from the pipe was directly supplied

and no pH adjustment was carried out. As Gutsche and Moody*

indicated, pH of the polishing fluid might be a important factor,

In our further investigation this point will be taken into

consideration.

3. As the polishing machine was designed so that the powder does

not "get away", replenishment of powder is not always necessary

even in dry polishing.

LEE

In your high pressure reaction study, how did you control the temperature

for the short reaction time you need.

YASUNAGA

In these high pressure experiments the mass of the mixed powder

samples is small (4mmtfx7mm). By supplying the graphite tube heater

an electric current under 200A at about 4 V, the sample can be rapidly

heated and cooled. In this experiment by manual control, rapid heating

from room temperature to 1230°C was realized within a few seconds.

Preparation of automatic control system would make it possible to

realize more rapid and precise control of the reaction times and

temperatures.

H.W^ Gutshe • J, W.Moody t J.Electrochem. Soc, 125, 1(1978)136.

183

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

EXTENDED DISCUSSION: HARDNESS AND ITS RELATION TO MACHINING

R. W. Rice

Naval Research Laboratory

Washington, DC 20390

A. G. Evans

University of California

Berkeley, CA 94720

During the conference, a number of questions arose about hardness and its

relation to machining. As a consequence, time was set aside for a general

discussion; A. G. Evans volunteered to present some specific considerations on

this subject as an aid in focusing this discussion. The following note, subsequently prepared by R. W. Rice and A. G. Evans, represents a summary of both

his specific considerations and the general group discussion. Basically, three

issues were addressed: 1) the meaning and technique of hardness measurements;

2) the appropriate hardness/load values to use in machining considerations;

and, 3) the relevance of hardness to multipoint machining.

1. The first topic, the meaning and technique of hardness measurements,

is an old and complex topic because of the many factors that can affect hardness and its measurement. These factors can largely be classified into the

following three main categories:

a. Load Application Factors — Of these, major importance is ascribed to

the dependence of" hardness on load; the effect of translational motion; and,

the effect of loading rate.

Several factors give an apparent or real dependence of hardness on load.

Actual errors in measurement of the true identation dimensions can occur.

Typically, the dimensions are undermeasured by a constant factor, and hence,

the errors are more serious for small indentations, i.e., harder materials and

lighter loads. Another long known factor giving an apparent load dependence

is elastic relaxation on unloading. It is generally accepted that elastic

relaxation results in a fixed shortening of indentation dimensions and, thus,

again becomes a serious effect for low loads and hard materials. Since elastic

relaxation also increases with Young's modulus (E) , the effect is more properly

described in terms of the ratio H/E, with the extent of elastic recovery increasing with H/E. In this regard, it is important to note that ceramics, in

comparison with metals, are characterized by relatively large H/E values.

Finally, work hardening of the surface due to machining or abrasive

lapping has been shown to produce an intrinsic load effect [1]. At low loads

the indentation deformation occurs primarily within a work hardened surface

layer, while at higher loads only a fraction of the deformation occurs within

the work hardened layer. This again leads to an increase in measured hardness

with decreasing load. For most ceramic materials, these effects (both real

and apparent) become negligible for loads above a few hundred grams. During

machining operations, normal loads in excess of this range are exerted on the

workpiece, so that any correlation between machining rate and hardness should

not be influenced by these effects. This may not be true for correlations

between hardness and fine abrasive surface finishing operations.

The second consideration, namely the effect of translational motion, is

extremely pertinent to machining applications. Since translational motion is

185
involved in machining operations, it would appear that scratch hardness is a more relevant parameter than the commonly used indentation hardness. Nevertheless, while detailed comparisons of scratch and indentation hardness are

not known, the relative ranking of materials by both techniques is expected

to be similar. Several authors have, in fact, shown good correlations between the Moh scratch hardness scale and indentation hardness [2,3].

The effect of loading rate on hardness is also pertinent to machining

considerations. Again, while only limited information on the effect of loading rate is available, the relative ranking of materials obtained under typical

quasi-state conditions is not expected to be significally altered. For ceram-. ics, the phenomena of indentation creep has been observed and appears limited

(primarily to relatively soft ceramics) though not extensively documented.

Another indication that loading rate is not a significant factor is the successful use of punch testing (i.e., impact hardness) on metals [4]. Furthermore, the static hardness of ceramics appears to be a factor of three, and

possibly less, times the dynamic yield stress (i.e., the shock yield stress

as obtained under ballistic impact) [3,7]. Accordingly, little difference between quasi-static and dynamic hardness is expected, and may be minimal for

materials of high H/E values.

b. Material Factors — Two material factors that effect hardness are

microstructure and the propensity for cracking. The latter depends on material properties, microstructure, and on load, i.e., cracking usually commences

above a threshold load and becomes more severe at higher loads. However,

observations on cracking indicate that the effect on hardness is generally

small. Since cracking that occurs during the indenting process, i.e., during

loading, is confined to a region outside of the plastic zone (the stresses

within the plastic zone are compressive) the effect of such cracking on hardness should be small. Cracking that occurs during unloading can, and often

does, extend into the indentation. This should have little effect since it

normally results in minimal distortion of the indentation or in the accuracy

of determining its dimensions. Some effect of cracking might occur, however,

if cracking is extensive enough to change the effective compliance and hence

elastic restraint of the material around the indentation (e.g., crushing

beneath the indenter)

.

Microstructure, i.e., grain size and the size and distribution of any

second phases or pores, also affects hardness [8]. The extent of the effects

generally depends on the scale of the indentation relative to the scale of

the microstructural features. Of the microstructural parameters, grain size

is the variable of greatest pertinence to the relatively pure and dense materials discussed at this conference. For small indentations, or very large

grains, single crystal values of hardness apply, but should probably reflect

an average over all crystal orientations. For larger indentations or smaller

grain sizes, the polycrystaline value should be appropriate. This may be

smaller than, but is often larger than, the single crystal value. However,

for hard ceramics, grain size effects do not appear to be very pronounced [8].

c. Environmental Effects — Chemomechanical effects have also been found

to affect hardness, mainly at quite low loads, and are at best of uncertain

pertinence to machining processes [9,10]. Lubrication effects of coolants

under the dynamic conditions of machining are also not well understood, but

appear to be limited. Local temperature excursions due to abrasive particleworkpiece interaction apparently can be quite substantial [10]. Although

hardness decreases substantially at the temperatures indicated, the volume of

material with substantial temperature rise should generally be much smaller

than the volume of materials deformed in machining so that any local temperature effect should be fairly limited.

2. In the context of the preceding discussion, the second major topic,

the selection of appropriate loads in hardness determinations, can be addresse

directly; generally, high loads in the kilogram range are considered more

186
appropriate in machining considerations. This can be seen by recalling the

typical hardness formulas:

« = £ (

where a = a constant determined by indenter geometry, P = load, and L = characteristic dimensions, e.g., indentation diagonal, and

where d = depth of the indentation and 8 = another constant determined by

indenter geometry. The hardness value pertinent to machining should generally

be that value representing similar depths of penetration of the hardness indenter and the abrasive particles. One easily obtains the load in terms of

the penetration (depth, d)

Thus, if' one assumes that abrasive particles will penetrate a reasonable

fraction of their diameter (e.g. , «l/4) during a machining operation, one can

calculate the load required for comparable penetration during identation.

As an example, consider Al2O3(H~2000 kg/mm ) abraded by a 325 grit (dia.«40

ym) ; taking d = 10 urn, the appropriate Vickers indentation (a = 1.854; 8=7)

load from eq. 3 is: P«5.3 kg. Thus, one can readily see that even for

smaller abrasive grits (with correspondingly smaller penetrations) and for

softer materials, loads of the order of a kilogram or more are required. At

such loads, the major load-dependent effects are minimal and hardness is nearly

unsensitive to load. One notable exception is the possible effect of surface

work hardening from prior machining steps or prior passes with the same abrasive. Generally, the full depth of slip, D, produced by a given finishing

processes varies roughly as 1/H [11]. For A1 2 0 3 ground with a 325 grit wheel,

D is found to be of the order of 10 ym. Since this is comparable to the

estimated depth of penetration (see above) , the effect of work hardening is

difficult to determine. However, for MgO, D«40 ym, and hence the effect of

work hardening can be of consequence. The effect of work hardening should be

of greatest importance during the early stages of each step in sequential

finishing operations, since the smaller particles must penetrate initially

into the residual damage layer of the preceding step. One of the uncertainties is the extent of work hardening in the deformed layer; a fair portion of

this zone may represent slip on only one or two systems. Also, transient

temperature effects, noted earlier will be greatest at the surface, i.e., in

the region of extensive deformation, and may offset some of the effect of

work hardening.

3. The final major topic, which is concerned with the relevance of hardness to multipoint machining, has two aspects. The first is the good correlation between indentation fracture and the primary mode of natural removal

during machining. In both cases, lateral cracking occurs about a central

plastic compression (or groove) , and expressions for material removal can be

derived from consideration of the extent of lateral cracking. These expressions invariably involve hardness because of the nearly direct relationship

between depth of lateral fracture and depth of penetration [12].

The second aspect of the relevance of hardness to multipoint machining

is the direct correlation of the process and its effects with hardness. For

example, the rates of both sawing and grinding directly correlate with hardness [11,13] as does the depth of slip, D, discussed earlier. Also, in some

earlier studies the rate of lapping was used as an extension of, or replacement for, the familiar Moh hardness scale [14].

In summary, it appears clear that hardness is a basic factor in the

machining of ceramics. Further, although there are many second order effects,

L = 8d (2)

(3)

187
some of them such as work hardening possibly being significant, the major

trends and implications of hardness for machining seem fairly well outlined

by present work. Thus, while hardness depends on a number of variables (such

that some may question calling it a material property), it is fortunate that

most of the inherent problems are not significant at the higher (i.e., kg)

load values that are generally implied for relating hardness and machining.

Further work is important to define in much more detail the hardness -machining

relationships, but a sound foundation on which to build is fairly well

developed

.

References

[1] Becher, P. F., Surface Hardening of Sapphire and Rutile Associated

with Machining, J. Am. Ceram. Soc . , [57] 2_ (1974).

[2] Rice, R. W. , Discussion in The Science of Hardness Testing and Its

Hardness Testing and Its Research Applications , J. W. Westbrook and

H. Conrad, eds., 72-73 (1973).

[3] Gilman, J. J., Hardness — A Strength Microprobe, in The Science of

Hardness Testing and Its Research Applications , J. H. Westbrook and

H. Conrad, eds. , 51-72 (1973).

[4] Li, C. H. and Enos, G. M. , Punch Hardness Testing: Theory and Uses,

Ibid, 300-310.

[5] Shaw, M. C, The Fundamental Basis of the Hardness Test, Ibid, 1-11.

[6] Rice, R. W. , The Compressive Strength of Ceramics, in Ceramics in

Severe Environments , W. W. Kriegel and H. Palmour, I I I~ eds . , 195-227,

Proceedings of 6th University Conference on Ceramic Science, North

Carolina State University at Raleigh (1970)

.

[7] Gilman, J. J., Relationship Between Impact Yield Stress and Indentation

Hardness, J. Appl. Phys . [46] 4, 1435-1436 (1975).

[8] Rice, R. W. , Microstructure Dependence of Mechanical Behavior of Ceramics, in Treatise on Materials Science and Technology , R. K. MacCrone,

ed., 199-381 (1977).

[9] Cooper, G. A., Some Observations on Environmental Effects When Diamond

Drilling (these proceedings).

[10] Gielisse, P. F., Choudry, A., and Ida, I., The Effects of Abrasive

Finishing Processes on Materials (oral presentation, this proceedings)

.

[11] Rice, R. W. , Machining of Ceramics, in Ceramics for High Performance

Appl ications , J. J. Burke, A. E. Gorum, and R. W. Katz7 eds . , 287-343,

Proceedings of 2nd Army Materials Technology Conference, Hyannis,

Massachusetts (1973)

.

[12] Evans, A. G., Abrasive Wear Damage in Brittle Materials (these proceedings) .

[13] Rice, R. W. and Speronello, B. K. , Effect of Microstructure on Rate of

Machining of Ceramics, J. Am. Ceram. Soc. [59] 7-8 , 330 (1976).

[14] Wooddell, C. E., Method of Comparing the Hardness of Electric Furnace

Products and Natural Abrasives, J. Electrochem. Soc. [68], 111-130

(1935)

.

188
Session II

Technology of Machining

and Surface Finishing

Session Chairmen

C. Bersch, Naval Air Systems Command

R. W. Rice. Naval Research Laboratory
i

I
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

OPTIMIZATION OF MULTISTAGE GRINDING OPERATIONS:

THE CHOICE OF CONSECUTIVE REMOVAL RATES

A. Broese van Groenou

Philips Research Laboratories

5600 MD Eindhoven, the Netherlands

The optimum choice of removal rates is determined in a model for

multistage grinding on the basis of remedying the damage introduced in

the previous stage. A set of second order difference equations is derived

from the relation between the depth of damage d and removal rate Z, written in the form d = CX Zm . The boundary conditions are the shape error

in the workpiece d

Q and the final, acceptable, depth of damage d

g

. The

solution shows that the minimum total time is achieved when the number

of stages increases; more than three or five stages are ineffective

both as regards the constant duration of each stage and the optimum duration. In contradistinction to normal practice, the first stage - to be

optimum- should ensure a high removal rate over a short period of time.

The following rates are one order of magnitude smaller and last longer.

A map with time contours is given for three-stage grinding; a shallow

minimum is found which allows wide margins for choosing removal rates.

Key Words: Grinding; grinding damage; lapping; polishing; polishing times.

1 . Introduction

Grinding and lapping are used to shape a given workpiece correctly and obtain the

required surface quality. Shaping is the main aim of a grinding operation, but it leaves

a damaged surface. This is then lapped and polished until the degree of damage is acceptable. These stages are interdependent: the higher the removal rate, the greater the damage

and the longer the final operation. Where economy is important, the rough grinding rate

must be balanced against the time spent in removing the damage by a finer treatment.

Various examples of such multistage operations have been described in the literature,

e.g. for glass [ 1,2,3] \ where two to four stages are used. In preparing samples for

microscopic observation five or more stages may be necessary. In such laboratory work

the duration of the various stages need not be controlled, in contrast to a production

process where the time for each stage is often equal [2,3] . In view of these variations

in grinding and lapping procedures, it is worthwhile to know how many stages give an optimum result and how the subsequent times and removal rates should be chosen.

A useful approach has been described by Mushardt [ h] and Scherf [ 5] who combine on-line

measurements of the volume rate of removal and the surface roughness with numerical

control of the grinder. They include wheel wear and sparking out in their analysis and

obtain workpieces with the roughness required. Roughness, however, is not always a good measure of the damage in the workpiece. In ground metals there is a zone of con- siderable depth where the material has been modified by plastic deformation and the

heat generated [6] ; also, in glass and ceramics, deep cracks may be masked by a smeared

surface [ 7] . For a high quality surface such layers must be removed.

The present treatment optimizes the total time of the multistage operation by

considering for each stage the time required to remove the depth of the damaged layer

created by the previous stage. This approach leads to a set of equations that can be

solved when the initial shape error and the final, acceptable, depth of damage are

given. To achieve this, the relation between the rate of removal Z (mm3/s) and the

depth of damage d must be known for each stage. It is possible to start from Preston's

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

191
equation [8,9,10] between Z and force F,

Z = F v/H (1)

s g

where v is the wheel speed and H the grinding hardness . Because of the force a damaged

layer is present. A quantitative delation between its thickness d and F can be given in

a few cases

:

1. d«F2/H2 when groove formation [ 10] occurs as in low-speed grinding in metals and in

fine grinding of glass and ceramics; H is then proportional to the Vickers hardness.

2. d<*F2/3/Kw3 when craking and chipping occur as in rough grinding of brittle materials

[ 11,12] ; K is then proportional to the fracture toughness Kj c . Generalising from these

two cases it seems appropriate to assume a power law between d and F: d « Fm . For practical purposes it is more useful to relate the damage to the rate of removal by writing

d = a Zm

(2)

using one value of m for a range of finishing processes. This has the added advantage over using F that eq. (2) includes deviations from (1) - the "size effect" which is so often found in grinding metals [ 13,1^+1 and ceramics [ 15,16] . In this form the treatment

becomes fairly general and may be applied to a wide range of finishing operations.

An example of a relation such as eq.(2) is shown in figure 1 for glass grinding with

three rates of removal . The damage was made visible by sawing through the grinding

tracks and polishing away the damage due to the sawing. The length of the biggest cracks

are given in the figure.

100, .

Figure 1 . Depth of damage in glass after

grinding on various machines with

different rates of removal Z. The line drawn has a slope of 0 .

5

Cl

II

11 1 1 1 I

I I

10 100 1000

rate of removal Z (mmVs)

In the next section the optimization procedure for a two-stage situation is considered

which is extended to k stages in section 3. The applicability of the model is discussed

in section k.

2. Two-stage optimization

Consider the example of a horizontal axis surface grinder which is used for two

stages of finishing. The shape errors of thickness d

Q must be removed; the final surface

damage must be smaller than a given thickness d

g

. In the first stage no more than d

Q is

removed, in the second stage no more than the damage that results from the first stage,

d-| . In total a thickness dQ + d-| is removed, which should be acceptable in view of the

tolerance on the final dimensions. The depth of cut a and the workpiece speed vw must

be chosen for each of the two stages. For a sample of width b the volume rate of removal

Z equals

Z = a.b.v^ (3)

192
For a sample length 1 the volume to be removed in the first stage is d .l.b; the grinding

time is then

T = d bl/Z. (1+)

1 o 1

where the index 1 denotes the first stage. For a given Z^ the damaged layer is known

from eq.(2), d^ = aZ^ m

; this layer is removed in the second and last stage by using

a removal rate Z2. This leaves damage over a thickness do = aZ^1., which should be the

known required surface damage d

g . Z

2

is therefore known from eq.(2),

Z

2 = (de /a)

1/m

(5)

The grinding time for the second stage is T^ = d^bl/Z^ and the total time becomes

d d

T = T

1 + T2 = {-f + ^).bl (6)

Here d is given, Z^ follows from eq. (5) and d^ is related to Z-| by eq.(2). Two choices

for Z-| will be considered, the equal-time case and the minimum-time case. In the first

one, the times for the two stages are taken equal, T-| = T2 and thus d

Q /Z^ = d^/Zg, thus

d

Q

Z

2 = d^l5 or using eq.(2),

%h-<" (t)

from which Z can be solved. The total time becomes

1

T = 2T

1 = 2 (ad

om

/Z

2

) m+1

.bl (8)

In the second case the duration of the two stages is not taken as equal, but Z^ is varied

to minimize T. In eq.(6) a minimum occurs when

d

Q

Z

2 = maZ^ +m

(9)

which gives a higher Z^ value than eq.(T). From eq.(9) follows that ^/Z^ = nHotZ^/Z^ = = mCd.j/Z ), or using eq.(V

T

1 = mT2 (10)

The first stage takes a shorter time than the second, it is also shorter than in the

equal-time case. The total time is now 5.5% lower than in eq.(8). A further reduction

in total time is possible by using more stages.

3. k stages

The extension from 2 to k stages does not essentially change the minimization problem.

As before, the initial shape error dQ and the final depth of damage d

g are given. The rate

of removal during the last stage is then found from eq.(2)

Zk

=

( d

e /ct)1/m

The two cases considered before - equal times for all stages and minimum total time - can

be calculated now (see the Appendix for details). For equal times the total time T is

given in figure 2a as a function of the number of stages used, for various values of m and

a fixed value of (d /d ) From the minimum in the T(k) curve it follows that, for

. o e * . . equal times, the number of stages should not be increased above a certain k that still

depends on the material behaviour, i.e. m. 193
50

20

10

d0/de = 316

tj varied

\ \

m =2/3

r

V2.

V,

i i i 1 I I I 2 3 4 5 6 7 8 9 10 — number of stages k

(a)

2 3 U 5

(b)

6 7 8 9 10

Figure 2. Total time needed for grinding in k stages as a function

of k. On the left the result for the case when the duration

of the stages is equal; on the right the times are not equal

and the optimum total time is shown, m is exponent in eq.(2),

d is the initial shape error, d is the final depth of damage.

Tne unit of time is given in the Appendix.

For the case of minimum total time the result is given in figure 2b. Again a minimum time

occurs, now at higher k than in the first case and the minimum is very shallow.

For five stages of grinding the results have been worked out in greater detail. The

removal rates in the subsequent stages are shown in figure 3 for m = ^ and (d /d ) 2 = 10 5. The constant-times case has lower Z values, but takes longer than the optimum wfiere the

times were varied. Here higher Z values are found, especially in the first stages; these

higher removal rates, however, act during a shorter time!

Figure 3. The removal rates during subsequent grinding stages as a function of

time for five stages . Equal duration of

each stage is shown by the broken line,

optimum: full lines.

194
For three stages, e.g. grinding, lapping and polishing, a simple representation is

possible. The total time is given by

T = {d

Q

/Z

1 + d^Zg + dg/Z } hi (12)

As before, the starting shape error d and the final depth of damage d are given. With

eq.(ll) the final removal rate Z-^ is ?hen known. This makes T a function of Z^/Z^ and Z0/Z3

only. For every possible operation where these ratio's are known, it is possible to indicate the T value on a graph where Z -j /Z^ and Zg/Z^ are plotted along the axes (fig.U).

From the model presented here T values have been calculated and plotted as contours of

constant time in the figure; m = \ and d

Q /de

= 3 1 6 were used for this case. The two standard

cases are indicated too: for constant duration of the three stages the time (in certain

units) is 6.9*+; if the times are allowed to vary, the total time has a minimum of 5.92.

The margin around the minimum is seen to be quite large (notice the logarithmic scales

for bbth axes). Near the T.| = Tg = T^ point the hatched area in the figure corresponds

to a variation in T

1 /T

2

and Tg/T^ by + 20%.

The influence of the choice of dQ /d on the optimum time T is not very strong:

for m = g, T is proportional to (

Q /d

e

)

1 / 12 , for m = I the exponent is 1/8U. The dashed

line in figure k indicates the locus of T when d

Q /de is varied. There is hardly any

effect on the Zg/Zo value, but rather more on the abscissa Z^/Z^. For lower m values the

minimum of T as a function of Z^/Z^ and Zg/Z-^ becomes even more shallow than the one shown in figure h.

Figure h. For a three-stage operation the total time is shown as contours in a graph

where the removal rates Zg and Z^ (normalized on the last rate Z3) are plotted along the

axes. The round dot is the optimum choice, the square corresponds to constant duration

of the stages (the hatched area refers to 20% variations in t-| and t

p

). The lines are calculated for m = \ and d

0 /de< For other dQ /de values the optimum shifts along the broken

line. The crosses are examples of a practical operation taken from [ 1] .

195
k. Discussion

The requirements for the application of the model described above are twofold:

firstly, the depth of damage after an operation with a known volume rate of removal

must be known; this d(Z) relation supplies the values for a and m in eq.(2). Techniques

for measuring d are available [6,7] • Secondly, the starting conditions (dQ

) and the acceptable damage (depth d

g

) must be known. The recipe for all stages follows from dQ /de = fDm

and the equations given. It should be pointed out that the d(Z) data are to be taken

on the grinding machine as it operates, e.g. at constant wheel speed and with the same abrasive, at the concentration and grain size to be used in practice and at the same wheel width. The results of the calculation show a minimum time for the total operation,

but the minimum is fairly shallow (fig. k) . With an increasing number of stages the rate

of removal in the initial stages rises considerably, although the duration of these

stages is short (fig. 3). Apparently the model favours boldness, in contrast to the

approach used in current practice. For example, in Balai 1 s paper [ 1] a three-stage

grinder (t. is constant) is described with consecutive depths of cut of 0.7-2.5 mm,

0.2 - 0.7 mm and 0.02 - 0.03 mm. Assuming constant workpiece feed these data are included

in figure h. is about three times Z

2 whereas figure h shows a Z

1 value that is about 25

times higher for constant duration of the stages. The total time is not far from the

optimum, for this value of d

0 /de (where no damage data were available, m and dQ /de need

not be correct for this example). No notice was taken of the time required for handling

the workpieces between the stages ; this is expected to lead to a sharper minimum in

figure 1. The minimization procedure for the total grinding time that was reported here,

is of economic use when the grinding costs are taken proportional to grinding time.

Other contributions to the total costs, such as wheel wear and the resultant cost of

dressing, were disregarded, but can be incorporated where quantitative data are available,

e.g. in the grinding of metals [ U-6] . High Z-| values (and thus high wheel wear) are expected to be penalized and the optimum in figure k will shift to the left.

Sparking out as a last stage in grinding improves the surface quality; the relieve

of the elastic deformation of the machine is easier to measure than predict [ ^,5] • Might it perhaps be possible to control sparking out and thereby eliminate many of the k

stages? The geometric progression in t^ suggests a continuous grinding operation where

Z decreases with time as a power law, Z <* t -n (n is about 2 for the data in figure 3).

Numerical control of feed or depth of cut would be necessary.

5. Conclusions

-A model has been given for the optimization of a multistage grinding operation.

The duration of each stage need not be larger than the time required for removing the

damage caused by the previous stage. In combination with the shape errors that must be

removed and the acceptable damage at the end of the operation, a complete method has been

worked out for the choice of removal rates in the various stages.

- The basis of the calculation is the relation between the depth of damage in a given

operation and its volume rate of removal. An example was given for grinding glass.

- The total finishing time shows a minimum for the proper choice of the consecutive

removal rates and the duration of the stages. The minimum time improves as the number

of stages increases (fig. 2). Depending on the behaviour of the workpiece, three to five

stages are sufficient. If the various stages are of constant duration the total process

takes longer. Variation of the removal rates around the optimum values shows that the

curve representing the total time has a rather shallow minimum (fig. k) . - The removal rate during the initial stages increases sharply with the number of stages,

but the duration decreases (fig. 3). In practice, however, the ratio of consecutive

removal rates is much smaller than predicted here. Improvement seems possible,

- For a three-stage operation a convenient method of presentation is to plot the reduced

removal rate during the first stage versus that of the second, both divided by that of

the third stage (fig. h)

.

- On the assumption that the depth of damage has been determined for the range of removal

rates that is of interest , the calculations may be applied to any multistage grinding

operation, be it on metals, glass or ceramics. By using data on wear, dressing and handling, the model may be extended to cover those situations as well.

196
References

[1] Balai, V.V., Method, technology and automated equipment for diamond polishing of

optical prisms, Sov. J. Opt. Techn. U3, [8] , U81-U83 (Aug. 1976).

[2] Ardamatskii, A.L. and Vorontsova, N.M. , The polishing of optical components vith a diamond tool, Sov. J. Opt. Techn. U3, [ 12] , 756 (Dec. 1976).

[3] Bohache, J.J., Chen, G.H. and Moore, D.T., Design of a high speed optical production

line, Opt. Eng. JJ5, k 16-1+23 (Sept. 1976).

[ k] Muschardt, H., Optimierung des mehrstufigen Schleifprozesses , Techn. Mitteil. 69 , h 18-U20 (July 1976).

[5] Scherf, E. , ACO-Pilotanlage fiir den mehrstufigen Schleifprozess , Techn. Mitteil. 69,

1+21-1+2U (July 1976).

[6] Peters, J., Snoeys, R. and Maris, M. , Oberflachenspannungen beim Einstechschleifen,

Techn. Mitteil., 69, 338-3^2 (July 1976).

[7] Rupp, W. , Surface Structure of fine ground surface, Opt. Eng. 15 [ 5] 392-397 (May 1976).

[8] Preston, W.J., The Theory and design of plate glass polishing machines, J. Soc. Glass

Techn., JM> 21^-256(1927).

[9] Kumanin, K.G. , Generation of optical surface (The Focal Library, London, 1962)

[ 10] Rabinowicz, E., in Friction and wear of materials , p. 168 (J. Wiley, New York, N.Y. 1966)

[ 1 1] Lawn B.R., Fuller E.R. and Wiederhorn, S .M. , Strength degradation of brittle surfaces:

sharp indenters, J.Am.Cer .Soc . , 59 [ 3] 193-197 (May 1976).

[ 12]Veldkamp J.D.B., Crack formation during scratching, (Proc. 9th Int. Conf. on Science

of Ceramics, Noordwijkerhout , November 13-16, 1977), Book, Science of Ceramics , 9., K.J. de Vries ed. , pp. U38-UU6 (Nederlandse Keramische Vereniging, 1977)

t 13]Bacher, W.R., Marshall, E.R. and Shaw, M.C., The size effect in metal cutting, Trans.

ASME, Jk, 61-72 (Jan. 1952).

[ 1U] Shaw, M.C., Fundamentals of grinding, Book, New Developments in Grinding , M.C. Shaw,

ed., pp. 220-258 (Carnegie Press, Pittsburgh, Pa., 1972).

[ 15]Broese van Groenou, A., Veldkamp, J.D.B., and Snip, D., Scratching and grinding

parameters of various ferrites, J. de Physique, CJ_, 285-289 (1977).

[ l6]Broese van Groenou, A., Grinding experiments on ceramics, see ref. 12, pp. U86-1+93.

Appendix

k stages

As in the two-stage problem in section 2 the initial shape error d

Q

, the final depth

of damage d and, from eq.(2), the rate of removal during the last(kth) stage,

are given. Using these a number of useful dimensionless quantities can be introduced. The removal rate during stage i (i = 1 to k) can be normalized by

introducing reduced removal rates fj_ = Z^/Z^. The rate f^ is a decreasing function of i

and satisfies a number of equations which wxll be derived below. In the last stage,

for i = k, fk

= 1. It is also useful for the following calculation to introduce f0 for

k = 0; f

Q

is related to d

Q by dQ = aZ^fJ1 ; also

f* = d /d (A1)

o o e

197
The duration of stage i is determined by the damage in stage i-1 , d.^_ ^

, and Z. . By analogy with eq.(U) can he written

T. = d. 1 hl/Z. = aZTM bl/Z. = (etZ

m_1 hi) {fTM Jf.) (A2)

1 1-1 1 1-1 l e 1-1 l

The first factor in the last expression is convenient for introducing a reduced time,

t^; consequently, by definition,

The total reduced time g for k stages is given by

k k

g(k) = E t. = t f? Jf. {Ah)

1=1 1-1

As before, two approaches are possible: 1. t^ is constant and 2. t^ is allowed to vary.

t- is constant

i

Each grinding stage has now the same duration, which we will call t, so t. = t,

The rate f. (i = 1 to k) can be derived from eq.(A3) taking the logarithm,

In t. = In t = m In f. , - In f. (A5)

l 1-1 l

or

In f. = m In f. , - In t (A6)

i 1-1

After elimination of In t this becomes a set of second order difference equations in In f

-

with boundary conditions for i = 0 (f is known) and i = k(f^ = 1). The solution is

In f. = In f {S(k) - S(i)} / {S(k) - 1} (AT)

1 o

In t = m

k

In f /S(k-1 ) (A8)

where

S(k) = 1 + m + + m

k

= (l-m

k+1

)/(l-m) (A9)

The total time given by eq.(A^) is now knownj g(k) = kt . Figure 2 gives g as a function

of k for various values of m and for f m = 316. The value of f. can be calculated from

eqs.(A6) and (A8).

t. varies

1

When f. is varied in eq.(Ait) to minimize the total time g(k), two terms contribute,

namely t^ and t. + ^; after differentiating to f| and putting the result equal to zero,

we find

m f

m+1

= f? . f. , (A10)

1 i-1 i+1

198
t.can now be expressed in t... by

1 1+1

t. = fm Jf. = m f^/f-^ = m t. , 1 i-1 i ii+1 i+1

as in eq.(io). This can be continued for i+1, i+2 to k-1: t^

Hence

m t. +1

(A11)

k-i.

= m t.

;(k) t. = a.k

, m

k~ 1 = t, S(k-1)

i=1 l k i=1 k

(A12)

Also from eq.(A10), by taking the logarithm,

In m +m In (f^f^) = In (f^/f^ (A13)

This is again a second order difference equation in f . , with the same boundary conditions

as before. The solution is similar to eq. (A7) , with an extra term containing In m,

ln

i Is(k)-1 * 1

In m

m-1

S(k)-S(i)

S(k)-1

ln f (A1U)

t

k

follows from eq. (A3), t

fc = f£ ^

f^_

1 follows from eq.(AlU) and g(k) from eq.(A12)

Figure 2b gives g(k) for f

Q

m

= 3 1 6 and for various values of m.

List of symbols

a depth of cut *i normalized duration of stage i

b width of workpiece V

s vw

wheel speed

d shape error before operation workpiece speed

4 depth of residual damage after operation F grinding force

d. depth of damage after sta ge i grinding hardness

4

= Z^/Ze , reduced removal rate Vickers hardness

f

o from d

?

/de = fQm K fracture toughness

g normalized total time S sum of powers of m

i integer Ti duration of stage i k number of stages T total time

1 length of workpiece Z volume removal rate (m3/sec) m exponent in eq. 2 Ze removal rate in last stage

n exponent

a

removal rate in stage i

proportionality factor

Discussion

NARKEN

1. In the abrasion process, regarding the normal and tangential force vectors (Fn, Ft) when considering the machine stiffness, the machine will yield

in some degree to both vector components (or resultant). The yield to F^

is probably of little consequence. However, the yield to Fjj (perpendicular

to grinding surface) affects the resultant flatness and also (probably)

the median crack depth (or surface damage). If these problems are of concern then it should be considered that the abrasion process should be designed (wheel and machine parameters) to minimum Fn.

2 » Regarding the optimization of consecutive removal rates

T. = —2 + —i + . . . + n

Z

l

Z

2 V1

199
It is implied that the total depth of surface removed is Dv = d +

d + . . . d . Ideally it appears that a more efficient approach would be

clearly, the more steps involved in the process, the more significant is the

effect.

VAN GROENOU

1. A frequency dependence of the compliance is know to lead to chattering

and an uneven surface (in steel grinding the wheel is the most compliant part

of the system, metal -bonded wheels are much stiffer). A yielding machine leads

to a lower grinding force and a lower removal rate. It is therefore possible to do fine grinding by mounting the workpiece on a piece of rubber, with

the risk of chattering.

2. This seems to be a good idea,

MALKIN

You have assumed that the depth of damage depends on the removal rate Z.

This is contrary to our results which show for plunge grinding of glass that

increasing the removal rate by taking a bigger depth of cut has almost no influence on surface damage and transverse rupture strength, whereas increasing the removal rate by taking a faster workpiece velocity causes more

damage and reduces transverse rupture strength. This would completely alter

the nature of your optimization problem.

VAN GROENOU

If the depth of damage is only a function of the workpiece velocity V

and independent of the depth of cut, a, it is clearly advantageous to use

w

large a values. The total time can then be minimized with respect to the

various V^ values, that can be chosen for the consecutive grinding stages.

Apart from factors that contain the ratio's of the subsequent depths of

cut the results will be similar as those given in the paper.

RICE

On your log - log plot of depth of damage (d) vs. removal rate (Z)

you had three points through which you drew a straight line. However , these

points would also suggest an initially slower increase of depth of damage

(d) with removal rate (Z) with d becoming progressively greater as Z increases in contrast to your linear log - log plot. Have you other information to clarify the possible distinction between these two trends? Our

data suggests the second of the above trends i.e. even more non-linear than

a straigth line on a log - log plot.

VAN GROENOU

I have no other information. The power law, d & Z

m

, was chosen because

of ease of computation. I think the extension to other relations should not

be difficult.

200
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

GRINDING STONE, CONCRETE AND CERAMICS WITH COATED ABRASIVES

R. C. Lokken and E. J. Duwell

Industrial Abrasives Div., 3M Co.

St. Paul MN 55101

Silicon carbide coated abrasives grind stone or ceramics most efficiently

when run dry and at high speeds. This high energy process appears to heat the

surface layer of the workpiece to temperatures which make "chip formation" easy

However, neither the ground surface of the workpiece nor the coated abrasive is

thermally damaged to any significant degree. Dry dust collection may be an attractive alternative to the removal of wet debris in water lubricated operations .

Key Words: Abrasives; ceramics; coated abrasives; grinding; stone.

1. Introduction

Ceramics are generally regarded as "hard-to-grind" materials. This may be caused by

their hardness, or perhaps by their inability to deform plastically to form continuous

chips as metals do. Chemical interactions may also affect the grinding process. Quartz,

for example, is undoubtedly difficult to grind because of its hardness (KHN = 820 kg/mm?),

which is twice that of a fully hardened high carbon steel. However, it may also react

with abrasives like aluminum oxide to form alumina-silicates at the grinding interface.

Thus, hard oxides are usually abraded with hard carbides such as silicon carbide or diamond. Water floods are generally used to minimize thermal or chemical damage to

the abrasive in the grind interface.

This project was started with the objective of using wide abrasive belts to machine

or finish large concrete or ceramic surfaces. Although wet grinding does not adversely

affect the mechanics of using an abrasive belt, the problems associated with the control

and removal of wet debris were a strong incentive to include an investigation of dry

grinding conditions. Granite was chosen as a workpiece, mainly because it is readily

available in our area in large pieces of uniform composition that were found to give

grinding data of excellent reproducibility. Granite is composed largely of quartz and

feldspar, but the exact composition varies considerably depending on location of the

deposit.

The method to be described in this paper involves the utilization of coated abrasive

belts under dry conditions and at high belt speeds and high horsepower levels. Grinding

debris is collected in a high velocity air stream and deposited in dry dust collection

equipment. Although the early development work was done on granite, marketing efforts

have led to use of the process in the concrete block industry for dimensioning and facegrinding block [1]. Grinding concrete block to accurate dimensions permits the use of

thin-line adhesives and rapid lay-up procedures, while face-grinding imparts an attractive

terrazzo-like appearance to the surfaces of the concrete block [1].

To date this process has not been applied to the machining of any other ceramic or stone composites. The purpose of this paper is to offer additional data on the performance of coated abrasives for dry grinding of other ceramics, and to relate, as far as possible, the grinding process to the properties of the workpiece being ground,

backing before the mineral is electrostatically drawn up to the backing. This procedure

serves to orient the mineral grain so that elongated particles have their long axis

perpendicular to the backing. A second adhesive layer, or size coat, is deposited over

201
2. Structure and Evaluation of Coated Abrasives

A coated abrasive consists of essentially one layer of abrasive grain on a flexible

backing (Figure 1). A first adhesive layer, called the make coat, is deposited on the

the mineral to provide further anchorage to the abrasive grain. The coated abrasive

material is converted to belts for use in this process.

Both the cost and stock removal capability of an abrasive belt tend to be closely

related to the surface area of the belt. A belt 160 inches long will in general remove twice as much stock as a belt 80 inches long, while doubling the width will result in

another two fold increase if the workpiece is correspondingly wider. It is, therefore,

convenient to define two terms which can be combined to give a measure of belt efficiency:

Cut: The total volume of workpiece material removed.

Path: The area of coated abrasive actually used to remove stock.

3 2

Efficiency: Cut (in )/Path (in ) when the abrasive belt is no longer

capable of removing stock at a satisfactory rate or when the

horsepower requirement exceeds the capability of the machine.

A simple schematic of a belt grinding operation is shown in Figure 2. Because the

belt is generally wider than the workpiece, some attention must be given to distribution

of wear across the belt. Uneven belt wear detracts from belt efficiency and also limits

the dimensional tolerance that can be produced in the workpiece. If the belt is substantially wider than the workpiece and the workpiece is not uni-dimenstional in the

grinding direction, uniform belt wear can be achieved by staggering the entry point of

the workpiece, or by oscillating the belt [2]. Belt oscillation can be controlled through

deflection of the idler drum (fig. 2).

BACKING

Figure 1. Schematic diagram of the cross section of a coated abrasive.

Figure 2. Simple schematic of a belt grinding

operation. Belt can be oscillated in" directions

a and b by appropriate deflection of idler.

Spindle is generally driven, and for dimensioning

is at some fixed distance (d) from conveyor bed

(III). Load on bed (F n ) and cutting force (F c ) increase as belt wears. Specific energy of

grinding (E s = hp-min/in-3

) is independent of F

n but is proportional to Fc

.

202
Power requirements depend on depth of cut, belt speed, workpiece infeed velocity,

state of wear of the abrasive belt, and on the "grindabil ity" of the workpiece. 1 It is

usually expressed as horsepower per inch of workpiece width (HP/in) and does not include

the no-load power requirement of the machine. Rate of cut, R, is also normalized with

respect to the width of belt actually used (in 3/min/in) . Belt efficiency (cut/path) is

a good indication of workpiece grindabil ity, as is the energy (E

s ) required to grind up

a unit volume of workpiece (HP-min/in 3

) . E,. is obtained from the relationship E, = P(HP/in)R(in3/min/in). s H s

3. Grinding Granite

Using Dakota granite, a wide range of machine conditions were used with coated

abrasives made with silicon carbide (SiC), boron carbide (B^C) and aluminum oxide

(A^Og). Grade 24 minerals were used since the objective was to remove stock rapidly

as a result of large downfeeds, high belt speeds, or large throughfeeds . Both dry and

water-flooded experiments were performed. The investigation showed that by far the

greatest belt efficiencies were obtained with the following conditions:

(1) Silicon carbide mineral

(2) Dry grinding

(3) A hard non-elastic backing to carry the mineral

(4) A smooth steel contact wheel

(5) High belt speeds (>9000 ft/min), (>2743 m/min)

(6) High throughfeeds (300-100 ft/min), (9-30 m/min).

Decrease in belt caliper, belt wear, and power requirement are plotted as a function

of stock removed in figure 3. The power consumption is noted to level off and the rate of

40i 1

J»30

1

C

Zzo

4>

10-

a

Figure 3. Belt wear and horsepower requirements for grinding Dakota granite as a function of stock removal. A 24-mesh silicon

carbide belt was run at 7800 ft/min over a smooth steel contact wheel. Infeed was 10

ft/min, downfeed was 0.01 in/pass.

_i i i i i i 0.1 02 0.3 0.4 _ 0.8 0.6 cut/path (ii^/in2

)

^ince the machine settings and abrasive tool dimensions are in English units, these units

will be used in this paper. For reference, a belt speed of 10,000 ft/min = 305 m/min; a rate of cut of 1.0 in 3 /min/in = 6.45 cm3 /min/cm; a cut/path of 1.0 in 3/in 2 = 2.54 cm3 /cm2

; 1.0 HP-min/in 3 = .653 Kcal/cm3

.

203
belt wear becomes linear after the initial sharp rise in both variables. This initial

sharp increase is believed to accompany the fracture and loss of a small population of

sharp grains that rise above the general level of the abrasive coating. Under the same conditions, the introduction of water results in an increase of 50% in power consumption

and a substantial loss in belt efficiency. Use of an aluminum oxide belt of the same grade resulted in a decrease in belt efficiency of approximately 95%. B»C performed

even more poorly, failure appearing to be a result of catastrophic mineral fracture.

Although belt speeds of 3000-6000 ft/mi n commonly give optimum belt performance

on metals, higher speeds are definitely better for grinding granite (fig. 4). Grinding

0.5

CMc 0.4

c 0.3

JCo

a. 0.2

O

O.I

I I I 1 1 1 1 1 ,

I 2 3 4 5 6 7 8 9x10

s ABRASIVE SPEED (ft/min)

Figure 4. Cut/path as a function of belt speed for granite grinding with 24-mesh

silicon carbide belts. All data was taken at a belt speed of 7800 ft/min over a smooth steel contact wheel. Infeed speed was 10 ft/min, downfeed was 0.010 in/pass.

granite dry at 9000-12000 ft/min produces highly incandescent grinding debris, while the

belt and the ground surface remain relatively cool (the surface is slightly warm to the

touch). Thus, it appears that grinding energy leaves the grind interface almost entirely

in the chips being formed. The situation appears to be improved by increased throughfeeds, and throughfeeds of 50-100 ft/min are now commonly used for grinding concrete block

(there are no commercial applications of this method in the stone industry). Since these

materials do not increase in temperature as a result of air oxidation, as is the case with

metals, it must be concluded that the conversion of mechanical energy to heat energy in

the grind interface increases the temperature of the surface layer being removed to

possibly 1000°C or higher.

Not surprisingly, examination of the ground surface and the swarf (grinding debris)

indicates that some of the particulates are globular and apparently became molten during

the process. However, angular particles are also present, indicating chipping or spalling.

From the data in figure 3, it can be calculated that

E

$ = 3.1 (HP/in)/l .76 (in

3 /min/in) = 1.76 HP-min/in 3

(= 1.15 Kcal/cm3

).

Based on the constituents of granite, a specific heat (C ) of 0.20 can be reasonably

assumed. Thus, the heat generated can be expected to increase the temperature (T) as follows:

T = ^ x 1000 = I'll

*

= 2090°C. (p = density)

Although some heat is certainly lost in the abrasive belt, to the atmosphere, and into

the remaining surface, it is not surprising that the swarf temperature is increased to

the melting point of the constituents in the stone.

204
4. Grinding Other Ceramic Materials

A porous alumina fire brick and a dense magnesite fire brick were abraded under the

high speed, dry conditions, used for granite. The porous alumina brick offered such

little resistance to stock removal that power measurements could not be determined with

any degree of accuracy. The magnesite brick offered much less resistance to abrasion

than granite, and was effectively ground under the conditions recorded in table I. Belt

efficiencies were not determined for all conditions, but based on data generated for

granite for which the specific energy of grinding is greater (M.75 HP-min/in-^), belt

efficiencies (cut/path) probably exceed 3 in^/in^. The magnesite bricks displayed distinct

grind lines on the abraded surface, indicating the possible existence of a chip forming

process by plastic deformation.

Table I - Grinding Data for Magnesite Fire Brick

Belt Speed: 8650 ft/mi n, Infeed: 10 ft/mi

Depth of Cut Rate of Cut ^s

(mils) (in /min/in) (HP-min/in ) 10 1.2 0.358

20 2.4 0.333

30 3.6 0.327

40 4.8 0.329

50 6.0 0.331

Although the aluminum oxide firebricks were easily abraded, dense forms of this com- pound are not easily ground, as evidenced by the data in table II. Dense, fused aluminum

Table II - Grinding Data for Fused and Natural Corundum (A1 ? 0.J

Belt Speed: 8650 ft/mi n

c 6

Workpiece Depth of Cut Rate of Cut

E

s

(mils) (in /min/in) (HP-min/in

3

) Fused A1

2

0

3

* 10 1.26 2.1

15 1.89 1.3

20 2.52 0.91 ** Corundum 10 0.6 1.35

ii 20 1.2 1.11

ii 30 1.8 1.00

ii 40 2.4 0.93

ii 50 3.0 0.91

ii 60 3.6 1.10

* at 10.5 ft/mi n infeed rate

* Corundum in terrazzo composite - at 5.0 ft/mi n infeed

oxide for the preparation of abrasive grain was somewhat more difficult to grind than

natural corundum in a terrazzo matrix. In both bases, examination of the swarf or ground

surface failed to reveal evidence of melting, but rather indicated a material removal

process by chipping and spall ing.

More recently, alumina has been co-fused with zirconia at or near the eutectic com- position to form mechanically tough, thermally stable structures. Grinding of this

material produced specific energy values far in excess of that noted for any other fused

ceramic (table III). As in the case of fused A1 9 0,, higher cut rates appear to result

205
Table III - Grinding Data for Fused Eutectic Alumina-Zirconia

Belt Speed: 8650 ft/min, Infeed: 10.5 ft/min

Depth of Cut Rate of Cut

E

s 3 3

(mils) (in /min/in) (HP-min/in )

2 0.25 10.20

3 0.38 6.64

5 0.63 3.69

in lower specific energy of grinding. For this material, chip formation appeared

definitely to involve plastic deformation or melting. The ground surface was metallic

and shiny in appearance with well defined grind lines. The chips exhibited the curl

shapes noted in metal swarf.

Quartz was ground both as a synthetic single crystal and in its natural polycrystalline form (flint). It was apparent that the grindability of the single crystal

depended on crystal orientation (table IV). From the limited data obtained, the flint

Workpiece

Crystal

Crystal

Flint

**

***

Table IV - Grinding Data for Quartz

Belt Speed: 8650 ft/min

Rate of Cut

(in /min/in)

2.52

2.52

2.52

grinding parallel to a axis

grinding parallel to c axis

(HP-mi n/in

-3

) 0.23

0.42

0.71

was more difficult to grind than single crystal quartz, although it is possible that some

crystal orientations may result in far greater resistance to chip formation. No evidence

was found of plastic chip formation, with all workpieces exhibiting extensive chipping

and thermal spall ing.

5. Discussion of Results

High belt speeds, high infeeds, and dry grinding conditions appear to create a material removal process in which the energy of grinding is removed from the grind

interface almost entirely by the chips. Composite workpieces such as granite, fused

alumina-zirconia, or concrete, appear to resist spalling and fracturing, with the result

that the chips are formed plastically or by actual melting of the surface layer being

ground away. A survey of specific heat values [3] for the materials under consideration

indicates that between 100°C and 1500°C, specific heat values generally rise from 0.2

cal/g/°C to 0.3 cal/g/°C. (Zr02 has a specific heat range of 0.11 to 0.18, but its

density is also greater than the rest of the material studied). The greatest values for

E s were noted for the alumina-zirconia eutectic. This ceramic has a density (p) of 4.7.

For an E s of 5.0, and assuming an average specific heat of 0.15 for alumina-zirconia,

the increase in temperature is calculated to be at least 1800°C (to melting point) with

about 2000 cal/g left over for fusion (assuming, of course, that all energy goes into the

chips). Thus, as in the case of granite, the swarf is noted to be incandescent as it

leaves the grind interface. By contrast, AT for magnesite is calculated to be 288°C and

no incandescent chips are observed during grinding. Various architectural marbles yield

similar values for E , and also did not exhibit incandescent grinding debris.

206
It should be noted that E

s generally tends to decrease with an increase in depth of

cut. This is probably caused by material removal as a result of sub-surfacing cracking,

and less chip formation by deformation processes. It should be noted that the energy

losses in the machine at higher rates of cut will increase, this tending to be a factor

contributing an error which would result in lower values for E than are calculated

assuming the no-load horsepower reading. s

Whether chip formation is induced by fracture or by plastic processes, it is apparent

that high stock removal rates by grinding can be achieved at high belt speeds and. under

dry conditions with silicon carbide coated abrasives. Commercially, concrete block are being ground with belt efficiencies (cut/path) of 1.0 - 5.0 in /in . At current coated

abrasive prices, this represents a tool cost of generally less than $0.01 for each cubic

inch of workpiece removed.

References

[1] Lokken, R. C, McDonald, W. J., and Duwell, E. J., Grinding Stone and Concrete with

Coated Abrasives, Am. Cer. Soc. Bull . , 5_3, [a], 646-649 (1974).

[2] McGibbon, T. R., Amundson, S. E., and Lokken, R. C, CAMOS: A Computer Aided System

for Establishing Optimal Parameters for Achieving Uniform Abrasive Wear During High

Rate Grinding, Wear , 37, [1], pp. 1-14 (1976).

[3] International Critical Tables, Vol. II, p. 85.

207

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

TURBINE COMPONENT MACHINING DEVELOPMENT

D. W. Richerson and M. W. Robare

AiResearch Manufacturing Company of Arizona

Phoenix, Arizona 85010

Various machining methods that have been evaluated to produce ceramic turbine

components from high-strength ceramic materials are discussed. These machining methods

are compared in terms of economics, dimensional tolerances, surface finish, and strength of

the resulting product. Special emphasis is on the machining methods developed for

fabrication of rotor blades from hot-pressed silicon nitride (Si^N^).

Key Words: As-machined material strength; dimensional tolerances; profile grinding; silicon

nitride; ultrasonic machining.

1. Introduction

The majority of chapters in this Proceedings discuss the basic theory of machining with respect to

interactions of a tool and a ceramic workpiece. This chapter will discuss the application of basic

machining theory to fabrication of complex gas turbine components and the effects of the machining on

the strength of the material.

The examples and data that will be described were obtained during fabrication development of

turbine components conducted under the DARPA/NAVY/AiResearch Ceramic Engine Demonstration

Program. Three types of configurations were involved:

(1) "Green machining" of silicon vanes, shrouds, and other static hardware.

(2) Finish machining of reaction-bonded Si^N^ stator vanes.

(3) Individual rotor blades of hot-pressed silicon nitride (Si^N^) requiring machining of an airfoil, a platform, and a dovetail attachment.

Items 1 and 2 will be discussed briefly, item 3 in detail.

2. Green Machining

Green machining refers to machining of a ceramic part prior to final densification while the

material consists of compacted, loosely-bonded powder. Such material is much softer than the ceramic

in its final densified condition and can be machined much more economically since diamond tooling is not

required. However, the material is relatively fragile, and great care is necessary in the design and

fabrication of the tooling and fixturing so that the parts can be accurately and uniformly held during the

various shaping operations. In addition, the machining parameters must be carefully controlled to avoid

over-stressing the fragile material and producing chips, cracks, breakage or poor surface.

Reaction-bonded silicon nitride (RBSN) is presently an important material being evaluated for

static hardware (stator vanes, shrouds, transition liners, pilot cylinders, combustors, etc.) in gas turbine

engines. It is prepared by reacting a compact of silicon particles in a nitrogen atmosphere at high

temperature under controlled conditions to achieve densification. This material is relatively unique in

that no significant dimensional change occurs during this nitriding operation. Thus, to achieve minimum

cost, it is advisable to conduct all shaping operations on the silicon compact prior to nitriding.

209
Prior to green machining, the silicon-compact is either partially nitrided or presintered to provide

adequate strength for handling and fixturing. This is a critical step, because insufficient prenitriding or i presintering will leave the part too fragile and an excess will increase the difficulty of machining.

Holding of the presintered silicon for machining is typically accomplished with a combination of

bee's wax and precision metal fixtures. The part must be held rigidly, but with no distortion or stress

concentration. Metal fixturing should normally be located either on the outer diameter or the face of j the part. This is to avoid expansion of the metal fixture into the ceramic causing breakage when

heat is applied to soften the wax during removal of the machined part from the fixture. If there is no way to locate the part except on the inside diameter, special care must be taken to heat the ceramic ! rapidly while keeping the metal cool.

Once a ceramic part has been secured rigidly in a fixture, machining can be conducted by a variety

of methods — turning, milling, drilling, form wheel grinding, and profile grinding. Machining can be

accomplished either dry or wet. In either case, the silicon compact is abrasive and results in tool wear. As little as 0.005-inch wear land on the cutting edge will result in a buildup of the cutting pressure and

damage to the ceramic.

It is possible to machine silicon compacts with high-speed, steel cutting tools, but this is not

recommended for critical gas turbine components. The tools dull so rapidly that extreme care is necessary to avoid damage to the workpiece. Figure 1 summarizes a study of other cutting tool

materials using a 5-degree positive rake and 10-degree clearance angle. The GE compact diamond costs

about 10 times as much as the tungsten carbide but results in a significant cost savings in terms of : increased life, less time changing inserts, and reduced risk of damage to the workpiece from a dull tool.

O BORAZON

Ti-COATED CARBIDE

TUNGSTEN CARBIDE

G.E. COMPACT DIAMOND

0.030 -

0.020

0.015 -

jJ 0.010 -

0.005

TOOL MATERIAL SFM

DEPTH

OF CUT

(IN.)

TOOL LIFE

(LINEAR

INCHES)

MATERIAL

REMOVAL

(IN.3/MIN.)

TUNGSTEN CARBIDE 100 0.02. 6000 1.2

Ti-COATED CARBIDE 100 0.02 4000 1.2 BORAZON 100 0.02 2400 1.2 GE COMPACT

DIAMOND

300 0.04 1,296,000 12.0

FEED/REV. = 0.010 INCH

20 —

r

- 25 —

i

-

30 —

i

-

35 —

i

- 40 45

i 50 55 60

TIME (MINUTES)

Figure 1. Tool Wear During Machining of Pre-Sintered Silicon.

210
The above study was conducted with single-point turning on an engine lathe. Milling with a twoflute end mill at 200 surface feet per minute (SFM) with GE compact diamond inserts showed the same

life characteristics.

As a portion of the DARPA/Navy/AiResearch Program, Norton Company and Jarvis Company [1]

1

have teamed to develop the capability to produce complex stator vane configurations by green machining

isostatically pressed, presintered silicon compacts. Most of the grinding steps were conducted with

form-grinding wheels manufactured by the Norton Process E, which utilizes a very accurately formed

steel substrate coated with a single layer of carefully sized diamond particles. Formed wheels with 120-

grit diamond were used to shape a blank having appropriate shroud contours and dimensions. The airfoil was then profile ground using a 1/8-inch diameter spindle-mounted, diamond-coated wheel. A finished

Si„N. vane from Norton is shown in figure 2. No machining was required after nitriding. Typical

material strength in the as-nitrided condition was 44 KSI, compared to about 35 KSI for similar vanes made by other processes (slip casting or injection molding).

Figure 2. A Finished Si,N . Vane.

3. Machining After Nitriding

The loosely bonded silicon compact that could easily be machined with a variety of tooling is converted to strongly-bonded silicon nitride during the nitriding operation. The resulting reactionbonded silicon nitride (RBSN) is about 85 percent of theoretical density with high hardness and a strength

typically over 35 KSI. It can only be machined effectively with diamond tooling.

Cost analyses have shown that it is important to make the RBSN part as close to net shape as

possible since even small amounts of post-nitriding machining can dominate the cost of the part. When

machining is required, however, the same type of careful procedures as for the previous silicon compact

machining are necessary. Holding is accomplished with rigid metal fixtures and bee's wax. Machining is typically done with resinoid-bonded diamond wheels operating at 6000 to 7000 SFM with a water-soluble

coolant. To avoid edge-chipping or other damage to the part, downfeed is restricted to about 0.0005 to

0.001 inch per pass.

4. Rotor Blade Machining Development

Norton Company NC-132 hot-pressed silicon nitride (HPSN) was selected as the rotor blade

material to be machined from solid blocks into individual rotor blades for insertion into a metal rotor

; hub. The final engine blade design was not available at the start of the program; therefore, a

,

t preliminary attachment design (fig. 3) and an airfoil design (fig. 4) were selected so that machining

process development could be initiated. The attachment design was selected to permit subsequent spin1 Figures in brackets indicate the literature references at the end of this paper.

211
TOLERANCE BAND ON

THIS SURFACE TO BE

±0.0005 FROM BASIC

CONTOUR. BLEND

BETWEEN SURFACES

MUST BE CONTINUOUS

AND SMOOTH (TYP)

TOLERANCE BAND ON

THIS SURFACE TO BE

±0.001 FROM BASIC

CONTOUR. BLEND

BETWEEN SURFACES

TO BE SMOOTH AND

CONTINUOUS (TYP)

TOLERANCE BAND ON

THIS SURFACE TO BE

+0.0002 FROM BASIC

CONTOUR. BLEND

BETWEEN SURFACES

TO BE SMOOTH AND

CONTINUOUS (TYP)

TOLERANCE BAND ON

THIS SURFACE TO BE

±0.001 FROM BASIC

CONTOUR. BLEND

BETWEEN SURFACES

TO BE SMOOTH AND

CONTINUOUS (TYP)

Figure 3. Attachment Configuration for Machining Development.

Figure 4. Airfoil Configuration for Machining Development.

212
testing of the machined specimens to assess the as-machined component strength for comparison with

previous strength data for NC-132 generated with simple test bars. The selected airfoil configuration was highly twisted and tapered to make certain that it would be at least as difficult to machine as the

final engine configuration.

Since machining of such complex airfoil configurations with tight tolerances had not previously

been demonstrated for ceramics, a variety of approaches and sources were evaluated. Efforts conducted

at AiResearch included Pantograph profile grinding, ultrasonic machining, and lapping. Efforts

conducted by Norton Company included contour grinding, formed-wheel grinding, cam grinding, profile

grinding, and lapping. Efforts at Western Gold and Platinum Co. (WESGO) included profile grinding and

lapping. The key evaluation criteria for these machining development studies included dimensional

tolerances, surface finish, orientation of grinding damage, and as-machined strength.

Since the importance of orientation of grinding damage may not be readily obvious to the reader, it will be explained with respect to a turbine-rotor blade. Peak stresses in an inserted rotor blade occur in

the neck region of the attachment perpendicular to the radial direction. The strength of a machined

ceramic is greatest parallel to the direction of grinding (subsequently referred to as the longitudinal

direction) and weakest perpendicular to the direction of grinding (the transverse direction) [ 2 ] . For NC132 HPSN surface-ground with 320-grit diamond and tested in quarter-point bending with a 0.125-by0. 250-inch cross section over a 1.5-inch outer span, the longitudinal strength is typically 100 KSI +15 KSI

and the transverse strength is 62 KSI +5 KSI. To achieve optimum strength for a rotor blade, therefore,

the highly stressed attachment region should be machined in the radial direction. Since this is not

possible by most of the machining approaches, the damage resulting from machining must be minimized

by using finer grit diamond for the final machining passes, followed by lapping.

The following sections describe the techniques that have been evaluated for rotor-blade machining.

4.1 Attachment contour grinding

Contour grinding of the blade attachment was conducted at Arthur Crafts Company (ARCO) under

the AiResearch subcontract to Norton Company. A schematic of the contour grinding technique is illustrated in figure 5. In this technique, a manual machine is equipped with a shadow lamp, which

projects an enlarged image of the work piece and grinding wheel on a screen. With the aid of this

projection, the operator is able to manipulate the machine to produce the proper shape. A very high

degree of accuracy may be obtained by a skilled operator. Typically, blade attachment specimens are rough machined at a down-feed rate of 0.0005 inch per pass, and finish machined at a rate of 0.0002 inch

per pass using a finer grit diamond wheel.

SHAPE TO BE

PRODUCED

MYLAR CHART

SHADOW COMPARATOR SCREEN

Figure 5. Contour Grinding.

The contour grinding process uses very small amounts of diamond, since dressing and reshaping of

the wheel are not as necessary as in some other techniques. However, contour grinding is still a very

expensive approach since an operator must be present continuously. An additional disadvantage of

contour grinding is that all the grinding grooves are aligned in the axial direction of the blade and are

thus perpendicular to the axis of maximum tensile stress. This is comparable to 4-point flexure test

specimens machined in the transverse direction.

213
The contour grinding approach has produced rotor blade attachments to the required tolerances

which, after lapping, have survived 130-percent design-speed spin testing.

4.2 Attachment formed-wheel grinding

Formed-wheel grinding is performed on a surface grinding machine using a diamond wheel that has

the exact mirror image of the shape to be produced. Typically, for a shape such as the blade

attachment, a series of formed wheels are used, with each subsequent wheel having a configuration

closer to the final shape and a finer grit size to achieve an improved surface finish.

The initial cost of formed wheels is high because they must be precision shaped, and during

machining they require frequent reshaping. In spite of these difficulties, formed-wheel configurations

and grit sizes can be chosen to allow very rapid stock removal to produce an overall low-cost machining

operation. As with contour grinding, formed-wheel grinding produces grinding grooves in the axial

direction of the blade, which is transverse to the principal stresses. This results in reduced material

strength. Lapping is required to minimize the strength degradation.

Acceptable blade attachments have been produced by formed-wheel grinding, and the cost

projections are very encouraging.

4.3 Attachment profile grinding at WESGQ

WESGO has machined the attachment sections on both spin-test specimens and turbine blades using a combination of form-wheel grinding and profile grinding. The attachment specimens were first ground

to within 0.010 inch of net size using a formed wheel. Final machining was accomplished on a singlespindle, numerically controlled, profile-grinding machine, resulting in orientation of grinding grooves in a

radial direction. The surfaces were then lapped and polished. Spin-test specimens and blades were

within required tolerances and have survived the 130-percent design speed spin testing.

4.4 Airfoil cam grinding

Initial airfoil fabrication studies by Norton Company were conducted at Excello using cam grinding

techniques that had been successfully used to fabricate a large ceramic airfoil configuration for

Westinghouse. However, machining was more difficult for the rotor blades of this program because of

their complex geometry, small size, and thin trailing edges.

A schematic of the basic cam grinding setup is shown in figure 6. This was not successful in

achieving the required geometry. Support of the airfoil at the tip severely limited the size and angle of

approach of the grinding wheel. This prevented grinding the base of the airfoil to tolerance. The set-up was modified, removing the blade-tip support and machining the blade cantilevered. This permitted a

smaller angle of attack of the grinding wheel, eliminating the initial restrictions. Excello demonstrated

the feasibility of airfoil machining but chose not to continue further.

Figure 6. Schematic of Cam Grinding Technique.

4.5 Airfoil profile grinding at Jarvis

Norton Company teamed with Jarvis Company to machine airfoils after Excello chose not to

continue. Jarvis used a two-spindle automatic profile grinding machine. Rough grinding was conducted

214
with a 2-inch diameter diamond wheel followed by final grinding with a 0.5-inch diameter wheel.

Airfoils were successfully machined onto attachments supplied by Arthur Crafts, and the complete

blades were successfully spin tested to the required 130-percent of design speed.

4.6 AiResearch rotor blade machining

AiResearch has conducted process-development studies of both attachment and airfoil machining

using a Gorton P-2-3 Pantograph Profile Grinder. The setup is shown schematically in figure 7 for

airfoil machining and pictorially in figure 8 for attachment machining. In both cases an 8:1 master is used to achieve the required close tolerances.

COPY MASTER

Figure 7. Schematic of Pantograph Profile Grinding Setup for Airfoil.

Figure 8. Pantograph Setup for Grinding Blade Attachment.

A major advantage of the Pantograph approach is the capability of aligning all the grinding grooves

radially in the attachment region. This results in the major tensile stresses being parallel to the grinding

grooves during engine operation. As previously discussed, this represents the condition that will have the

least detrimental effect on the material strength.

215
To evaluate the strength achieved by Pantograph profile grinding and to help define techniques to improve the strength, 4-point bend-test specimens were machined with the Pantograph, simulating the

same surface finish and orientation as in the blade attachment region. Initial specimens were machined

with a 60,000 rpm spindle varying the final diamond grit size. A 220-grit finish resulted in an average

strength (in the longitudinal direction) of 73 KSI. A 320-grit finish resulted in an average strength of

84.5 KSI. Both of these strengths were lower than typically measured for specimens prepared by

standard surface grinding using comparable diamond grit size. This indicated that the Pantograph

procedure was producing more severe surface damage. The equipment was modified to permit a spindle

speed of 100,000 rpm, resulting in an average strength of 98.8 KSI, comparable with standard surface

grinding. Slight further improvements occurred as the Pantograph operators became more experienced,

and as the diamond tooling was optimized. Additional significant strength improvement (>110 KSI) was

achieved by a post-machining oxidation treatment which either blunted or removed surface cracks

resulting from machining [3, 4].

Initial-blade attachment specimens were machined by a multi-step approach. A blank was cut from

a 9-inch-diameter billet and rough-machined on a surface grinder to remove as much stock as possible.

The remaining stock was then removed by Pantograph machining using 220-grit wheels for rough passes

j and 320-grit for finishing. This procedure was slow and expensive. Table 1 summarizes material removal ! rates for these various machining operations. Removal rates are extremely low for the profile grinding ; operations, which is mostly due to the small size of the diamond wheel and the minimal area of contact | with the workpiece.

TABLE I. MATERIAL REMOVAL RATES AND OPERATIONS.

OPERATION MACHINE USED TOOL USED

WHEEL AND/OR

GRIT SIZE IN.3/HR

GRIND BILLET

TO THICKNESS

SURFACE GRINDER DIAMOND WHEEL 1/2 X 6 IN. - 150 GRIT

2.4

SLICE BILLET SURFACE GRINDER DIAMOND WHEEL 0.050 X 6 IN. - 150 GRIT

1.2

FORM-GRIND

ATTACHMENT

SURFACE GRINDER FORMED

DIAMOND WHEEL

1/2 X 6 IN. - 150 GRIT

1.4

PROFILE-GRIND

ATTACHMENT

(ROUGH)

PANTOGRAPH MOUNTED

DIAMOND WHEEL

1/4 X 1/4 IN. - 150 GRIT

0.040

PROFILE-GRIND

ATTACHMENT

(FINISH)

PANTOGRAPH MOUNTED

DIAMOND WHEEL

1/4 X 3/16 IN. - 320 GRIT

0.005

TREPAN AIRFOIL

SHAPE

ULTRASONIC SHAPED TUBE 320 GRIT

B4C SLURRY

1.4

ROUGH-GRIND

AIRFOIL

PANTOGRAPH MOUNTED

DIAMOND WHEEL

1/8 X 5/8 IN. - 150 GRIT

0.02

FINISH-GRIND

AIRFOIL

PANTOGRAPH MOUNTED

DIAMOND WHEEL

1/8 X 7/16 IN. - 320 GRIT

0.01

:

! pi Ultrasonic trepanning was developed to remove stock that could not be removed by conventional I si diamond grinding, and thus reduce the stock removal required by profile grinding. A schematic of the ji ultrasonic system is shown in figure 9. j «

Trepanning is analogous to cookie-cutting in that a thin-walled tool produces the cut, avoiding the

necessity of grinding away the adjacent material. Figure 10 shows the trepanning tool used to ultrasonically machine a blade-attachment spin-test specimen directly from a solid HPSN billet in a

single step to within 0.010 inch of the required tolerance. This resulted in a significant cost reduction by

reducing machine time and operator time. In addition, tooling cost was reduced because the cutting was

done with boron carbide (B.C) abrasive which is much less expensive than diamond tooling grit-abrasive. I (ij

216
PUMP

ABRASIVE

SLURRY TANK

Figure 9. Ultrasonic Machining System.

Figure 10. Ultrasonic Trepanning Tool and Blade Attachment

Machined by Ultrasonic Trepanning.

Ultrasonic trepanning has been further refined to significantly decrease the time and cost of

machining complete rotor blades. The steps are shown in figure 11, First, ultrasonic trepanning

produces a blank containing the attachment to within about 0.010 inch, plus adequate stock from which

the airfoil and platform can be machined. An airfoil envelope is then ultrasonically trepanned to the

proper depth 90 degrees to the initial cut. A thin diamond slicing wheel is then used to cut in from the

sides and intersect the bottom of the trepanning cut. This removes a large amount of stock without

grinding, and thus yields a shaped blank that requires relatively little finishing by profile grinding. The

completed blade is shown in step 4.

As part of the machining process development, the strength of an ultrasonically machined 4-point

bend specimen was evaluated. Two different ultrasonic tool motions were evaluated, one referred to as transverse (as would occur in trepanning) and one referred to as frontal. Surfaces machined

ultrasonically had a nondirectional surface texture compared to the highly directional surface produced

by diamond surface grinding. This is clearly visible with a scanning electron microscope as shown in

figure 12. Figure 13 compares the strength of ultrasonically machined HPSN with specimens prepared by

217
standard surface grinding in the longitudinal and transverse directions, and by profile grinding in the

longitudinal direction. Based upon this and other data, the strength of the ultrasonically-machined HPSN

appears unaffected by grinding direction, and is comparable to longitudinal grinding by the other

approaches.

SHELL BLADE FINISHSEPARATED GROUND BY

WITH DIAMOND PANTOGRAPH

SLICING WHEEL

Figure 11. Reduced-Cost Blade Machining Using Ultrasonic Trepanning

and Pantograph Profile Grinding.

218
TRANSVERSE ULTRASONIC ON LEFT

(320-GRIT BORON CARBIDE)

CONVENTIONAL 320- GRIT DIAMOND

GRINDING ON RIGHT

FRONTAL ULTRASONIC ON LEFT

(320-GRIT BORON CARBIDE)

CONVENTIONAL 320-GRIT DIAMOND

GRINDING ON RIGHT

TRANSVERSE ULTRASONIC ON LEFT

(600-GRIT BORON CARBIDE)

CONVENTIONAL 320-GRIT DIAMOND

GRINDING ON RIGHT

ULTRASONIC EDGE-WASH OCCURRING

DUE TO UNOPTIMIZED TOOLING

AND PROCEDURES

320-GRIT BORON CARBIDE

Figure 12. Typical Surfaces Prepared by Conventional Diamond Grinding and

by Ultrasonic Machining.

T CONTROL

L—LONGITUDINAL

T-TRANSVERSE

L CONTROL

T ULTRASONIC USING 320-GRIT B4C

T ULTRASONIC USING 600-GRIT B4C

40 50 60 70 80 90 100 110 120

FOUR-POINT FLEXURE STRENGTH (ksi) 95-PERCENT CONFIDENCE INTERVALS

Figure 13. Strength Versus Machining Technique.

219
5. Conclusions

Machining of complex rotor blades to the required close tolerances has been successfully

accomplished at AiReseareh and other sources. Fabrication time and costs have been significantly

reduced by development of formed-wheel grinding and ultrasonic machining techniques.

Material strengths resulting from the various machining techniques have been measured, and the

data has been used to modify the machining procedures to achieve improved strength. Pantograph

profile grinding results have been significantly improved by this iterative procedure.

The HPSN ceramic rotor blades resulting from this machining process development effort have

been successfully tested with high yield at 130-percent of design speed for the intended engine

application.

The authors wish to thank the Defense Advanced Research Projects Agency for supporting this

effort.

References

[1] Torti, M. L., Baker, S.H., Tuffs, S. and Olson, B.A., Norton process optimization studies for hot

pressed silicon nitride blades and reaction sintered silicon nitride stator and static components,

Proceedings of the 1977 DARPA/NAVSEA Ceramic Gas Turbine Demonstration Engine Program

Review, MCIC-78-36 pp. 341-363 (March 1978).

[2] Mecholsky, J. J., Freiman, S. W. and Rice, R. W., Effect of grinding on flaw geometry and fracture

stress of glass, J. Am. Ceram. Soc. 60 3-4 pp. 114-117 (1977).

[3] Richerson, D. W. and Yonushonis, T. M., The effect of surface condition on the strength of hot

pressed silicon nitride, presented at the fall meeting of the American Ceramic Society, San

Francisco, California, (November 1, 1976).

[4] Richerson, D. W., Schuldies, J. J., Yonushonis, T. M., and Johansen, K. M., ARPA/Navy ceramic

engine materials and process development summary, in Ceramics for High Performance

Applications—n, ed. by John J. Burke, et al., pp. 625-650 (Brook Hill Publishing Co., Chestnut Hill,

MA, 1978).

220
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

MACHINABILITY STUDIES ON MACORTM GIASS-CERAMIC

D. G. Grossman and D. L. Taylor

Research and Development Laboratories

Corning Glass Works, Corning, N. Y., 14830

MACORTM glass-ceramic is a class of ceramic material which lends

itself to conventional machining tools and equipment. This machinability is attributed to the unique microstructure which consists of

a highly interlocked array of plate-like mica crystals dispersed in

a glassy matrix. The machining of this material is unlike machining

ductile metals and this paper makes some comparisons of the energy

requirements, cutting forces and tool wear in turning and drilling

operations.

Key words: Corning code 9658; glass -ceramic ; machinable; MACORTM

1. Introduction

The word machinable as applied to a ceramic material most often connotes

the use of abrasive techniques such as grinding. With MACORTM machinable

g lass- ceramic , parts are mostly shaped by single point turning, milling,

drilling, sawing, etc. using metal-working tools. Evaluation of this material in these operations is made through comparisons with metals and plastics

for, even though MACORTM is a brittle ceramic, its unique microstructure

allows it to be shaped with the ease of conventional machining. Furthermore,

it is ready for immediate use after machining because it does not require

the subsequent firing needed by ceramic materials that are machinable in the

"green" or unfired state.

MACORTM glass-ceramic is distinguished from most other ceramics by its

microstructure, which consists of an interlocked array of plate-like mica

crystals dispersed throughout a glassy matrix. A description of the microstructural development has been previously described through a series of

photomicrographs [1J. The "house of cards" microstructure is the key to

MACOR' sTM machinability . Like natural mica sheet, it is difficult to propagate fractures through the thin, flat mica crystals. Typically, fractures

follow the glass -mica interfaces and cleavage planes of the mica as revealed

in figure 1. The microscopic fractures are repeatedly deflected, branched

and blunted - in effect, turned around and arrested. Since the fractures

are localized, MACORTM glass-ceramic does not sacrifice strength or hardness

for the sake of its machinability.

Evaluation of machinability can be made using the following considerations: energy (forces, temperature) tool life (wear) and work result (stock

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

221
Figure 1. Crack deflection by mica crystals in MACORTM glass

ceramic . White marker = 10 ym.

removal rate, precision, surface quality). Since MACORTM, Corning Code 9658,

machinable glass-ceramic is a relatively new material, particular emphasis

was given to tool design in an effort to establish the optimum machining

conditions under various operations.

2 . Power and Energy Requirements

The unit power required to remove a specific volume of material was

measured using a drill dynamometer which simultaneously monitored the torque

and thrust imparted to the work piece through two sets of strain gages mount

ed on the shank of the test stand. The arrangement is shown in figure 2.

A split collet specimen holder is joined to the base by a hollow torque tube

of 5/8 inch O.D. and 1/2 inch I.D. A pair of two-element 90° torque sensitive strain gages is mounted so that they face each other diametrically on

the tube. A second pair of two-element gages is mounted in the same plane

at 90° and respond respectively to the axial strain and circumferential

(Poisson) strain. Both sensors act as four-arm bridges. Calibration for

both torque and thrust was accomplished using certified weights. The crosssensitivity between torque and thrust outputs was determined to be less than

1%. The output signals from the strain gage conditioning system were fed to

a datalogger.

The unit power requirements, P, for drilling can be calculated from the

following equation [2]:

Autodata Nine datalogger with VISIG II signal conditioning system, Acurex

Autodata, Mountain View, California.

222
TEST

PIECE

Figure 2. Drill dynamometer.

n STRAIN

GAGES

BASE

P = 1

50,000 d

2 f

where

P = Unit power requirements in drilling, hp/in3/min.

T = Torque, in- lb.

d = Diameter of drill, in.

f = Feed, in/ rev.

Note that only the torque enters the power equation since the horsepower due to thrust accounts for a small percent of the total power

developed [3]

The test stand was accurately located on the center line of a Bridgeport

drill press equipped with a variable feed rate. The drill speed was held

constant at 300 rpm or 20 fpm using 1/4 inch diameter drills.

A number of different drill styles including high-speed steel, solidcarbide, carbide-tipped, masonry, etc. were compared using the torque and

thrust data shown in table 1„ A somewhat surprising result was that masonry

drills designed for brittle materials performed poorly on MACORTM glassceramic. This new machinable ceramic responded particularly well to ordinary

high-speed steel twist drills. Where wear or tool life is a problem, the

carbide-tipped twist drills also developed low forces at the point of the

drill.

Drill sharpness was also investigated to examine both the effect of

wear and whether grabbing with a sharp drill is a problem as is the case

when drilling brass. Figure 3 indicates less torque is imparted to the

stock by a sharp drill but in both conditions, stability was attained after

the drill penetrated the glass-ceramic. Figure 4 indicates no particular

benefit in reduction of force from lubrication during drilling. Water-soluble coolants have been shown useful in other operations, such as turning,

particularly to trap the loose powder which is released.

223
Table 1.

Torque and Thrust at 20 Second Drill Time

MATERIAL DRILL STYLE TORQUE

in- lb

THRUST

lb

MACORTM #917 High-speed steel

surface treated

1.65 27 .

5

#924 Plastics 1.8 19.

#957 Hi -Helix 1.5 24.

#960 Lo-Helix 1.5 24.

#4947 Cotter pin w/135°

split point

3.5 36.

#77653 CID style carbide

tipped

1.75 28.

Style 400 solid-carbide

twist drill

1.5 24.

Style HD straight flute 6.5 65.

Style KK masonry 8. 64.

DULL DRILL VS SHARP DRILL

Time, sec Time, sec

Cleveland 1/4" High Speed Drill Cleveland 1/4" High Speed Drill

118° Point Angle *917 118° Point Angle *917

6 Mil/Rev Feed Rate,300RPM 6 Mil/Rev Feed Rate, 300 RPM

Figure 3. Effect of drill sharpness. Figure 4. Effect of lubrication.

To examine the effect of point angle, a series of high-speed drills was

ground with angles from 90° to 140°. The forces at the point of the drill

remained fairly constant over this range; therefore, it was unnecessary to

224
apply a special point angle for a particular drill.

The form of the torque, T and thrust, B, equations are [4]:

T = Cf

k

d

A B = Kfm

d

n

where C and K are constants, the values of which depend on the material,

drill and cutting fluid; f is the feed per rev; d is drill diameter; and k,

l, m and n are exponents the values of which depend primarily on the material

being drilled. For most steels the value of k = 0.78 and the value of i = 1.8. For cast iron, k = 0.60 and X = 1.7. A series of drilling experiments

varying the feed rate and drill diameter for MACORTM glass-ceramic gave exponents which more closely follow the brittle metals as shown in table 2.

The slopes of the lines for torque and thrust are much lower for the machinable glass-ceramic. Figures 5 and 6 show the torque as a function of feed

rate and drill diameter, respectively, for MACORTM machinable glass-ceramic.

Table 2

Constants for Torque and Thrust Equations

MATERIAL EXPONENTS CONSTANTS

k & m n C K

MACORTM Code 9658 0.66 1.67 0.71 0.60 39.4 2,305

Leaded brass 0.73 1.9 0.60 1.0 418. 6,938

screwstock

S.A.E. 33 aluminum 0.83 1.9 1.1 1.2 646. 75,030

Cast iron 0.60 1.7 0.73 370. 160,000

S.A.E. 1020 steel 0.78 1.8 0.87 2,186 906,000

Carbon tool steel, 0.78 1.8 0.87 2,089 933,000

annealed

TORQUE VS FEED RATE

MACORTM

.01 .02 .03

FEED RATE (IN/REV)-66

Figure 5. Torque-feed rate relationship for MACORTM glass-ceramic.

TORQUE VS DRILL DIA.

MACORTM

0.1 0.2 0.3 0.4

DRILL DIA. (IN) 1 - 67

Figure 6. Torque-drill diameter relationship for MACORTM glass

ceramic .

225
The average power consumed during drilling can be compared for a number

of machinable materials by the bar graph shown in figure 7. Note that all

the materials were drilled using high-speed drills at a feed rate of 6 mil/

rev and a speed of 300 rpm. It is interesting to compare these results with

tests of stock removal rate described in an earlier publication [1] as

"Machinability Index" „ The test measures the cut depth using a weighted

hacksaw for a given number of strokes . The values are arbitrarily scaled

against that of a standard, namely cold-rolled 1018 steel, assigned a value

of 111 on the scale. The scale is such that the index value increases as the

difficulty of machining increases. The machinability index values are compared to the unit power consumed in drilling in table 3. The correlation

between these values indicates that the tool's efficiency does not vary widely over the range of materials investigated.

AVERAGE POWER DURING DRILLING

x 0.8

5 0.4

t 0.2

Figure 7. Average power during drilling for various materials.

Table 3

Power, Torque and Thrust at 20 sec. Drill Time

MATERIAL POWER TORQUE THRUST MI*

hp/ in. 3 /min in- lb lb

Graphite 0.03 0.5 5. 1.

Teflon® TFE 0.06 1. 7.5 2.5

MACORTM Code 9658 0.11 2. 40. 25.

Brass -free machining 0.35 6.5 52. 36.

alloy 2024

Aluminum T-4 0.48 9. 90. 50.

1018 Steel cold-rolled 0.75 14. 122. 111.

Copper alloy 10 0.92 17. 168. 97.

#304 Stainless steel 1.00 19. 260. 229.

Machinability Index is an arbitrary unit and increases with difficulty

of machining.

226
3. Tool Life

Frequently the machinability of a material is assessed by how hard it

is on the tool, for example, number of holes completed per drill bit or

inches of work travel per milling tooth or simply the time taken to produce

a set amount of wear on the face of the tool. Tool life does affect production economics through tool costs and change time. However, for MACORTM

machinable glas s -ceramic , tool life is not the limiting factor in machining.

Instead, surface finish requirements set constraints on speeds and feed

rates

In single point turning, for example, feeds and speeds that produce

acceptable surface finishes do not cause excessive wear on either high-speed

steel or carbide tools. The wear on the tool body was measured under deliberately harsh conditions for both high-speed steel (fig. 8) and carbide

(fig. 9) tools.

How long a tool will last depends on factors other than just the material being machined; tool design and parameters are also important. This

effect is shown dramatically by altering the side cutting-edge angle using

C2-carbide tooling. Figure 10 shows that the allowable cutting speed, that

is, the maximum surface speed before the onset of chipping, may be increased

markedly with higher side cutting-edge angles.

Another factor in tool life is lubrication. The use of water-soluble

coolants such as Supercut S67, Johnsons Cold Stream or Quaker 103 is strongly

recommended not only to prolong tool life but also to achieve better results

These materials improve the cutting action and also trap and wash away the

loose powder produced during machining.

OlOiS

008

l±J I

o§ 006

<

g 004-

TURNING MACORTM MACHINABLE GLASS - CERAMIC

.000

HIGH SPEED STEEL TOOL

TOOL

SR

BRECEA

SCEA

RELIEF

J L 10

FEED

CUTTING SPEED DEPTH

CUTTING FLUID

.009 in /rev 45 ft/min .062 in SUPERCUT

20 30 40 50 60

TIME (MINUTES)

70 80 90

J

100

Figure 8. High-speed steel tool

wear in turning under deliberately

harsh conditions.

OI2i—

.010

008

tuIo

2

006 <

004 oo

002

TURNING MACORTM MACHINABLE GLASS - CERAMIC

CARBIDE TOOL

FEED .009 in/rev CUTTING SPEED 100 ft/min OEPTH .100 in CUTTING FLUID NONE

TOOL

SR

BRECEA

SCEA

RELIEF

CARBOLOY 370 (C-6) 0°

I" 0°0"

Figure 9. Carbide-tool wear in turning

under deliberately harsh conditions.

12 18 24 30

TIME (MINUTES)

36 42

_J

48

227
4 . Work Results

The effect on surface roughness was investigated for a variety of

machining parameters in single point turning. The surface quality is

simultaneously controlled by the feed rate, depth of cut and cutting velocity

The manner of material removal can be divided into two classes: Mode I results from lower cutting forces; under dry machining conditions it occurs

when the velocity ^ 400 sfpm, depth of cut ^ 0.100 in. and the feed S 0.0035

ipr. This form of material removal results in a surface quality represented

by the measured profile shown in figure 11(a) . Mode II results when the

cutting forces exceed a given value.

TURNING MACORTM MACHINABLE GLASSCERAMIC EFFECT OF TOOL GEOMETRY

800

~700

Depth-. 100 in Feed - .005 in /rev

Cutting Fluid- Supercut

Tool :Carboloy 523 (C2)

Rake:0°to+5°

Relief :0°

_L 15 30 45

SIDE CUTTING EDGE ANGLE"

SURFACE ROUGHNESS OF MACORTM

MACHINABLE GLASS- CERAMIC

V 100 SfptH d 0.025 in

f 0.0025 ipr

CO

o

S

O

O.lin

i- -i

b)

V =100 SfpM

d 0.025 in

f 0.0035ipr

v 100 SfpM

d 0.025 in

f =0.0050 ipr

Figure 10. Effect of tool geometry

in turning MACORTM machinable glassceramic

Figure 11. Effect of machining parameters on the surface quality of

MACORTM glass-ceramic in turning .

With higher cutting forces, random but densely populated imperfections

are produced on the surface. The onset of this condition for dry machining

is depicted in figure 11(b) . The population of these imperfections increases

with still higher cutting forces as shown in figure 11(c) . The two modes of

material removal operating simultaneously at higher cutting forces result in

a bimodal distribution of the collected chips. The ratio of the number of

small chips to the number of large chips is definitely related to the machining parameters

.

The use of a wet environment during turning allows for greater speeds

and feed rates without sacrificing surface quality. It is quite probable

that the Mode II material removal mechanism is thermal-shock related.

228
The authors wish to acknowledge the assistance of E. G. Turker of the

Industrial Engineering Department, Cleveland State University who conducted

the surface roughness experiments. We should also like to thank the members

of the Machine Shop at Technical Staffs Services of Corning Glass Works who

helped us with the drilling experiments.

References

[1] Grossman, D. G., Machining a machinable glass -ceramic , Vacuum 28, 55-61

(1978).

[2] Anon., Power requirements and coefficient of friction in machining,

Machining Characteristics of High Strength Thermal Resistance Materials , pp. 129-137 (AMC Technical Report 60-7-532, Curt is s -Wright Corp.,

Cincinnati, Ohio, 1960).

[3] Boston, 0. W., Power required in drilling, Metal Processing , pp. 268-

280 (John Wiley & Sons, Inc., New York, N.Y., 1941).

[4] Marks, L. S., Torque, thrust, and power in drilling, Mechanical Engineers Handbook, 4th Ed., pp. 1828-9 (McGraw Hill Inc., New York, N.Y.,

1941) .

!

229

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

EFFECTS OF VARIOUS POLISHING MEDIA AND TECHNIQUES ON

THE SURFACE FINISH AND BEHAVIOR OF LASER GLASSES*

Richard L. Landingham, Alton W. Casey and Roy 0. Lindahl

Lawrence Livermore Laboratory

Livermore, California 94550

The advance of high-power laser technology is dependent on the rate of advancement in laser glass forming and surface preparation. The threshold damage

of glass surfaces continues to be a weak link in the overall advancement of

laser technology. Methods were developed and used in the evaluation of existing glass surface preparation techniques. Modified procedures were evaluated

to reduce surface contamination and subsurface defects. Polishing rates were monitored under controlled polishing conditions (purity, pH, particle size

distribution, particle concentration, etc.). Future work at LLL for this on- going investigation are described.

Key Words: Laser glass; surface preparation; polishing compounds; particle

characterization; surface contamination; and subsurface defects.

1. Introduction

High-power lasers are used to implode D-T targets at LLL. Results from such implosions

demonstrate the feasibility of generating electrical power by inertial confinement fusion.

The 30-TW SHIVA laser (the world's most powerful laser) was designed and built at LLL for

this purpose. This first design used neodymi um-doped oxide glass as the amplifying medium.

A typical laser chain (50-m long) is shown in figure 1. Large increases in laser power are anticipated by using fluoride-base laser glass instead of oxide-base glass in the next upgrade: 200-to 300-TW NOVA solid-state laser design. Such improvements are possible because

fluoride amplifiers will have high gain and low nonlinear refraction index. High gain is

achieved with optimum balance between emission cross section of the lasing ion and duration

of the laser pulse. A smaller cross section is generally needed with increasing pulse

width. The emission cross section is host-dependent.

The nonlinear refractive index is most important to the performance of short-pulse

high-power lasers. Glasses with low linear refractive index will also have lower nonlinear

index coefficients. Glasses with low coefficients propagate laser beams with less distortion. Classes of glasses in order of decreasing index are silicates, phosphates, fluorosilicates, fluorophosphates, and fluorides (BeF? is the lowest). The advanced laser designs will will need these fluoride glasses, which are now under intense development. Both

passive components (lenses and windows) and active components (amplifiers and Faradayrotation media) of the fluorophosphate glasses (LG-810, E-181, LHG-10, etc.) are being

developed commercially for the NOVA design. Scaleup problems similar to those encountered

in scaling up the silicate glasses for the large high-power lasers are anticipated.

A more persistent problem common to these glasses is their low resistance to laser damage at their surfaces. This low damage threshold is reduced further when these surfaces

are coated. Properly controlled laser damage tests generally confirm that, the damage

threshold decreases in the following order: bulk glass (highest), bare entry surface, bare

exit surface, coated entry surface, and coated exit surface (lowest). Laser pulse duration

and wavelength can affect this order[l]. The lower damage threshold of the exit surface as compared to the entrance surface is attributed to a greater optical electric-field strength

231
Fig. 1. The installation of components in the Shiva laser system at the Lawrence

Livermore Laboratory.

at the exit surface[2]. Conflicting results and lack of consistent evidence still prevent

adequate explanation of the lower damage threshold for bare and coated surfaces as compared

to their bulk glass damage threshold. Surface damage is generally attributed to surface

finish and various techniques have been tried to obtain imperfection-free surfaces[3]. No

improvement in damage threshold is anticipated after achieving surface roughness (rms) of < 100 A[4]. Surface damage threshold could also be affected by a variety of other surface

related features as dust particles, surface films, embedded polishing particles, subsurface

scratches and cracks, and residual stresses[4-7].

An investigation was initiated this summer at LLL to evaluate the polishing procedures

used to prepare our laser glass components. In this document, we report some of our findings on polishing compounds, subsurface defects, surface chemistry of the polished glass,

and techniques developed to evaluate glass surfaces. Since this is an ongoing study, we

will also comment on the direction being taken as a result of existing data.

2. Procedure

This investigation is divided into three parts; (1) evaluation of existing processes,

(2) laboratory-scale experiments on modified processes, and (3) development of techniques

to analyze the glasses prepared by these processes. Effort is ongoing in all three areas.

The proprietary nature of most processes used in preparing the surfaces of laser

glass hampers our direct correlation of surface conditions with processing stages. The

surface conditions from well-established grinding and polishing practices are being evaluated by preparing these surfaces in our Optics Laboratory from glass provided by the laser

glass suppliers. Suprasil-II silica disks (50 mm diam, 10 mm thick) are being evaluated

Reference to a company or product name does not imply approval or recommendation of the

product by the University of California or the U. S. Department of Energy to the exclusion of others that might be suitable.

232
first. The surface damage introduced by coarse grinding and polishing operations is being

monitored quantitatively by the Ernsberger technique[9]. The surface from each state of

preparation is polished on cloth laps with fine polishing compound (< 1 pirn diam TiC^) to remove the surface roughness without introducing more deep subsurface defects. Ion ex- change or etching techniques are used to expose these subsurface defects and the concentration of defects per mm are determined by image analysis with a Model 720 Imanco image

analyzer. Alternate fine polishing and etching steps are used to determine the depth of

the defects to within a few micrometers.

The polishing compounds affect the degree of subsurface damage and surface impurity

level. The characterization of these compounds includes particle size distribution, sur- face area, true density, particle shape, phase identification, composition, and impurity

level. Impurity levels of other potential contamination sources (carrier solution, antifoam and dispersion agents, wheel pitch, dressing blocks, etc.) are also checked. The

impurity level of the polishing medium is monitored during and after polishing the glass.

The impurity levels of the glass are determined to establish the lowest attainable impurity

level on the polished surface.

Several analytical techniques were compared to establish a method of monitoring the

surface impurity level after the finer polishing and finishing stages. Ion microprobe,

Auger electron spectroscopy (AES), energy dispersive system (EDS), electron spectroscopy

for chemical analysis (ESCA), atomic absorption spectrometry, and atomic emission spectroscopy techniques were evaluated.

The characterization of the final surface will include roughness and flatness determinations, laser damage threshold tests, surface impurity analyses, and/or subsurface defect

tests. A unique laser damage facility has been developed at LLL to provide rapid and

reliable damage threshold results to accelerate the development of laser glass components

[1].

We have begun applying some of the experience gained in preparing the Suprasil-II

silica surfaces to the preparation of fluorophosphate glasses. The pH level in a highpurity polishing slurry was controlled to determine its effect on polishing rate.

3. Results and Discussion

Various final polishing compounds were characterized because most of the contamination

and defects on the surface of the glass should come from this final polishing stage (see

table I). The particle size distribution was larger than desired in every case although

the mean particle size was < 3 ^m. The large particles were primarily bonded agglomerates,

which should breakdown if recycled in the final polishing stage. The subsurface damage

a. Single large particles in MgO (Code 920) b. Large agglomerate particles in SnO(AV-116)

Fig. 2. Photomicrograph of typical large particles discovered in

final polishing compounds by scanning electron microscopy.

233
caused by hard agglomerates is difficult to polish out in this final stage. Neither large

single particles nor agglomerates (see figure 2) should be present in the final polishing

stage. We have removed such particles by three methods: (1) allowing large particles to

settle out in a liquid and decanting off smaller particles for polishing, (2) continuouscentrifugation of the polishing solution to remove large particles during polishing and

(3) filtering the polishing solution to remove large particles.

Table I. Glass Polishing Compounds.

Particle

Surface area, size

Density, Phases detected by X-ray range,

Compound g/cnr x-ray diffraction nr/g fluorescencef

nm

MgO (Code 920) 3.58 Cubic + unident

1 ines

26.84 0.01-15a

SnO (AV-116) 6.45 Tetrga Sn09 +

Sn304

11.27 0.01-10

ZnO (HSA) 4.84 Hexag. + unident

1 ines

50.79 0.01-30

TK68 6.77 Cefi W0i ?Typeb 2.82 Ce(s),La(s),W(vvw) 0.03-10

D . C\J ueu2 isj ueuriwj 0

"i . AOId U.Ul-10

Rhodite 50 5.91 Ce02 (s) CeF

3 (w) 12.49 Ce(s),Ca(m),W(vvw) 0.01-203

Rhodite 76 6.38 CE0

2 (s) Ce0F(w) 6.21 0.01-20

ZOX CE89 6.05 Ce6W0 12 Typeb 16.36 0.01-10

ZOX E 5.78 ZrO^ 14.59 0.01-10

Cerium 85 6.603 Ce6W012 (s) Typeb +

CeC03F

(w)

29.01 Ce(s),La(s),W(vvw) 0.01-7

Lustrox 1200 4.97 Zr0

2 (s) d + Zr02 (w) c 20.32 Zr(s) 0.01-12

Lustrox PG 3.41 Zr0

2 c 27.78 Zr(s) 0.01-20

ZnOe 5.61 Hexag. (5-660) 7.18 0.02-15

a Large particles are solid and not agglomerates of smaller particles as determined

for other powders listed.

b CegW0j

2 crystal type structure but La replaces W.

c (13-307) Zr02 monoclinic (variable brown to colorless) Baddeleyite.

d (14-534) Zr0

2 Tetragonal.

e Reagent grade powder (8832).

Relative intensities; s-strong, m-medium, vvw-very very weak.

234
No direct correlation between surface area and particle size distribution was observed

(see table I). The surface area of powders is primarily affected by the relatively small

amount of fines (< ljim) in the powder.

The densities of these powders varied both above and below the theoretical density of

the supposed major phase. Multiple phases or impurities generally contributed to these

types of shifts. Multiple phase compounds make the control of the polishing operation

more difficult unless the same ratio of phases can be obtained and behavior of both are understood during polishing operations. Impurities can alter polishing conditions and

contaminate the surface of the glass. We determined impurity levels of several compounds

being used to polish LLL laser glasses (see table II). Table II includes impurity levels

Table II. Impurities in polishing compounds as determined by emission spectrometry.

Impurity levels in polishing compounds, 9 ppm

Impurity

elements

MgO

Code 920

SnO

AV-116

Liquid

85b

ZOX

CE-89

A1 2°3

Microgrit

30T

A1 2°3

Microgrit

9T

Reagent

grade

ZnO

Al 300 < 5000 nd nd major major < 100

Cu 6 1 10 10 4 nd 3

Si 600 1000 nd nd 1 Wt% 1 Wt% < 2

Fe 200 600 >1000 1000 1000 1000 nd

Ca 3000 100 2000 20 3000 3000 <2

Zn < 30 100 nd < 1 <30 <30 major

Ni 10 60 nd <1 5 5 <3

Cr <10 30 nd 1000 <10 <10 < 10

Ag <1 <1 nd nd < 1 3 < 1

Sr 10 < 1 >1000 1000 100 1000 <1

V <10 <10 nd nd 40 20 < 30

Pb <10 <10 600 60 <30 <30 <10

La nd nd nd > 10 wt% • nd nd nd

Ba <100 <10 > 1 WW 5000 <30 <30 <100

Nb <30 <30 >1000 300 <30 <30 <100

Mn 10 < 1 30 20 5 5 <10

Mg nd nd >1000 10 500 500 < 1

Na <300 nd 100 3000 3000 nd nd

Ti 10 nd nd 800 800 nd < 10

Ga nd nd nd 30 30 nd nd

Zr <10 nd >1000 nd nd nd nd

a nd means not detected.

Y and Th level at 100-1000 according to x-ray fluorescence.

235
of two coarser polishing compounds also being used. The impurity levels of other potential sources of contamination are shown in table III. These impurity levels should be compared

with those in the bulk of the glass (see table IV) before selecting polishing materials.

The impurity levels of all these polishing materials are above those of the bulk glasses

so we can expect contamination of the surface and/or difficulty in getting reproducible

polishing conditions.

The surfaces of Suprasil-II silica disks were analyzed after final polishing with

Liquid 85 1 by a commercial vendor or with ZOX CE-89 in our Optics Laboratory. Surfaces

polished with Liquid 85 were highly contaminated with Ce, Pb, Na, Al , K, and Ca. Trace

amounts of Be, B, Mg, Ti , agd La were also detected by ion microprobe analyses at depths

under the surface of < 150 A. Ion microprobe analysis is not sensitive for P, Sc, and Fe

in the presence of Si molecular igns. Auger electron spectroscopy (AES) could only detect

C, P, S, CI, Ca, and N within 10 A of the surface. Photoelectron spectroscopy (PES) was

able to detect Ce, Sn, Ca, N, and C at the same, depth. Trace quantities of Al , Pb, and

water were detected at greater depths (^ 3200 A). As expected, none of these surface

impurities could be detected by the energy dispersive system (EDS) in the scanning electron

microscope (SEM). Embedded particles of > 0.1 urn could be identified by SEM and analyzed

by EDS when exposed at the surface.

The interpretation and turn-around time prevent the use of electron spectroscopy for

chemical analysis (ESCA) on routine surfaces. The sensitivity level for several elements

of interest is too low to use induction-coupled-plasma (ICP) with direct-reading spectrometry (DRS), flame atomic absorption, or furnace atomic absorption techniques. Evaluation

of these surfaces are still in progress by secondary ion mass spectrometry (SIMS).

Table III. Other sources of contamination during laser glass polishing operation.

Impurity Pitch Ever-F

elements 64"C melt 73°C melt 82 U C melt Sol uti

Al 3 2 2 1

Cu 0.1 0.1 0.1 <1

Si 10 2 2 2

Fe 3 10(60) a

7 1

Ca 3 1 2 2

Zn < 1 <K5) a

7 < 1

Ni 7 6(10) a

2 < 1

V 10 10(20) a

7 < 1

Pb <1 < K5) a < 1 2

Mg 1 0.6 0.7 < 2

Na < 6 < 6 < 6 25

Ti 2 < 0.3 < 0.3 < 1

K nd nd nd 25

a X-ray Fluorescence Analysis (xx)

(Note: nd means element was not detected)

Reference to a company or product name does not imply approval or recommendations of the

product by the University of California or the U. S. Department of Energy to the exclusion of others that might be suitable.

236
Table IV. Impurities in laser glasses as determined by emission spectrometry3

T mm i vi +"w el ements BK-7

Impurities, ppm

Laser glasses

LG-812 ED-2

Mg 100 major 10

Mo 200 2000 1000

Ba major 1000 60

Si major 20 major

Cu 3 6 3

Sr 60 major major

B major 100

Sn < 30 < 10 30

Ti 200 < 10 < 3

As 1000 < 30 300

a Major = additives or major elements of glass composition; blanks mean

not detected.

We have developed a rapid and more quantitive method to detect impurities on these

surfaces. We attach teflon fixture to the surface of a glass to contain an etchant

2% HF) within a 13 mm diameter area on the surface. Etchagt time and temperature

is controlled to dissolve the desired depth of glass (< 3000 A) for subsequent analysis.

The actual etch-out depth Q is profiled with a Det Tak surface analyzer (background threshold of approximately 200 A). The solution of dissolved glass is evaporated and analyzed

by by emission spectroscopy. Three surfaces polished with Z0X CE 89 were analyzed by

etching to depths of 1200, 2000, and 2600 A, respectively. Surface impurities detected

by this method include Na, Al , Ca, Cu, Mg, Fe, Ti, Ba, Ni, Pb, Cr, Mn, Sr, Zr, and Zn. No

Ce was detected although Ce02 was used to final polish these surfaces. This analytical

method is not as sensitive to Ce as the ion microprobe method. The bulk glass is contaminated with Al . Additional analyses and development of standards are in progress to

improve the accuracy of this method.

A second feature of the above method is that the etched region can be examined

optically. Surface and near-surface flaws are revealed by etching or ion exchange techniques[9]. The deep microcracks caused by rough surface removal operations (blanchard

grinding, diamond saw cuts, etc.) have patterns similar to those in figure 3. The shallow

defects of the coarser polishing operations (30 and 9 pm AI2O3 polish) used to remove the

previous defects are shown in figure 4. If proper amounts of the surface are removed during each stage of polishing, these deep defects will not be observed after etching a sur- face polished with fine (< 1 nm) polishing compounds. Etching such a finely polished sur- face does expose the polish scratches previously filled in and/or covered over by a sur- face film (see figure 5). The concentration of scratches and defects can be determined by

image analysis and compared with laser damage results to establish a quantitative correlation. Such information would be useful in establishing inspection limits on surfaces

before coating or bare-surface use.

The laboratory scale polishing unit (305 mm diam wheel) was set up to final polish

the Suprasil-II silica disks with high-purity polishing slurries. We stopped using pitch

on this wheel after we observed that the pitch reacted with the polishing slurry and was picked up on the glass. Nylon covered wheels allowed better monitoring of polishing rates

and impurity levels. Evaluation of surfaces suitable for commercial polishing applications

will be investigated after the questions about impurity effects have been resolved.

237
as-pol ished

etched

a. Defect pattern in diamond saw cut surface

as-pol i shed

etched

b. Defect patterns in blanchard ground surface

Fig. 3. Surface defects revealed on Suprasil-II silica surfaces

after etching with dulite HF

The characterization results are given in table V and VI for the high-purity powders

used on the laboratory-scale polishing studies. These powders were used on the final

polishing stage to remove the defects of the two previous coarse stages (A^Oo microgrit

30% and 9% in table II). To date, we have polished Suprasil-II silica disks (51 mm diam)

with MgO, CeC^ (B-1407), and both AI0O3 powder and a fluorophosphate glass (ZG-812) disk

(25 mm diam) with the ZnO powder. The adjusted polishing parameters and polishing rates

for these initial studies are listed in Table VII. Polishing with MgO was discontinued

because additives would be required to control excessive foaming. High polishing rates

with Ce02 powder were obtained after the coarse particles (> 5/im diam) were removed

while maintaining

slurry

a constant concentration of particles (< 5 mdi diam) in the polishing

238
as-pol ished

etched

a. Coarse polish with 30 pm A1 ? 0

? compounds

as-pol ished

etched

-*! 100pm

b. Coarse polished with 9 pm Al ^0^ compounds

Fig. 4. Surface defects revealed on Suprasil-II silica surface

after etching with dulite HF

239
as-polished with

CE-89

\ etched

g. 5 Surface scratches on Suprasil-II silica are apparent after etching with HF.

Table V,

Powder

type

Material characterization of high-purity powders under evaluation

as polishing compounds.

Particle Status of

Density, Phases detected Surface _ size large

g/cnr by x-ray diff. area, (m /2) ranges, nm particles9

ZnO 5.61

(Code 8832)

MgO(lOOm) 3.58

CeOo

B-1407

CeO? B-1408

CeOo

B-1409

Al ?0o

ERC

J

AloOo

Lot 4071

7.03

7.09

7.58b

3.97

3.93

Hexag.

Cubic + Moderate

Mg(0H)

2

F.C.C.

F.C.C.

F.C.C.

*** Alpha Al203+Unknown

1 ines

*** Alpha Al203+Unknown

1 ines

7.2

76.6

1.4

1.2

6.6

7.4

9.6

.02-15

.01-25

.05-20

.05-20

.05-25

.03-20

.03-18

A

A

A

A

A

A

A

a S = single particle, and A = agglomerate.

b After vacuum bakeout at 230°C this density dropped to 7.13 g/cnr*, but

the particle size range remained the same.

*** Unknown = unknown lines also detected.

240
Table VI. Impurities in high-purity powders under evaluation as polishing

compounds (as determined by emission spectrometry) 9

Impurity

piv. i pments t; i hv_ i i u o

ZnO

8832

MgO

100m

Impurities,

CeOo

B-1407

ppm

Ce02

B1408

CeOo

B-1409

A1 20

3

ERC-HPDBM

AloOo

4071

Cu 3 10 < 10 nd nd 4 1

Mq < 1 nd 6 6 < 2 100 100

Si < 2 100 nd nd nd 110 100

Ca < 2 3000 1000 nd nd 25 30

Fe < 10 300 15 < 5 < 5 30 30

Ga nd < 3 nd nd nd 100 20

Mn < 10 30 10 < 1 nd < 1 3

B < 3 300 100 nd nd < 3 < 3

Ba < 100 < 100 100 nd nd < 30 nd

Pb < 10 < 10 2 2 2 < 30 < 30

Al nd 200 nd nd nd major major

Sr nd 30 nd nd nd < 1 nd

nd mean not detected.

Polishing with Al 2O3 (ERC-HP-DBR) powder was stopped because of excessive foaming.

This foaming action was attributed to the Mg doping agent in this hot press powder.

Satisfactory polishing rates were obtained with the AI2O3 (4071 powder, but subsurface

scratches persisted after 35 h of polishing with the same slurry. Starting with a finer

agglomerate size (< 5 ^m diam) A^O^ slurry and continuously removing large agglomerates

(> 5 (im diam) from the slurry reduced subsurface scratches. Preliminary evidence indicates

a rapid buildup (snowballing) of agglomerates comprised of AI0O3 particles and silica from

the glass (see figure 6). This buildup was also observed by determining the particle distribution size of the slurry before and after polishing. Further evaluation of this buildup mechanism is in progress because it could account for the common belief that AI2O3 is

"too hard" to polish glass.

Our first attempt to polish fluorophosphate glass was made with ZnO (8832) powder with

controlled pH slurry (see table VII). At lower pH values (<7.0), the foaming was not ex- cessive, and a maximum polishing rate was observed at 6.5 pH. Fewer subsurface scratches

were observed with ZnO on the fluorophosphate glass than observed with Al on the silica

glass. Additional polishing studies are in progress on both glasses.

241
Table VII. Results vs adjusted parameters of laboratory scale polishing

with high purity powders.

Glass Removal Rate

Polishing Polishing Polishing Wt^ Thickness,

compound pH* time, h /ig/h jiiti/h Remarks

MgO (100m) 9.8-9.7

Ce0

2 (B-1407) 8.3-6.8

as received

Ce0

2 (B-1407) 6.8-7.2

as received

Ce0

2 (B-1407) 8.9-7.5

-5 fxm fraction

A1 2

0

3 ERC

HP-DBM

8.0-7.7

A1 2

0

3 (4071) 9.0-7.9

as received

A1

2

0

3 (4071) 8.3-7.4

-5 nm fraction

ZnO (8832)

as received

ZnO (8832)

as received

AnO (8832) as received

5.9-5.8

6.5-6.6

9.0-8.5

27.5

27.5

27.5

34.5

17.5

35 .

24.5

13

13

13

250

990

3,900

4,600

1,800

1,600

760.

0.11

0.44

0.48

2.06

0.82

0.71

0.34

420 1.80

780 3.32

500 2.15

Excessive Foaming

Control pH

Excessive foaming

Decreasing A1

2

0

3 cone.

(50%) due to decant of

large particles (>5/jm

dia)

.

Control pH

Control pH

Control pH. Excessive

foaming.

* pH values before and after polishing

Laser damage tests have begun to establish the threshold damage of Suprasil-II silica

surfaces after various polishing and etching procedures. A separate report will be prepared when we have sufficient results to show a direct relationship between specific

impurity levels, defects, and surface threshold damage.

4. Conclusions and Recommendations

The rapid advance of high-power lasers for fusion applications requires advances in

glass fabrication, glass surface preparation, and glass coating techniques, to sustain

laser stability and durability. Since the laser threshold damage of glasses seems to be

the weakest link, we examined existing surface preparation procedures as well as evaluated

alternate methods to reduce surface contamination and subsurface defects. Analytical techniques were developed to quantitatively determine surface impurities and subsurface defects.

The origin of many of these impurities and defects were identified.

242
Closer control of final polishing compounds is needed to assist opticians in meeting

the increasing constraints on surface finishes. Characterization of polishing compounds

should at least include impurity content, phase content, particle size distribution, and

nature of large particles and/or agglomerates. Such information could prevent problems in

subsequent polishing operations, which might trigger unwanted changes in those procedures.

The contamination of the polished surface by the polishing slurry has been confirmed.

Efforts to maintain a low contamination level in this slurry should improve polishing

operations as well as laser threshold damage. The chemical reactions at the glass-slurry

interface seem to control polishing rates as well as polishing particle breakdown or growth

growth (snowballing) rates. More work in this area is in progress at LLL.

5. Acknowledgements

This investigation was initiated and supported by Norman Brown of the Optics Laboratory

and Howard Lowdermilk and Lee Smith of the Laser Damage Group^at LLL. Their assistance

Fig. 6. Buildup (Snowballing) of large agglomerates in the final polishing

slurry were photographed in the SEM and analyzed by EDS.

These agglomerates were made up of aluminum, silicon, and oxygen and

formed during final polish of silica glass with A1 9 0, polishing

compounds. c 6

243
and guidance made this report possible. Dave Milam of the Laser Damage Group has been most

helpful in directing us in matters relating to and evaluation of laser damage tests. Al

Ulrich has prepared and evaluated many of the glass surfaces by metal lographic techniques.

Charlie Farrell developed the procedures used to detect impurities and defects on the sur- face of these glasses. The unique support of Carol Weaver and Chuck Slettevold was needed

in the characterization of all the polishing compounds.

6. References

[1] Milam, D., "Laser-induced Damage at 1064 nm, 125 psec", Applied Optics, Vol. 16,

May 1977, pp. 1204-1213.

[2] Crisp, M. D., Boling, N. L., and Dube, G. , "Importance of Fresnel Reflections in

Laser Surface Damage of Transparent Dielectrics", Appl . Phys. Lett., Vol. 21. No. 8,

October 15, 1972, pp. 364-366.

[3] Fradin, D. W. and Bass, M., "Comparison of Laser-Induced Surface and Bulk Damage",

Appl. Phys. Lett. Vol. 22, No. 4, February 15, 1973, pp. 157-159.

[4] Milam, D., Smith,, W. L., Weber, M. J., Genther, A. H., House, R. A., and Bettis,

J. R., "The Effects of Surface Roughness on 1064-nm, 150 ps Laser Damage", in Laser

Induced Damage in Optica , A. J. Glass and A. H. Guenther, eds., Nat. Bur. Stand.

(U.S.) Spec. Publ. 509, pp. 166-173.

[5] House, R. A., Guenther, A. H., and Bennett, J. M., "Surface Roughness Statistics of

Fused Silica as a Function of Surface Preparation and Treatment", in Laser Induced

Damage in Optical Materials: 1977 , A. J. Glass and A. H. Guenther, eds. Nat. Bur.

Stand. (U.S.) Spec. Publ. 509, pp. 157-165.

[6] Bloemberg, N., "Role of Cracks, Pores, and Absorbing Inclusions on Laser Induced Damage Threshold at Surfaces of Transparent Dielectrics", Applied Optics, Vol. 12, No. 4,

April, 1973, pp. 661-664.

[7] Ringlien, J. A., Boling, N. L., and Dube, G., "An Acid Treatment for Raising the Surface Damage Threshold of Laser Glass", Applied Physics Letter, Vol. 25, No. 10,

November 25, 1974, pp. 598-600.

[8] Boling, N. L., Ringlien, J. A., and Dube, G., "Q-Switched Laser Inducted Surface Damage at 1.06 Microns", in Laser Induced Damage in Optical Materials: 1974 ", edited by

A. J. Glass and A. H. Guenther, Nat. Bur. Stand. (U.S.) Spec. Publ. 414, December,

1974, pp. 119-130.

[9] Ernsberger, F. M. , "Origin and Detection of Microflows in Glass", Fracture Mechani cs of Ceramics, Volume I, Fractography , edited by R. C. Bradt, D. P. H. Hasselman, and

F. F. Lange, Plenum Press, N.Y. - London, 1974.

Discussion

BENNETT

The industry has a good feeling, I believe, for the depth of subsurface

damage introduced by various finishing techniques but not for the depth of

the impurity level added. If most of the impurities are only a few hundred

anystroms below the surface, perhaps they could be removed after final

polishing by a weak etch or by light sputter etching. We have succeeded

in removing several thousand anystroms of fused Si02 by sputter etching

without adversely affecting surface roughness, for example, although BK 7

glass does roughen using this technique . Would such an approach reduce

the need for high purity polishing compounds?

244
LANDINGHAM

o

The major amount of contamination was within the first 150 A of the

surface for quartz glass polished with CE 89 (Ce02

). This contamination

layer will vary with polishing conditions (materials being polished.)

In general, improvements have been observed in surface damage threshold

when the surface is etched. We are testing material that has been etched

to confirm this improvement. If impurities are the main problem and we can

remove them without causing other defects (like stress corrosion of defects) , the improvement in high-purity polishing may not be needed. I would expect some intermediate solution between both approaches will be

necessary

JOHNSON

Comment on the results of coolant pH - hydrogen ion effect in relation to

the previous paper results. One paper said No enhancement; the other paper said,

Yes, there is and it is H ion increasing brittle behavior.

LANDINGHAM

The interaction of the coolant with the surface of a workpiece is dependent on many conditions (pressure, temperature, time, etc). The previous

speakers are working in very different environments and observe different

behaviors. We are polishing at relatively low rates and pressures with

relatively chemical active surface so pH variations can influence our

results. At higher polishing rates with higher temperatures and pressure,

other parameters may become more important. We plan to investigate these

parameters

BOWLES

Has variation in surface quality been observed in polishing with media at

different pH?

LANDINGHAM

Under our present polishing conditions, the f luorophosphate glasses are

significantly affected by small shifts in pH. Such pronounced affects on

polishing rate will also alter the surface quality in some manner. The

affect is not independent of other polishing parameters and requires specific attention to the system (glass, polishing compounds, carrier, solution,

etc.) of interest.

245

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

THE FABRICATION OF ALUMINA TOOTH

ROOTS BY DIAMOND CONTOUR GRINDING

G. L. Messing, L. G. McCoy, and C. R. Hassler

BATTELLE

Columbus Laboratories

505 King Avenue

Columbus, Ohio 43201

To determine if the process of fabricating alumina tooth roots by diamond

contour grinding significantly reduces the strength of the tooth root, flexure

strengths were determined for four different grinding conditions. The grinding conditions studied were grinding direction (longitudinal or transverse) and

the sintered state of the alumina during grinding (bisque or fully sintered).

Only for one condition (transversely ground, sintered alumina), did the grinding process lead to the strength-limiting flaw. For the other conditions the

strength limiting flaw was entirely intrinsic (i.e., pores) or an interaction

of a pore with surface damage from grinding.

Key Words: Alumina; diamond contour grinding; flexure strength; prosthetic

device; serrated surface design; tooth roots.

1. Introduction

In the past 10 years there has been considerable interest in the possibility of re- placing missing teeth with a new generation of implant designs. In ceramics, the implant

is characteristically an alumina-based root structure designed to accept a precious metal or composite porcelain crown. Alumina is used because of its established biocompatibil ity

and good strength behavior in vivo . Artificial tooth roots had been fabricated in earlier

studies by slip casting [l]i, but because it is difficult to maintain current design details, this process has been replaced by diamond contour grinding [2]. Diamond contour

grinding is a process whereby either a sintered, bisque, or green preform is ground with a diamond tool that has the negative contour of the desired design. Contour grinding can be

cost and efficiency advantageous over other processes, such as injection molding, when

small lots of tooth roots, having a variety of sizes and shapes, are required.

At Battelle, a three-component prosthetic system was designed using an alumina root as the implanted component. The root has a serrated surface designed for stress distribution.

After implantation the serrations are filled with and stabilized by ingrown bone. In implant studies using a baboon model, only a single failure due to implant fracture has oc- curred out of ^1 00 implanted roots. To assess the cause of the failure and to assure the

viability of this process for fabrication of human implants, it is necessary to establish

whether the grinding operation significantly flaws the tooth root surface. It has been

well demonstrated in the literature [3] that the presence of surface grinding flaws can be

strength-limiting in ceramics. This could be a significant effect as is apparent in Figure

1 which shows the flawed surface of the serration in a contour ground root. Furthermore,

with the recent advent of a clinical study at The Ohio State University College of Dentistry

using BCL-designed and fabricated implants, it is necessary to demonstrate that tentative

ASTM standards [4] for high-density alumina implants could be met. These standards require

a flexural strength in three-point loading of 58 ksi (400 MNm"^).

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

247
Figure 1. Scanning electron micrograph of surface

damage (both uniform and non-uniform) on ground serration of a tooth root

2. Tooth Root Design and Fabrication

Before describing the grinding operation, it is informative to first describe the design of the tooth root. The tooth root is either elliptical or rectangular in cross section with a tapered, serrated surface (Figure 2). The depth of the serrations is 1.5 mm to

maximize the horizontal area available to support occlusal loads. To avoid a stress riser,

the inside of the serration is rounded. A 5-degree taper of the tooth structure permits

easier fitting into the vacant site. The top of the tooth root is recessed 2 mm and a hole

is drilled into the tooth root to a depth of 7 mm. A combination gold post and core is

investment cast to fit each implant and is cemented to the tooth root 3 months after it has

been implanted. This initial 3-month period allows for bone ingrowth into the serrations

and stabilization of the root structure. A gold crown is then fitted to the post and core

to provide correct occlusion.

Contour grinding was chosen as the process for fabricating this design because it

gives detail definition, structural reproducibility, and size and shape flexibility. The

process of contour grinding is used extensively in the ceramic industry to fabricate precision electrical insulator components. However, since the cross-sectional shape of the

tooth root is either elliptical or rectangular, a machining operation more involved than a simple lathe operation is required. For this purpose, a computer-controlled milling machine is used. To generate the elliptical and rectangular shapes, the alumina rod stock

blank is supported on an indexed rotary table. Variations in shape are generated by moving

the cutting tool relative to the x-y plane of the rotary table. Thus, it is a simple

matter to generate a tooth root having either an elliptical or rectangular cross section of

varying size.

To fabricate a tooth root, high-purity alumina powder 2 with 0.025 percent MgO is hydropressed at 345 MNm-2 and bisque fired at 1120 C for 2 hours. The bisqued alumina rod is

preground to a solid truncated cone having a 5-degree taper on each side, so that only a minimal amount of material (i.e., 2 to 2.5 mm) needs to be removed during contour grinding.

The serrations are ground in water with a 200-grit diamond-coated tool which has the

negative relief of the tooth root. The feed rate is approximately 12.7 cm/min with a

2Reynolds' Alumina RC-HP-DBM

248
Figure 2. Design details of a rectangular tooth root

maximum depth per pass of 0.5 mm. Consequently, the number of passes required to grind the

full depth of the serration is either 4 or 5, depending on the cross-sectional shape. A

high-speed (23,000 rpm), air-driven 3 chuck is required to attain a machining rate of 1,828

surface m/min (6,000 surface ft/min) with the present tool. There is a minimal loss in

speed if the tool is sharp and appropriate feed rates are maintained. To finish the tooth

root, which is ground 15 percent oversize to allow for sintering shrinkage, the recessed

top is ground with a conical diamond tool and a post hole is drilled. The tooth root is

then sintered for 1 hour at 1538°C. The density of the tooth root is ^99.5 percent of

theoretical and the average grain size is 3 to 4 micrometers.

3. Experimental

To determine whether the described grinding process and sintering procedure affects

the mechanical properties of the implants, rectangular test bars were prepared for fourpoint flexure experiments. All of the test bars were prepared by the same two-step sintering cycle described above. Half of the test bars were ground to size after the bisque

sinter and the other half were ground after the final sinter. Half of the bars in each

group were ground transversely (i.e., perpendicular to the tensile axis) to reflect the

processing conditions and the other half were ground longitudinally to simulate a control

condition. The four grinding conditions are:

(1) Bisque, transversely ground (BT)

(2) Bisque, longitudinally ground (BL)

(3) Sintered, transversely ground (ST)

(4) Sintered, longitudinally ground (SL).

To simulate the contour grinding operation, the test bars were ground in water with a 200-

grit diamond-coated grinding wheel

1*. The edges on the tensile side of the finished test

bars were rounded and polished with 320-grit diamond to reduce corner fractures.

Fracture stresses were determined in air with a special four-point bend test fixture [5] at a stressing rate of 6.9 MNm~2/sec. The specimens had an outer span of 3.17 cm

and an inner span of 1.9 cm. After the samples were fractured, optical and scanning electron microscopy were used to identify the fracture origin.

To determine whether the four-point bend strengths yield an average three-point bend

strength of 400 MNm" 2

, the procedure of Davies [6] was used. This calculation takes into

account the Weibull modulus, stress distribution, and specimen geometry. These data were calculated with the Weibull moduli presented in Table 1 and are given in parentheses.

The surface roughness due to grinding was characterized using a Talysurf IV profilometer with a 2.5-micron-diameter diamond stylus drawn perpendicular to the grinding

direction. The center-line average of the roughness profile was determined for each of the

grinding/ sintering combinations.

sstarlite Industries, Inc.

^Norton SD200-R75B56-1/8.

249
Table 1. Bend Strengths of A^Og

Average

No. of Strength,

Observations (MNm~2) Std. Dev. m

ST 8 218

SL 8 414

BT 9 359

BL 9 361

[276] 24 7.4

[493] 32 12.2

[416] 26 15.9

[457] 48 7.1

4. Results and Discussion

4.1. Grinding of Bisqued Alumina

The strength data obtained for the bisqued alumina bars are plotted in Figure 3 and

additional information is given in Table 1. These data indicate that the grinding direction has no apoarent effect on the strength of the bisqued, ground alumina. A comparison

of the standard deviations with the statistical F-test at the 95 percent confidence level

demonstrates that there is a significant difference in the variability of the data. This

could be explained by assuming that flaws are generated during grinding which are oriented

normal to the tensile direction. As a consequence, these flaws have a more common character and/or orientation than those generated by grinding in the longitudinal direction. This, of course, would explain the higher Weibull modulus of the BT samples. However, fractographs of representative samples do not substantiate this claim. Instead, BT

and BL samples fracture because of apparent intrinsic flaws (Figure 4) such as porous

regions (50 to 100 ym) or large grain areas. The porous regions probably arise from

agglomerates in the original powder which were not compacted during hydropressing.

/

/

/

/

/

Figure 3. Stress-failure probability curves for test bars ground transversely in sintered state (ST), longitudinally in sintered

state (SL), transversely in bisque state (BT), and longitudinally

in bisque state (BL)

.

250
65X 325X

Figure 4. Scanning electron micrograph of internal

pore fracture site for a BT test bar.

4.2. Grinding of Sintered Alumina

The average strengths and standard deviations obtained from SL and ST samples are given in Table 1 and plotted in Figure 3. It is apparent that samples ground transversely

have a significantly lower strength than those ground longitudinally. This would be the

expected result if one assumes that microcracking is associated with diamond grinding of

dense alumina. Then for transverse grinding, the potentially critical flaw would be

oriented normal to the tensile axis, and would essentially be larger [7] than that produced

when the microcrack is aligned parallel to the tensile axis. In fractographs of ST samples

(Figure 5), a small region of transgranular fracture near the surface is seen to depths

of 20 to 30 ym and lengths of 0.7 to 1.0 mm and the remaining fracture is intergranular.

50X 250X

Figure 5. Scanning electron micrograph of transgranular

fracture region which originated during machining

and acted as the fracture origin site (ST sample).

251
The existence of this type of fracture surface with the ST samples demonstrates that microcracking is the predominant flaw type for these samples. The particular significance of

the orientation effect is demonstrated by the fact that the fracture origins for SL samples were intrinsic in nature, i.e., pores and large grains (Figure 6). While in this

particular case the mean strengths were sufficient to indicate a difference due to grinding

one would also expect a significant difference to be demonstrated by comparing the Weibull

moduli (i.e., standard deviations). In Figure 3 the data for the ST samples indicate that

there may really be two different flaw populations instead of one. But, because of the

small population size, it is not possible to discern these differences with any degree of

rel iabil ity.

50X 500X

Figure 6. Scanning electron micrograph of internal pore

fracture site for an SL test bar

When the sintered ground (SL) strength is compared with the bisque ground strengths

(BT and BL), it becomes apparent that the SL process results in the least effect on the

strength of the alumina. Because all test bars were fabricated at the same time and failed

by the same mechanism, this would indicate that the BT and BL samples have been flawed to some degree during grinding. By comparing the surface roughness generated during the

grinding operation (Figure 7) it is apparent that the bisque ground surface both before and

after sintering is appreciably rougher than the sintered ground surface. This difference

is attributed to the mechanism of material removal which is chipping for the dense alumina

and a plowing or smearing process for the relatively soft, bisqued alumina. This surface

roughness can easily coincide with the surface pore to yield an effectively large flaw than

for a pore only. This increase in flaw size would account for the observed strength

differences noted above. This could also explain the directionality seen in the bisqued

alumina data.

One might conclude that the SL sintering/grinding operation is the best fabrication

process. But, it must be remembered that significant microcracking is associated with

material removal. The propagation of these cracks is strongly orientation dependent and

susceptible to subcritical crack growth in vivo . Therefore, for a tooth root, which is

used under multiaxial stresses rather than simple uniaxial stress, grinding of the tooth

root from dense alumina is less desirable than grinding from a bisqued alumina because of

the possibility for failure after implantation. While most of the stresses are compressive,

there is considerable stress generated at the serrations, which makes breakage a real

possibil ity.

252
B

i 0.25pm I

Figure 7. Surface stylus traces of surfaces machined with alumina

in sintered state (ST), bisque state, then sintered (BT)

and bisque state only (B).

From the above results, it is apparent that porosity is the main intrinsic flaw.

Therefore, to significantly improve the strength of the tooth roots, a milling operation is

needed to eliminate the agglomerates within the powder. With a significant reduction of

this type and size of flaw, and improved intrinsic strength, it might be expected that

grinding flaws would have a greater effect on the strength.

5. Conclusion

It has been shown that the main strength-limiting flaw for three of the test conditions

was related to intrinsic flaws (pores) rather than to the grinding operation. But, the

grinding operation does lead to some reduction in strength for the bisque ground material

and to a significant reduction in strength for the sintered ground material, depending on orientation. Thus, the strength of the ground alumina can be increased by powder processing to eliminate agglomerates, after which it is suggested that flaws associated with the

grinding operation could be the strength-limiting flaw type.

Assistance of Battelle staff members

respectively in diamond contour grinding,

Helpful discussions with G. Bansal and W.

J. Van Bibber, R. Higgins, and A. Skidmore

testing and fractography is greatly appreciated.

Duckworth contributed to this work.

253
References

[1] Driskell, T. D., O'Hara, M. J., and Niesz, D. E., "Surgical Tooth Implants, Combat

and Field", Report No. 2, Contract No. DADA 17-69-C-9181 (April, 1973).

[2] Hassler, C. R. and McCoy, L. G., "Surgical Tooth Implants, Combat and Field", Report

No. 7, Contract No. DADA 17-69-C-9181 (October, 1976).

[3] Rice, R. W. , "The Effect of Grinding Direction on the Strength of Ceramics", The

Science of Ceramic Machining and Surface Finishing , S. J. Schneider and R. W. Rice

(Eds.), NBS Special Publication 348, Washington, D. C: Government Printing Office,

pp 365-76 (1972).

[4] Tentative ASTM Standard F4-XXX, "High Density Alumina for Surgical Implants", American

Society of Testing Materials, Philadelphia, Pennsylvania (1976).

[5] Hoagland, R. G., Marschall, C. W., and Duckworth, W. H., "Reduction of Errors in

Ceramic Bend Tests", J. Am. Ceram. Soc, 59 [5-6] pp. 89-92 (1976).

[6] Davies, D. G., "Statistical Approach to Engineering Design in Ceramics", Proc. Br.

Ceram. Soc. No 22, pp 429-52 (1973).

[7] Bansal , G. K., "Effect of Flaw Shape on Strength of Ceramics", J. Am. Ceram. Soc,

59, [1-2], pp. 87-88 (1976).

254
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SIMULTANEOUS FINISHING OF BOTH SURFACES

OF HEMISPHERICAL DOMES1

R. Gentilman

Raytheon Research Division

' Waltham, MA 02154

F. Cooke

Frank Cooke, Inc.

North Brookfield, MA 01535

P. Miles

Raytheon Missile Systems Division

Bedford, MA 01730

The high cost of finishing durable oxide materials is the major factor

limiting their use as infrared domes for future missiles. A novel technique

has been developed for lapping and polishing both surfaces of hemispherical

domes simultaneously. The technique, which halves finishing time and

costs, has been successfully demonstrated for sapphire and glass.

Key Words: Dome finishing; finishing technique; hemispherical domes;

oxides, sapphire, two-surface lapping, two-surface polishing.

1. Introduction , More durable materials are being considered for domes for infrared air-to-air

missiles. In addition to optical transparency in the 3 to 5 micron wavelength band,

increased resistance to thermal shock, mechanical abrasion, rain erosion, and handling

damage is desired. Magnesium oxide, sapphire, and spinel (MgA^O^) are among the

candidate materials being considered.

The actual domes are hemispherical sheets approximately 75 mm diameter and

3 mm thick. Because it acts as a lens in the optical system which guides the missile,

the dome must be finished to within specific tolerances of sphericity and concentricity

in addition to being optically polished.

The conventional procedure involves generating the two spherical surfaces by

grinding followed by lapping and polishing one surface at a time. Unfortunately, the cost

of finishing domes composed of these hard oxide materials is high. In fact, finishing

costs can be as much as twice the cost of raw materials and blank fabrication combined.

This is the major factor which may limit the use of more durable domes on future

systems.

Reported herein are preliminary results using a novel technique for grinding and

polishing both surfaces of hemispherical domes simultaneously. Finishing both surfaces

at once is expected to halve finishing time and costs, and simplify the optical test procedure, thus reducing the total product cost to a more acceptable value.

This work is supported by the Air Force Materials Laboratory, Wright -Patterson AFB, OH 45433 under Contract No. F33615-77-C-5007.

255
2„ Conventional Lapping and Polishing Techniques

After a hemispherical blank has been generated by cutting and grinding, the dome is

cemented on a fixture and the inside spherical surface is lapped and then polished with

successively finer abrasive slurries (Figure 1-A). A pitch or metal lap is used with

either alumina, silicon carbide, or diamond abrasive. When finishing the inside dome

surface, the bottom lap rotates while the mounted dome undergoes a reciprocating motion

during which it is free to rotate about its axis in response to varying frictional forces.

DOME

LAP

Figure 1. The conventional finishing of

domes requires separate operations for

the (a) inside and (b) outside surfaces.

INSIDE SURFACE OUTSIDE SURFACE

The accuracy of the surface is checked against a spherical reference test plate, and,

if necessary, changes are made in the lapping motion to correct for errors. The final

operations produce a polished surface which is required to eliminate light scattering at

the surface,.

After the inside surface has been finished, the dome is remounted on the bottom

spindle and the outside surface prepared by similar procedures (Figure 1-B). The

accuracy of the second surface is checked independently. Then, when the dome is fully

polished and unmounted, the concentricity of the two surfaces is checked.

Unlike optical glass or the common infrared materials such as germanium, zinc

selenide, or magnesium fluoride, the durable dome candidate materials require extended

machine time of several tens of hours for lapping and polishing, much of it with expensive diamond abrasives. Rework of domes which do not meet specifications also requires

correspondingly long times.

3. The Two-Surface Finishing Technique

The two-surface lapping and polishing technique for domes is based, in part, on the

highly effective technique first introduced near the turn of the century by C. E. Johansson

for the production of high precision gauge blocks, and used recently for the fabrication of

multiple, high precision laser window elements. In this process, five or more blanks are

mechanized to execute a planetary movement between two counter-rotating annular lapping

surfaces [ 1 ]. The proper choice of speed of the planetary motion and of the absolute and

relative lapping speeds produces twin polished surfaces that are flat and parallel to a high

degree of precision.

Two-surface finishing of a hemispherical dome is shown schematically in Figure 2„

The dome workpiece is placed onto the inside lap, and the outer lap then placed atop the

dome. An oscillating support ring causes the dome to rock back and forth between the two

counter rotating spherical laps. In order to avoid contact with either of the laps, the

thickness of the support ring is less than that of the dome. The bottom lap is on a conven256
tional spindle while the top lap is driven by means of an off-axis ball-and-socket joint.

The top drive axis can be fixed at an arbitrary angle or made to swing through a range of

angles relative to the axis of the bottom spindle.

Figure 2. Schematic description of the simultaneous two- surface dome lapping and polishing

machine

= ±30°)

Figure 3 -A shows the inner lap and the support ring removed from the machine. In

Figure 3-B, the dome blank is shown being placed onto the inner lap. With the top lap in

place (Figure 3-C), the two-surface finishing process can commence.

Because both the blank and the laps are mechanized to execute four different movements, the dome is carried over a large fraction of each lapping surface and through a wide combination of positions of the lapping surfaces. These movements ensure a symmetry in the optical working of the dome as a function of dome azimuth and, with suitable

adjustment of relative movements,uniformity of material removed as a function of dome

angle.

The speeds of the four machine movements - top and bottom lap rotation, support

ring oscillation, and optional top drive shaft swing - are controlled independently with

variable speed motors. The relative rotation rates of the two laps are adjusted so that

there is little or no net rotation of the dome. In this way, lapping or polishing occurs

simultaneously on both surfaces.

4. Results and Discussion

Domes of glass and sapphire have been successfully lapped and polished using the

two-surface technique. As discussed below, optical quality of these domes is equivalent

to others finished by the conventional single surface technique. It should be noteds however, that finishing of sapphire domes by either single surface or two-surface techniques

has presented problems resulting not only from the hardness of the material but also

from its anisotropic machining characteristics, as described by Becher [ 2 ]

Diamond impregnated metal laps have been used for the initial lapping operations.

Abrasive slurries are used for polishing. Lead foil covered pitch laps have given the

best results for polishing the spherical surfaces.

257
A) Inside lap and support ring B) Dome blank being placed on

inside lap

Figure 3. Photographs of Two-Surface

Dome Polisher.

4

C) Polisher loaded. Free floating

outside lap atop dome.

For domes, which are transmissive elements, the control of optical thickness over

the dome is more important than the control of the optical figure of each individual surface. In conventional dome fabrication, tolerances on the local sphericity of each

surface are held to very tight tolerances to attempt to ensure acceptable control of transmitted wavefronts. The simultaneous two-surface technique, because of continual free

access to the dome, allows the dome figure to be observed in transmission, as well as

the individual surfaces to be spot checked. The degree of perfection of the transmitted

wavefront can be observed in a conventional interferometer using a correction lens to

compensate for the inherent power and spherical aberration of the dome., Thus the optical quality of the dome can be determined directly from the transmission interferogram

at any time in the polishing sequence.

258
To date, two sapphire domes have been fabricated using the two-surface polishing

technique, with encouraging results. Both were cut from single crystal boules such that

the dome axis lies close to the crystal c-axis. The design thickness of the domes was

2. 540 mm. Actual measurements of the better of the two gave a value of 2. 527 mm at

the zenith as compared to 2.500 ± 0.03 mm around the lip of the dome.

The optical quality of each dome was examined in transmission by placing it, along

with the precision correction lens, in a Zygo interferometer set up in the Fabry Perot

mode with plates adjusted to n§ar exact parallelism. The resultant interferogram for

the better dome,taken at 6328 A, is reproduced in Figure 4-A. The fringe pattern

results from a double pass through the dome and shows some 10 fringes of irregularity,

corresponding to a single pass optical path difference of some 3 nm over the 5 cm aperture of the system. While by no means perfect, this degree of uniformity is very close

to that achieved in similar c-axis domes polished by conventional techniques, where the

best result of several attempts produced a dome with a zenith thickness of 2. 522 mm and

lip thicknesses of 2.510 ± 0.005 mm. The transmission fringe pattern of this dome is

shown in Figure 4-B„

(a) (b)

Figure 4. Two-pass interferograms of the best domes produced by

(a) the two-surface technique and (b) conventional single

surface techniques.

In a related set of experiments, the spreading of a line image was measured in an

optical system capable of < 25 juradian image spread in the absence of the dome-lens

pair. With the dome-lens pair placed in an otherwise parallel light beam, the observed

focal line spreads were 75 /urads for the best conventionally polished dome and 88 /urads

for the best two-surface polished dome. It is clear that further improvement is needed

in the technique, but already domes can be produced that give image spreads of less than

100 /urads, a value approaching aperture limited resolution for infrared radiation beyond

3 /Ltm wavelength.

259
5. Summary

A new technique has been demonstrated for lapping and polishing both sides of

hemispherical domes simultaneously. The optical quality of domes produced to date

approaches aperture limited resolution at mid-infrared wavelengths. This technique

gives free access to the workpiece permitting optical evaluation in transmission at any

stage of the finishing process. When perfected, the technique will halve finishing time

and cost of domes made of durable oxide materials.

References

[1 ] Cooke, Frank, "About Flats and Parallels," Applied Optics 5_(7) 1104 (1966).

[2 ] Becher, Paul F. , "Abrasive Surface Deformation of Sapphire, " J. Am. Ceram.

Soc. 59 (3-4) 143 (1976).

Discussion

DECKER

Do you continuously add diamond abrasive during your final polishing

operation?

GENTILMAN

At the present time, the abrasive slurry is applied manually „ Lead

foil covered laps appear to extend the useful life of the diamond, but it

must be recharged periodically.

REMPER

You mentioned extensive time for polishing, can you be specific?

GENTILMAN

I do not have the time on the two-side technique, but it requires fort

hours per side with conventional single side polishing.

SCHMID

For zero degree sapphire domes, (domes whose axis is aligned with the

C axis) you state that the area on the top surface of dome polishes easier

than near the sides. You suggest that this is due to easy polishing of the

top zero degree material as opposed to the sides that is about 60° from

zero

.

I suggest that this is a mechanical problem. We have found that zero

degree material is hardest to polish. The dome is probably not polishing

well on the sides because of the polishing set-up. The pitch has give and

therefore there is insufficient pressure on the sides of the dome for adequate polishing.

260
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

ABRASIONLESS MACHINING METHODS FOR CERAMICS

R. F. FIRESTONE

I IT Research Institute

Chicago, Illinois 60616

Abrasionless machining methods for ceramic materials are reviewed with

extensive references to the recent literature. These methods offer significant

advantages over traditional abrasive machining methods since they are not

limited by the hardness of the ceramic and easily machine delicate or complex

parts.

Keywords: Ceramics; chemical milling, electric discharge machining; electron

beam machining; glass; hydrodynamic machining; ion beam maching; water jet

cutting.

1 . Introduction

Ceramics are the hardnest and potentially the strongest of all materials. Many applications require them to be machined to complex shapes with precision, turbine blades for

example. The hardness of ceramics makes traditional abrasive machining slow and expensive,

especially for complex shapes. Abrasionless machining methods circumvent the limitations of

traditional abrasive machining.

. 2. Background

Ceramic machining is traditionally accomplished by mechanical abrasion, an art whose

methods have been developed empirically and whose origins disappear into the mists of prehistory. Six thousand years ago the Egyptians were cutting blocks of granite for obelisks

using copper saws set with emery teeth, dressing the surface with dolerite balls and

polishing with fine emery power [1]. In these processes, one can recognize the precursors of

the modern diamond saw, grinding wheel and polishing lap. Advances on these ancient techniques have been made by improving the abrasives, e.g., diamond instead of natural emery, and

substituting power machinery for the sweating fellaheen. But the basic processes for shaping

hard materials by abrasion remain the same.

Abrasion is the most widely used and highly developed machining method for ceramics.

Hard abrasive grains are pressed against the workpiece to introduce cracks and remove stock

by controlled fracture. The cracks can shatter fragile parts, reduce the strength and provide nexus for delayed failure. The rate of stock removal depends upon the ratio of the

hardness of the abrasive to the hardness of the ceramics. Machining hard ceramics such as alumina, silicon carbide, silicon nitride, etc. is slow and costly even with diamond, the

hardest abrasive. The cost of machining can compose from 50% to 90% of the cost of the

finished part [2]

Abrasive machining methods, being mechanical, are generally restricted to simple parts

that are flat, or have axial or radial symmetry. Machining complex or asymmetrical parts may

be nearly impossible and phohibitively expensive. Thus, ceramics which are not formed to

final shape tend to have simple geometrical shapes, a severe limitation on their use.

Numbers in square brackets indicate the literature references at the end of the paper.

261
FUNDAMENTAL MACHINING

ENERGIES

STOCK

REMOVAL

Mechanical

Abrasion

Thermo-Electric

Vaporization

TRANSFER MEDIA

ENERGY

SOURCE

Physical

Contact Abrasive

Tool

High Velocity

Particles

Erosion

High Velocity

Liquid

Photons

Pneumatic/

Hydraulic

Pressure

Electrons High Voltage

Amplified

Light

Sputtering Ions Ionized

Material

MATERIAL REMOVAL

PROCESS

Traditional

Machining

Abrasive

Jet

Ultrasonics

Hydrodynamics

Chemical Ablative

Reaction

Reactive

Environment

Corrosive

Agent

Chemical

Milling

Electric

Discharge

Machining

Electron

Beam

Laser Beam

Ion Beam

Figure 1 . Ceramic machining processes

The objective of this paper is to survey the most important abrasion"! ess machining

methods for ceramics. The principles of each method will be discussed and the advantages

and disadvantages compared with abrasive machining will be pointed out. The literature,

mainly since 1970, will be reviewed. References to earlier literature can be found the

proceedings of the first ceramic machining symposium [3],

3. Abrasionless Machining Methods

Abrasion! ess machining methods do not have the limitations of abrasive methods which

require solid-to-solid contact. The hardness of the ceramic is not a factor in determining

stock removal rate. Some of the abrasionless methods produce ceramic parts without cracks

or flaws which are therefore stronger than as-fired or abrasive machined parts; allow

asymmetrical, three dimensional shapes to be produced to close tolerances; or have cubic

centimeter per second stock removal rates. Because of the advantages, abrasionless methods

are finding increasing application to ceramics.

Machining is the application of energy to the workpiece to remove stock. The fundamental sources of machining energy can be separated into: mechanical, chemical and thermoelectric. The relation of these sources to the machining processes is shown schematically

in figure 1 after Springborn [4]. This diagram includes traditional abrasive machining, two

non-traditional abrasive methods, and the six major abrasionless machining methods.

Mechanical energy is the oldest energy source applied to machining. It is used for

both traditional abrasive machining; non-traditional abrasive machining: abrasive jet and

ultrasonics; and for abrasionless machining: hydrodynamic machining (HDM). Chemical energy

is used for chamical milling (CM), the oldest abrasionless machining method, and for its

hybrid with the thermo-el ectrical energy, electro-chemical machining (ECM). Thermo-electric

is the final category and includes four-recent abrasionless methods: electric discharge

machining (DM), electron beam machining (EBM), laser beam machining (LBM) and ion beam ma- chining (IBM). The applications of these methods to metals have been recently reviewed [5].

Machining methods may be described by the process by which the workpiece is shaped.

The generation process focuses energy onto an area of the workpiece and by relative movement

generates a new free surface of desired configuration. Hydrodynamic machining, electron

beam machining and laser beam machining use this process which may be numerically controlled.

The transfer process uses a pattern or model to direct energy from a diffuse source. Chemi262
I cal machining and electric discharge machining use this process which is analogous to

printing. Ion beam machining is unique since it can use either process.

The material properties and environment which are required for the use of abrasionless

methods are listed in table 1. Hydrodynamic machining and laser beam machining are the most

flexible methods since they can be used to machine any material under ambient conditions.

Hydrodynamic machining is the most portable since it was originally developed as a field

method. The other four methods require immersion of the workpiece in a liquid or a vacuum.

Table 1. Conditions for abrasionless machining methods

MACHINING METHOD MATERIAL ENVIRONMENT

Hydrodynamic (HDM) Any Ambient

Chemical (CM) Non-crystalline Reagent

Electric Discharge (EDM) El ectrical ly conductive9 Dielectric

Electron Beam (EBM) Any Vacuum

Laser Beam (LBM) Any Ambient

Ion Beam (IBM) Any Vacuum

a Resistivity less than 300 ohm-cm.

The applications and results of abrasionless machining methods in the present state of

development are listed in table 2. Hydrodynamic machining (HDM) is best for rapid cutting

of large holes, slots and grooves in porous materials. Chemical milling (CM) is good for

transferring patterns of shallow holes and slots. Electric discharge machining (EDM) drills

small holes well but is unexcelled for pocketing, i.e., making cavities with straight sides

and flat bottoms, and contouring surfaces. Electron beam machining (EBM) and laser beam

machining (LBM) are excellent for drilling small holes. Ion beam machining (IBM) can be

used either as a focused energy source for drilling and contouring or as a diffuse source

for pattern transfer. Ceramics that have been ion beam machined are usually stronger and never weaker than as-fired or abrasive machined parts. Chemical machining and laser beam

machining may also increase the strength of ceramics.

All abrasionless machining processes, with the exception of chemical milling, are

capital intensive compared to abrasive machining. Minimum cost of equipment ranges from

approximately $10,000 for electric discharge machining up to $1,000,000 for high powered

lasers. Specially trained personnel are required and operating costs are usually higher.

However, for many applications, the total cost of machining a ceramic part is less than

abrasive machining.

3.1 Hydrodynamic Machining

Hydrodynamic machining uses a high velocity fluid jet to erode material [6]. The fluid

is commonly plain water and the method is often called water jet machining. There are three

types of jets: pulsed jets which are used to fracture material, cavitating jets which are

used primarily underwater, and continuous jets which are the subject of this section.

The water jet is a nonwearing tool with no frictional drag in the cut and virtually no forces at right angles to the jet. The kerf is insensitive to dwell of the jet which allows

contours to be accurately controlled. It is as versatile as a jigsaw on wood and is, indeed,

used to cut wood and other organic materials.

The equipment for a continuous jet is simple and robust (fig. 2). It consists of a fluid reservoir, a hydraulic intensifier to produce high pressure, a pump to drive the intensifier surges, pressure and flow control systems, and a nozzle. A relatively low flow of

fluid is required for cutting, only 1 to 10 1/m which is less in fact than the normal coolant flow for diamond grinding.

263
Table 2. Comparison of abrasionl ess machining methods for ceramics

METHOD

"

HDM

a CM EDM EMB LBM IBM

APPLICATION

Hoi es

:

Standard (>.005") Good

,

Poor Fair Poor Fair Poor

Small (.005-. 001") N .A. Good Fair Good Good Fair

Micro (< .001") N.A. N.A. N.A. Fair Good Good

Shal lnw (\ /D < ?C\) C

OIICil IUW \L/ U ^ ^- ^ / uUUU uuuu Fa "i v Fa i Fa "i rfl I i ro. \ i Deep (L/D >20) Fiar N.A. Fa i r Poor Poor N.A.

Through Cutting:

Shallow (L/D >20) Good Good Poor Good Good Fair

Deep (L/D< 20) Fair N.A. N.A. Fair Fair N.A.

Contouring Poor N.A. Fair N.A. N.A. Good

Pocketing Fair Poor Good N.A. N.A. Fair

Through Cavities Good Good Good Poor Poor Good

RESULTS:

Maximum Precision (in) .005 .001 .0005 .001 .001 .000001

Best Finish (y-inch rms) 100-1 000 5-1000 10-500 20-100 20-50 1-20

Effect on Strength D

d

M

6

D D M I f

See text for meaning of acronyms.

^Not applicable.

Length to diameter ratio.

^Decrease in strength of abrasive machined ceramic.

May either increase or decrease strength of abrasive machined,

^Increase in strength of abrasive machined ceramic.

Figure 2. Hydrodynamic machining (HDM) apparatus

264
The process parameters are pressure (50 to 1000 MPa), nozzle diameter (0.01 to 1 mm),

standoff distance between nozzle and workpiece (2 to 25 mm), jet angle (usually normal incidence, 90°) and traverse rate (0.1 to 10 m/s). The power requirements are related to the

pressure and nozzle diameter. The cutting rate is a complex function of the process

parameters and the workpiece properties. In general, it increases with increasing pressure

and decreasing nozzle diameter and standoff distance. The hardness of the workpiece is not

a major factor but weak, porous materials cut more rapidly than dense, strong ones. The jet

may be augmented with abrasives to increase the cutting rate in some materials. The abrasive may be either inspirated by the nozzle or placed on the surface of the workpiece to be

driven in by the jet.

The depth of cut in ceramic materials varies from 0.1 to 100 mm for a single pass of a simple jet using plain water. Multiple passes of the jet in the same kerf will increase the

depth of cut by a factor of two approximately. The depth produced by multiple passes is

limited since the standoff distance increases with each pass. By using a penetrating nozzle

which descends into the kerf to maintain constant standoff distance, deep holes may be

drilled and multiple passes may be used to cut grooves and slots. The depth to kerf ratio

may be 150 to 1 or greater and the depth may be 3000 mm or more being limited only by the

length of the nozzle.

The kerf is a function of nozzle diameter and material grain size. For small nozzle

diameters and fine grain material, the kerf may be only 0.5 mm greater than the nozzle diameter. For coarse grain materials, it may be as large as the nozzle diameter plus two grain

sizes. Since the jet spreads as it leaves the nozzle, the kerf is slightly larger at the

bottom than at the top. This may be reduced by using water to which has been added long

chain polymers, such as polyethylene oxide. Penetrating nozzles cut wider kerfs to accomodate the nozzle as it descends into the cut. The minimum kerf for a penetrating nozzle is

dependent on the depth of cut but can be less than 5 mm.

Water jets have been used to cut metal, rock [8], composites, concrete [9], and many

softer materials such as wood, cloth and paper [10]. Their use on ceramics is just beginning. Some of the preliminary results on the use of a water jet to cut refractory bricks by

the author are shown in table 3. The values are averages for each class of brick. The

surface exposure rate (m^/s) is the depth of cut multiplied by the traversing rate.

Table 3. Water jet cutting of refractory bricks

Refractory Depth of Cut (mm)

Insulating Firebricks 90

Porous Bricks, Al

2 0

3 & Si C 25

Dense Firebricks 15

Fusion Cast Bricks 1

CONDITIONS: Nozzle Diameter: 0.5 mm Pressure: 480 MPa

Traverse Rate: 1 m/s

Fluid: Plain water

Jet Incidence: Normal

Standoff Distance: 10 mm

Surface Exposure Rate (m /s)

90 x 10" 3

25 x 10" 3

15 x 10' 3

1 x 10" 3

The results indicate that the porosity and tensile strength of a material is more important than the hardness in determining the cutting rate. The alumina and silicon carbide

bricks are very hard and abrasion resistant but quite porous. Rough calculations show that

for the same power input, the water jet cuts these bricks five times more rapidly than a diamond saw. The silicon carbide bricks will abrade the firebricks but the firebricks are denser and stronger and so do not cut as rapidly. Fusion cast bricks are very strong and

non-porous. The cutting rate is very slow. Increasing the pressure shatters the brick

265
rather than increase the cutting rate. Future work may result in a more exact correlation

of these parameters.

The effect of hydrodynamic machining on strength is similar to abrasive machining.

However, since it works best on porous materials which are inherently weak and in which the

pores tend to serve as crack stoppers, the strength loss is minimal.

3.2 Chemical milling

Chemical milling is the controlled solution of the workpiece by a chemical reagent.

The areas to be milled are defined by a chemical resistant, strippable mask. Since solution

proceeds radially from the opening in the mask, the resulting undercutting limits the pro- cess to shallow cuts. Undercutting control on deeper cuts or contour milling may be accomplished by successively stripping masks as the cutting progresses. The amount of material

removed, i.e., the depth of cut, is controlled by the length of time the workpiece is left

in the reagent and the temperature.

Although there are reagents which will dissolve almost all ceramics, chemical milling

is restricted to glass by the lack of suitable reagents which will dissolve polycrystall ine

ceramics uniformly and at low temperatures where masking can be used economically. For

chemical milling glass (glass etching), the reagents are based on hydrofluoric acid and the

mask is traditionally wax through which the pattern is scratched, although polymers and

photoresists are now also used. For an abrasionless machining method, glass etching is very

old, beginning shortly after K.S. Scheele prepared hydrofluoric acid in 1771, and rapidly

supplanting copper wheel engraving of glass 11 . Much of the technique was developed

empirically but there has been some scientific study [12].

A special proprietary photosensitive glass has been developed by Corning Glass Works

which is self-masking [13]. When exposed to ultraviolet light, the glass records a latent

image which is developed by heating. There is a pronounced difference in solubility between

the exposed and unexposed areas so that the image in the glass can be etched out with hydrofluoric acid. Any pattern may be produced with resolution of 25 um. Further heat treatment

after etching can change the glass into a crystalline glass-ceramic.

A variation on chemical milling which can only be applied to conductive ceramics, such

as carbides, is electro-chemical machining (ECM). The electric field supplies the additional energy which allows polycrystall ine ceramics to be dissolved uniformly at room temperature. There is an extensive literature which was reviewed in 1974 [14]. Recent work

on the process for carbides has been done in Japan [15] and Russia [16]. A detailed description of the process has been given by Wilson [17].

Chemical milling need not decrease the strength of ceramics. Glass etching, if done

under proper conditions, can eliminate surface flaws and restore glass to its pristine

strength [18]

.

3.3 Electric discharge machining (EDM)

Electric discharge machining is the controlled melting and vaporization of the workpiece by multiple arcs between the workpiece and tooling to form a three-dimensional reverse

image of the tooling. The components required are: a pulse generator to deliver either DC

bipolar pulses 200 to 500,000 times a second at 50 to 100 volts; a tool, usually graphite,

copper or tungsten; a dielectric, usually oil or deionized water; a pump to circulate the

dielectric; and a feed control to maintain a constant standoff, usually 10-500 um, between

the tool and workpiece (fig. 3). The workpiece must have an electrical resistivity of less

than 300 ohm-cm unless a conductive electrolyte is used in place of a dielectric.

An excellent review of electric discharge machining up to 1968 was made by Bucklow and

Cole [19], who point out that it has developed broad applications "by virtue of its special

abilities to sink any profile, however elaborate, into a workpiece in one basic operation,

to machine any electrically conductive substance irrespective of hardness, and to remove

metal without exerting any appreciable force on the workpiece." There are no other articles

266
Figure 3. Electric Discharge Machining (EDM) Apparatus. PG, pulse generator; SCF, servo controlled feed; E, electrode; D, dielectric; WP, workpiece; WT, work table; F, filter;

P, dielectric pump.

which: treat the physics of the process [20]; compare electric discharge machining with

electrochemical machining [21,22] and with laser beam machining [23]; and analyze the cracks

produced in cemented tungsten carbides which propagate through the cobalt phase [23,24].

Since electric discharge machining causes melting of the material, there is a thermally

affected zone (recast layer) beneath the machined surface which produces residual stress and

cracks. In fact, this method is often used to introduce automically sharp cracks into

fatigue specimens. Electric discharge machining decreases the strength of cemented carbides.

In applications, such as tool inserts, where high strength is required, the material is

diamond ground rather than electric discharge machined. For hot pressed silicon carbide,

the author's experiments on flexure bars show that the room temperature strength is reduced

by half after electric discharge machining 100 ym from the tensile side. However, the

machining conditions were not optimized.

3.4 Electron Beam Machining

A beam of electrons is focused to produce an energy density of 10^ to 10^ watts per

square millimeter on a part of the workpiece which causes it to melt and evaporate. The

components required are: electron gun, a high voltage supply (50 to 150 kV), a vaccuum

chamber (10~5 to 10-6 Torr) and beam deflecting coils (fig. 4). The beam positioning is

very fast since it is electrical and can be computer controlled. For example, holes may be

drilled in ceramic substrates at the rate of 12,000 per second and the beam may be traversed

at 10 m/s [25]

Workpiece Viewing

System

Figure 4. Electron Beam Machining (EBM) Apparatus. C, cathode; A, anode; V, valve; ML

magnetic lens; DC, deflection coils.

267
Electron beam machining is particularly good for drilling small, deep holes. Hundreds

of holes 0.1 to 1 millimeter in diameter can be generated in seconds. The process is not

affected by the kind of material being drilled or its hardness [26]. Not only cylindrical

holes can be drilled but also tapered holes and holes with irregular cross-section. The

holes may be drilled at normal incidence or at an angle to the workpiece surface. For

drilling very small holes in ceramics, beam focusing must compensate for the space charge on the surface and above the ceramic [25,27].

Since electron beam machining is essentially a thermal process, there is a melted zone around the cut and residual stresses and cracks are introduced into the ceramic. While thin

(.040 in. thick) substrates can be cut by performation, thicker ceramic pieces must be

heated before machining to avoid shattering. Another approach is to machine unfired ceramics,

the binder volatiles which removes excess heat and firing heals any cracks [28].

3.5 Laser Beam Machining (LBM)

Laser beam maching uses a laser to form an intense beam of light which focused on the

workpiece causing it to melt and vaporize due to the intense heating [29]. The components

required are a laser with controls, a means of focusing the beam and workpiece positioner

(fig. 5). Beam positioning is mechanical so it is neither as fast nor as precise as electron beam machining, a process with which it directly competes. On the other hand,

machining is done under ambient conditions which gives the process greater flexibility.

Laser are of two types: continuous (C.W.) and pulsed, and the pulsed are further

divided into normal and Q-switched. The pulse duration of these types is tabulated below.

Laser Type Pulse Duration

Continuous (C.W.)

Normal Pulsed

Q-switched

10"^ seconds to infinity

10"^ seconds

_g

10 seconds

The wavelength of the light emitted depends upon the lasing material of which Nd-glass at

1.06 ym and CO2 at 10.6 ym are, perhaps, the most common.

Laser beam machining is not limited by the hardness of the material. For example,

graphite and diamond are cut with almost equal rapidity and, indeed, lasers are used for

commercial production of holes in diamond wire drawing dies [30-33]. Lasers may be used to

drill, groove and cut ceramics and are nearly as versatile as an end mill on metal.

The rate of laser beam machining is controlled by the rate at which material is removed

which is melted and vaporized by the beam. This occurs mainly by thermal convection but

also by the radiation pressure of the beam [29]. Experiments with metals have shown that

the machining rate is greatly increased if a gas jet is used to blow the melted material

away [34]

.

Figure 5.

WorMeca WrfVA Work Table

| [ | | Laser beam maching (IBM) Apparatus

268
The thermal gradients produced in the material around the focus of the laser beam produce stresses which can caused cracking and fracture [23,35-38]. Laser damage is a complex

phenomenon which is not well understood. The amount of damage is a function of the laser

wavelength and power, the duration of the beam, and the material properties, particularly

absorption of the laser light. Combining pulse and continuous laser beams produces

enchanced damage in some materials [39]. In spite of the thermal stresses, laser beam

machined ceramics can be as strong or stronger than diamond machined ceramics with the proper

choice of machining conditions. For example, alumina substrates sawed with a Q-switched

Nd-YAG laser (1.06 m) were 45% stronger than substrates sawed with a continuous C0

2 laser

(10.6 m) and 13% stronger than diamond sawed substrates [40].

Damage may be avoided by machining unfired ceramics, by using very short pulses to re- move only small amounts of material so the thermal stresses produce only elastic deformation,

or by putting the ceramic in compression to suppress cracking. This is done with diamond

dies which are mounted in a metal matrix if holes larger than 0.12 mm are to be drilled.

Preheating the workpiece before laser beam machining also reduces thermal stresses by

decreasing the temperature gradient and making some materials more ductile. This approach

has been used to cut sheet glass successfully. A 1 5 kw CO2 laser made smooth, damage free

cuts through 1.6 mm thick soda lime glass at 250 mm/s and through 3.2 mm thick glass at

75 mm/s after the glass had been preheated to 593°C [41]. At this temperature, the viscosity of soda-lime sheet glass is about 10^ poises so that thermal stresses may have been

accomodated by local plastic flow. Efforts by the author to use this method for laser cutting of hot pressed silicon carbide and silicon nitride were not successful. Even when

preheated to 1000°C, specimens, 3x6x50 mm, shattered when cut with a 500 watt continuous CO2

laser beam. Of course, these ceramics are still very strong at this temperature.

Laser beam machining is established in the semiconductor industry where lasers are routinely used for machining thin substrates [42-44], for drilling small holes with precise

dimensions [29,45,46] and for scribing grooves in substrates and silicon for controlled

breakage [47-49]. For example, a 150 watt CO2 pulsed laser can scribe a 0.63 mm thick

alumina substrate at 230 mm/s, cut it through at 20 mm/s or drill a through hole in 10 ms

[50]. Silicon wafers can be scribed with less than Mo yield loss [51]. The lasers which

are available for semiconductor [52] and ceramic machining [53] have recently been reviewed.

Laser beam machining of larger parts has not been as successful since it is difficult to

avoid catastrophic fracture. This can be used as an advantage in removing large volumes of

material, as in concrete breaking [54], or in laser assisted machining in which a single

point tool is used to remove the laser weakened material [55]. Thermal shock resistant

materials such as silica glass [44,56] and porous refractories may easily be laser machined

without fracture. The refractories may also be grooved or dimpled with a laser without

melting the surface by sintering and densification [44].

A number of studies have been made of laser cutting of silicate glass [41,57-63] and

also of laser welding [64]. However, glass which has been laser cut without preheating is

weaker than conventionally cut glass. Hot glass may be cut without weakening it, this is

inconvenient except immediately after forming. Although laser cutting is fast, it is no faster than glass cutting and much more expensive than this well established process. Be- cause of these factors, laser glass cutting has found few commercial applications. Cemented

tungsten carbides have been drilled using lasers [65,66] . However, the drilling initates

cracks which propagate through the cobalt phase [67] as in electric discharge machined

carbides [23,24]

Laser beam machining is finding increasing application to ceramics. The use of lasers

on ceramics and other inorganic materials has been the subject of several reviews of the

process and its applications [68-71] and of the manufacturing technology of computer con- trolled laser cutting [72,73]. A bibliography through July 1977 on laser applications to

non-metals has been compiled [74].

3.6 Ion Beam Machining

Ion beam machining uses a directed stream of ions to sputter material atom-by-atom from

the surface of the workpiece [75-79] with minimum damage leaving a surface which is atomically clean [80,81]. Atoms are ejected from the surface by the transfer of momentum from

269
the incident ions when the transferred energy exceeds the binding energy (-10 eV) and the

momentum vector of the atom points away from the surface. Atoms whose momentum vector points

into the surface are implanted as are some ions of the incident beam. Sputtering has some- times been called enhanced sublimation but it is not a thermal process and there is very

little heating. For example, paper may be machined to nanometer thickness without burning or charring. This behavior of surfaces under ion bombardment has been recently reviewed [82].

A system for ion beam machining is composed of an ion gun, a neutral izer, a workpiece

positioner and vacuum chamber (fig. 6). The gun uses a confined plasma discharge to generate

ions which are then extracted and accelerated to form a beam of ions which are positive if

the usual inert gas is used for the plasma. The neutral izer sprays electrons onto the workpiece to eliminate space charge buildup. The worktable allows translation of the workpiece

in three dimensions, variation of the angle of incidence between the beam and the workpiece,

and workpiece rotation. Accelerating voltages vary between 500 and 10,000 eV and beam

current density is usually from 0.5 to 2 ma/cm. The vacuum chamber pressure during machining

is 1 0-4 Torr or 1 ess.

Vacuum Chamber

E

Workpiece

Work Table

11

Vacuum

Pump

Figure 6. Ion Beam Machining (IBM) Apparatus. E, accelerating electrodes; F, filament; M,

magnet; N, neutral izer.

There are two types of ion beam systems: diffuse which is shown in figure 6 and

focused. The focused system is similar to the electron beam apparatus shown in figure 4

except an ion gun is used instead of an electron gun and a neutral izer is added. The deflection coils are used to sweep the ion beam across the workpiece surface. The focused ion

beam is used for drilling, cutting and surface contouring by computer control.

The diffuse ion beam system uses multiple ion sources to form a uniform ion beam which

may be as large as 30 cm diameter [83]. The areas to be machined are defined by openings in

a mask which may either be placed on the workpiece as in chemical milling [84] or close to

the ion gun and its image projected on the workpiece (fig. 7). The latter method allows

multiple images to be projected [85] at up to fifty times demagnification [86] so that very

small patterns may be machined.

The ion beam machining rate is determined by the sputtering yield which is defined as the number of sputtered workpiece atoms per incident ion. The sputtering yield depends on the material being machined and its temperature, the ions used and their energy and angle of

incidence, the ion beam current density, and, in some cases, the residual gas in the vacuum

chamber. The machining rate increases with increasing ion energy and usually saturates above

a few kilo-electron volts as penetration of ions into the workpiece increases. Ion penetration is also affected by the angle of incidence since normally incident ions penetrate more

270
ION APERTURE LENS TECHNIQUE

©

ION SOURCE

MASK

APERTURE

LENS

TARGET

Figure 7. Ion beam projection lithography

deeply into the workpiece and produce fewer workpiece atoms with momentum vectors pointing

away from the surface than grazing incidence ions. For most materials, the maximum machining rate occurs at an angle of incidence of about 15° from the workpiece surface. Since the

rate is related to the binding energy of the workpiece atoms, reactive gases can change the

rate. For example, oxygen absorbed on machined silicon surfaces froms silica and reduces

the rate while fluorine forms volatile silicon compounds and increases the rate.

A compilation of ion beam machining rates was published in 1971 [87]. Data from more

recent literature and from the author's work are in table 4. The rates are not affected by

the hardness of the material, hot pressed silicon nitride and soda lime glass machine at

about the same rate, but are dependent on process parameters. Ion beam machining is very

slow, the slowest of all ceramic machining processes, but in spite of this it is being used

in a number of applications where delicate parts are involved and/or it is necessary to produce a clean, damage free surface of precise dimensions. Ion guns are now being developed

which produce intense ion beams with current densities a thousand times or more greater than

present systems [89] and which can be used to produce concomitantly greater stock removal

rates

Surface textures may develop during ion beam machining [90] which are similar in

appearance to wind erosion features on rock [91] and show characteristic topography such as cones, facets, ridges, pyramids, hollows or even hoodoos. These textures are particularly

pronounced on polycrystal 1 ine and/or polyphase surfaces but also occur on single crystal and

normally clean glass surfaces. Ultra clean single phase glass and most crystalline semiconductors, whose surfaces become amorphous during machining with inert gas ion [92], are the

only materials which do not develop texture.

The development of texture is caused by three mechanisms: the sputter yield is different on different crystalline faces [93]: and it depends on the angle of incidence [94]; and some of the sputtered material may redeposit on the higher parts of the surface. Materials

which are initially smooth and featureless will develop a stable texture during machining

whose specific topography depends on the interaction between the mechanisms. The initial

workpiece contour [96] and its angle of incidence to the ion beam also influence topography.

The relief of even the most strongly textured surfaces is only a few microns and may be

ignored for many applications. It may be minimized by rotating and rocking the workpiece

during machining to vary the angle of incidence and by starting with an ultra clean workpiece surface. However, texture development is not always a disadvantage. Some strongly

textured surfaces show high light absorptance and have been suggested for solar absorbers [97].

271
Table 4. Ion beam machining rates

MACHINING ION BEAM PARAMETERS*

3

MATERIAL RATE (ym/hr) a ENERGY (eV) CURRENT (mA/cm2

) ANGLE REFERENCE

Polycrystal

s

Al 1 .8-4.2 500 90° [84]

4 .

2

1 000 0.5 ICO

1 5 2.7-4.5 1000 90° [84]

AIAT 0

2

U

3 0 5 500 90° T841 Lot j 0.8 1000 1 90° [84]

C (graphite) 3 1000 0.5 15°

C (diamond) 120 7500 1 90° [88]

LiNb0

3

3.8 1000 1 90° [84]

Si 1.3-3.0 500 1 90° [84]

2.2-4.5 1000 0.5 15° [84]

SiC (HP)

f 3.5 1000 0.5 15° —

Si C (HP) 5.6 1000 0.5 15° —

Single Crystals

A1

2 0

3 {lT02} 0.5 500 1 90° [74]

A1 9 0 Q {1123} 2.2 1000 0.5 15° _

Si {100} 1 .3 500 1 90° [74]

Si0

2 {100} 2 500 1 90° [74]

Glasses

Sil ica 3.9 1 uuu u . D 1 0

Soda-1 ime-sil ica 5.9 1000 0.5 15°

1 .2 500 1 .0 90° [74]

a for milling except where noted.

Argon ions.

c normal incidence.

^numbers in square brackets refer to the 1 iterature references at the end of the paper

.

When no reference is given, the data are the result of the author's research.

edrill ing.

^HP = hot pressed.

272
Ion beam machining can implant ions into the workpiece and causes radiation damage in

semiconductors and zirconia [98]. Most of these effects can be avoided by using low energy

ions at grazing incidence. If damage occurs, it may be eliminated after machining by annealing, often with a laser [99]. However, ion beam machining can also be used to produce

desirable changes in the properties of thin semiconductor films [100].

Precision ion beam machining of contours on optical components has been developed into

a process of great precision for glass polishing [101-109], figuring of aspherical surfaces

[110-116], and grooving and blazing diffraction gratings [76,117,118]. By 1970, computer

controlled ion beam machining could remove glass with depth tolerances of 1 nm from 750 mm diameter components and "superpol ish" the surface for increased reflection efficiency [119].

Today, even larger components may be ion beam machined.

The advantage of ion beam machining for optical components is the precision with which

stock can be removed to form an arbitrary surface and the freedom from the residual stress

and deformed layer which are introduced by mechanical abrasion and can caused delayed

distortion.

Diffraction gratings are now commercially produced by ion beam machining. A diffuse

beam is used and masking is used to define the areas to be machined. X-ray gratings are

produced at the National Physical Laboratory, London [116] and blazed holographic gratings

and grating couplers are produced at the Institute of Physical and Chemical Research, Japan

[117].

Ion beam machining is a rapidly expanding method for producing semiconductor devices,

especially since the ion beam can not only be used for machining but for thin film deposition

(SID, secondary ion deposition) [120-122]. This is also useful for materials such as sodium chloride which are hydroscopic since they may be machined and then coated with a moisture resistant film before being removed from the vacuum chamber to the ambient air [123].

The first ion beam machined semiconductor devices were made using a modified ion beam

thinner for TEM foils with a 1 cm diameter beam [124] but ion sources of 30 cm diameter are

now used [83]. The linewidths possible with ion beam machining have decreased from 10-4 m

[125] since the initial experiments to 10"6 m using carbon masks [126] which is much less

than is possible with chemical etching [127] to 10"' m using STEM lithography [128]. There

are numerous reviews of ion beam techniques for device fabrication [89,124,125] and of

production methods and equipment [83,129-132].

The strength of a ceramic which has been ion beam machined is greater than one which

has been conventionally abrasive machined and may be greater than the strength of the

pristine as-fired ceramic. This is due to the removal of surface flaws by ion beam machining which does not introduce new flaws so that the strength of the ceramic is controlled by

processing defects like pores and agglomerates. Ion beam polycrystall ine ceramics have been

found to be 1 0 to 60% stronger than diamond machined ceramics [133] and strength increases

of over 100% have been reported for silica glass after ion beam machining [134]. However,

the strength of the glass was limited by the texturing caused by ion beam machining since

the fracture origins were always at the bottom of depressions in the surface. The author's

ion beam machined a number of ceramics which had been carefully diamond machined. The

strength increases varied from about 50% for a processing immature silicon carbide to over

300% for a processing mature silicon nitride. In all cases, once the damaged surface produced by diamond machining was removed, the strength was controlled by processing flaws.

Because ion beam machining does not introduce damage, it has been used to shape a number

of delicate components. Diamond machined sawn-cavity heterostructure lasers rapidly degrade

in use but ion beam machined lasers do not [135]. Other components which have been ion beam

machined with success are: feedback lasers [136], reflective arrays [137], quartz resonators

[138], acoustic transducers [139], pyroelectric vidicons [140], magnetic arrays [141-142],

surface acoustic wave devices [143] and grooves, aerodynamic [142] and hydrodynamic [144],

for bearings. The literature on ion beam machining and applications is extensive. An incomplete bibliography contains 215 references for the years from 1954 to 1975 [145].

273
4. Conclusion

Abrasion"! ess machining methods for ceramics offer a variety of techniques which complement conventional abrasion methods. As the requirements for ceramic materials become more

stringent or the uses more exacting, the application of these unconventional method will

grow.

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[136] Schinke, D.P., Smith, R.G., Spencer, E.G., and Galvin, M.F., Thin-film distributedfeedback laser fabricated by ion milling, Appl. Phys. Lett. 21_ [10] (1 972).

137] Smith, H.I., Williamson, R.C. and Brogan, W.T., Ion beam etching of reflective array

filters IEEE Ultrasonics Symposium , J. Deklenk, ed., 1 98-201 (1 972).

[138] Vig, J., Wasshausen, H., Cook, C, Katz, M. and Hafner, E., Surface preparation and

characterization techniques for quartz resonators (Symp. on Freq. Control, 27th Ann.

Proc. Pap. Cherry Hill, N.J. 93-112 June 12-14, 1973).

139] Beecham, D., Sputter machining of ferroelectric potassium sodium niobate, Ceram. Bull.

52 [5] 454-455 (1973)

.

[140] Yamaka, E., Teranishi, A., Nakamura, K. and Nagashima, T. Pyroelectric vidicon with

grooved reting of PbTi03 ceramics Jpn. Ferroel ectrics TJ_ [1-2] (IEEE Symp. on Appl. of

Ferroelectric Pt. 2, Alburque, NM, 304-308 j 1 975)

.

[141] Schmidt, P.H. Spencer, E.G. and Wlaters, E.M., Ion milling of magnetic oxide platelets

for the removal of surface and nearsurface imperfections and defects, J. Appl. Phys.

41_ [11 ] 4740-4742 (1970).

142] Spencer, E.G., Schmidt, P.H. and Fischer, R.F., Microstructure arrays produced by ion

milling, Appl. Phys. Ltrs. 21 [8] 328-332 (1970).

[143] Murray, R.J., Application of ion-etching of grooves into quartz for surface acousticwave devices, J. Mater. Sci. 1_3 [10] p. 21 05 (1 978).

280
[144] Pinegin, S.V., Production technology of manufacturing aerodynamic grooves by ionic

etching on steel and ceramic bearing plates, Russ. Eng. J. 56^ [2] 52-54 (1 972).

[145] Bulter, A., Generation of hydrodynamic bearing grooving by ion machininq, J. Lub. Tech.

334-336 (1 975).

333-7 s»

[146] Hawkins, D.T., Ion milling (ion-beam etching), 1 954-1 975: A bibliography, J. Vac. Sci.

Technol. 12. [6] 1 389-1 398 (1 975).

Discussion

RUSSELL

With reference to hydrodynamic machining:

1. Is the unit used portable?

2. What water flow rates and pressures would you use to cut dense alumina brick with

10% porosity?

3. Any special safety precautions.

4. Can this replace the water-cooled diamond blade cut-off saw?

5. Can you give me some idea of the cost of a unit for cutting refractory brick?

FIRESTONE

1. We have two water jet machining units; a stationary laboratory unit and a portable,

gasoline powered, field unit. The laboratory unit was used for these experiments.

2. For cutting dense alumina bricks, a 0.5 mm (0.020 in) nozzle and a pressure of 700

MPa (100,000 psig) should be used. This requires a flow rate of 70 cm 3 /s (4 GPM) and a power

of 150 kw (200 HP). The cost of such a unit is about $100,000.

3. A water jet is as dangerous as a buzz saw and similar safety precautions are necessary.

4. Water jets can replace diamond saws for cutting refractories although I know of no commercial operation. However, it is used commercially for cutting concrete, rock and

composites.

5. See answer #2 above.

LANDINGHAM

You indicate that cutting or drilling of high density ceramics was not possible with

hydrodynamic machining. I assume you meant on an economical basis. Richard Wesley at the

Lawrence Livermore Laboratory has conducted many tests with water jets and drilled holes in

large blocks of dense ceramics and dense plates of Bi+C , AI2O3, sapphire, glass, etc.

FIRESTONE

You're right. It's not economical to consider water jet cutting of fully dense ceramics

since the cutting rate is from ten to a hundred or more times slower than for ceramics with

open porosity. The rate can not be increased by increasing the pressure because above a certain pressure the ceramic shatters along the cut. My interest in water jet cutting is as

a replacement for diamond sawing.

281

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

SHAPING SILICON COMPOUND CERAMICS WITH A CONTINUOUS WAVE CARBON DIOXIDE LASER

Stephen M. Copley 1

, Michael Bass 2

, and Russell G. Wallace2

department of Materials Science

2 Center for Laser Studies

University of Southern California

Los Angeles, CA 90007

Results are presented on the use of a high power continuous wave carbon

dioxide laser to shape workpieces of Si 3N4 , Si Al ON and SiC. Several shaping

operations have been investigated in the turning configuration including

grooving, threading and the production of convex and concave surfaces. A

straightforward analysis is presented, which predicts the feed and beam

power corresponding to a specific surface roughness grade and effective

material removal rate.

Key words: Laser; shaping; silicon compound ceramics.

1. Introduction

The application of the continuous wave carbon dioxide laser to processes such as straightline cutting and hole drilling is well established [1]. In our work we have employed the CW-C02 laser to shape metallic materials [2] and recently ceramics in the turning configuration. In this paper, we present results on the use of a 450 W continuous wave carbon dioxide laser to

shape Si^Nij, SiAlON and SiC workpieces. A straightforward analysis is presented, which predicts

the feed and power corresponding to a specific surface roughness grade and effective material

removal rate.

2. Experimental Procedures

The apparatus for laser machining of ceramic materials is shown schematically in figure 1.

The laser beam is directed along a path that is parallel to the turning axis of the lathe. It

WORKPIECE

Figure 1 Apparatus for laser machining ceramic materials

283
is reflected from a mirror mou.ited on the carriage along a direction perpendicular to the turning axis to a mirror mounted on the cross-slide. After reflection by the cross-slide mirror 1tl is focussed by a ZnSe lens on the workpiece along a radial direction for turning or along the

turning axis for facing. Because the direction of motion of the carriage is parallel to the

turning axis and the direction of motion of the cross-slide is perpendicular to the turning

axis, the angular relationships among the reflected beams remains constant during normal

machining operations. A small He-Ne laser is employed for alignment purposes.

In our work, we have employed a 450 W, CW, CO2 laser [\ = 10.6 pm) operating in the TEM

mode (Gaussian spatial distribution of intensity). For the experiments reported in this paper,,

the beam diameter was 160 ym (1/e 2 diameter) giving an incident power density of 22.3 GW m" 2

.

A description of the materials we have shaped in our experiments is given in Table 1.

Table 1. A description of materials investigated.

Material Description

SiC NC 435 (Norton Co.) A

sintered SiC containing

20 vol pet free Si

.

NC 132 (Norton Co.) A

hot-pressed Si 3Ni+

.

SiAlON A hot-pressed SiAlON.

3. Results

Figure 2 shows the cross section of the grooves of a laser machined spiral in SiC. The

grooves were cut in the facing configuration; the high numbered grooves were closest to the

turning axis and thus correspond to the lowest cutting velocities. It is evident that the

groove depth increases with decreasing cutting velocity while the groove width decreases

slightly, being approximately equal to the beam diameter. In figure 3, the groove depth is

plotted as a function of dwell time (0). The dwell time is given by the equation, 0 = b/v,

where b is the beam diameter and v is the cutting velocity.

12 13 14 15 16 17 18 19 20 21 22

Figure 2: Photomicrographs of laser machined spiral grooves in SiC

284
' I 1 1 2 3

Dwell time (Sec) xlO

Figure 3: Depth of laser machined grooves

in 3iC versus dwell time

A major concern in evaluating the potential of shaping with a laser is the quality of the

resulting surface. Figure 4 shows the topography of laser machined grooves in SiC. Although

no cracks were observed in the grooves of this material or the others examined, deposits

suggestive of a condensate or oxide were observed. The mechanical properties of specimens

with laser machined surfaces has not yet been investigated.

Figure 4: Scanning electron micrograph of

laser machined surface in SiC

Figure 5 shows the cross-sections of grooves 2 and 18 at a higher magnification than in

figure 2. It appears that their shape is approximately Gaussian. The mechanism of material

removal is believed to be evaporation.

285
(a) (b)

Figure 5: Cross section of laser machined grooves in Si C (a) groove number 2,

(b) groove number (18).

4. Analysis of Laser Machining

The shape of the laser machined grooves can be described employing the normal curve

2 D(y) = D exp (- K ) (1)

a Tracings of grooves 2 and 18 along with normal curves corresponding to the parameters given

in table 2 are shown in figure 6.

The total area under the normal curve is

A = D aVT (2)

The material removal rate is

I = Av (3)

and the specific cutting energy neglecting power expanded by the lathe is

P= j (4)

where P is the incident beam power. Values for A, Z and p for grooves 2 and 18 are also

given in table 1. It is interesting that the cross section of groove 2 is less than that

of groove 18 but the material removal rate of groove 2 is the greatest. Since the incident

Table 2. Parameters describing laser machined grooves,

Groove DQ (ym) a(ym) A10" Bm

2

) zOCTVs" 1

) p(GJm" 3

)

2

18

125

300

100

75

2.22

3.99

3.42

2.19

131

204

286
y

(a) (b)

Figure 6; Tracings of laser machined grooves and normal curves based on parameters

given in Table 2: (a) groove 18 (b) groove 2.

beam power was the same for both grooves, the specific cutting energy the groove 2 must be

the least. Thus, the faster the beam scans the surface, the more efficient becomes the

laser machining process. Although the origin of this effect has not been determined, it may

be due to lower conductive losses or possibly better coupling of the laser beam to the workpiece at high velocity.

Actual turning or facing involves the partial overlapping of laser machined grooves as shown in figure 7. In the following analysis, we shall assume that the shape of the groove

cut by the laser in turning is the same as in figure 2 even though in turning it is cut into

the edge of a step. The surface roughness (R), which is equal to the sum of the absolute values

of all the areas above and below the mean line divided by the sampling length [3], is given by

the equation

2

287
The effective material removal rate, which is the material removal rate corrected for the over- lapping of grooves is given by

V = vD

Q

a y/jT - v(geh) (6)

The areas abc, cde and geh needed to evaluate the surface roughness and the effective material

rate can be evaluated with the help of tables giving the ordinates of the normal curve and the areas under the normal curve [4]. Accordingly, the depth of the laser machined groove is

given by

D(y) = DVH f(z)

2 = yv^~

a

and the area of the laser machine groove is given by

-y

D(y) dy = D

Q a/F F(z)

(7a)

(7b)

(8a)

(8b)

where

f(z) exp (

f(i) -Jl m d

(9)

(10)

The surface roughness can be determined from eq. (5) and is given by

r = i [d

0

^

H^f ) - d y*j (ID

where y* is the abcissa value when D(y) equals the mean diameter see figure 7Z The mean diameter can be found by equating abc (the quantity in brackets in eq. (11)) to cde, and solving

for D, where

cde - D( f

-y*) -[30

avrfT F(f V272a) ^ ^ ^^,3 (12)

Following this procedure, we obtain

2D avT ~K"

n = —o f( — ) (13)

The parameter y* can be calculated with the equation

+ a ^-1| (14)

288
The effective material removal rate is obtained from eq. (6) and is given by

i = 2vD

Q

a vf F( (15)

Equations (11) and (15) are employed in figure 8 to make a plot of surface roughness and_

3 effective material removal rate for a groove cut with incident power density of 22.3 GWm »

beam diameter (1/e 2

) of 160 ym and dwell time of 1 .05 x 10

5 s (see groove 2, table 2,

figure 5). The feeds corresponding to ISO roughness grades N8, N9 and N10 are indicated

on the abcissa of figure 8. A feed of 0.05 mm rev is normally available on engine lathes.

As a rough quide for comparison purposes [3], turning operations in metals normally give

surface roughnessess greater than 1 ym.

Figure 8: Roughness and effective material

removal rate versus feed

0.05 0.10 0.15

Feed (mm rev )

It is easiest to understand the meaning of the results given in figure 8 by considering

a specific machining task. Consider, for example, the turning of a cylindrical workpiece of

SiC 12.5 mm in diameter to the final shape shown in figure 9. The total depth of cut is

1.75 mm; for conditions of cutting corresponding to groove 2, 14 passes of the laser beam

would be required. For a dwell time e * 1.05 x 10" s" , nw

= b/Tred = 4.6 s" 1

, where 3 is

1/2C12.5 t 9.0) mm. If it is assumed that a return time of 2s is required for each pass, then

the time for each pass in seconds including the return time is given by the equation

<5t = (L/n f) + 2 where L is the length of the pass (12.5 mm in figure 7). The total machining

time t =w

14 at.

Figure 9: Workpiece for laser machining

cal culations

Dimensions in mm

289
i

Table 3 gives the time for each pass and the total machining time for the workpiece shown

in figure 9 for feeds corresponding to roughness grades N8, N9, and N10. The times tabulated

in the column marked P = 450 W correspond to the irradiation conditions considered in figure 8. The times tabulated in the column marked P = 1400 W correspond to machining at a higher velocity

or shorter dwell time such that the product P © is the same as the previous case. These values are based on the assumption that the shape of the groove depends primarily on the total energy

absorbed by a circular area equal in diameter to the spot size. This assumption has not yet

been experimentally tested.

Table 3. Time to machine workpiece shown in figure 9 with a specified surface finish for two

values of incident beam power.

P = 450 W P

= 1400 W

Roughness f (mm) 6t (s) t

M

(min) fit (s) t

M

(min)

N8 0.060 47.7 11.1 14.7 3.4

N9 0.083 35.0 8.17 12.6 2.9

N10 0.130 21.1 4.92 8.8 2.0

5. Applications

Figure 10 includes two macrophotographs of a 1/4" x 20" thread turned on Si 3N4 by laser

machining. In figure 10a the retangular rod from which the threaded section was turned can

be seen. Figure 10a also shows a region in which we demonstrate the ability to produce

concave surfaces.

Figure 10: Macrophotographs of laser

machined Si^^

The negatively curved surface was made by focussing the laser beam so that it struck

the center of the rectangular surface at normal incidence. Thus the corners were out of

the focal region and irradiated at near grazing incidence where the light-to-material coupling was less. As a result, more material was removed near the center than near the corners

and a concave surface was established. A convex surface (i.e., the round section which was threaded in figure 10) was made by adjusting the focus and angle of incidence so that the corners were near the focus and the light near normal incidence at the corners. This makes

the central part of the surface out of the focal region and illuminated at near grazing incidence. Thus more material could be removed near the corners than near the center when the

light was so focussed.

290
When laser machining was used to thread Si 3 n h the focussed beam was at normal incidence

on the cylindrical surface. However, to make a proper thread the position of the focus with

respect to the previous spiral groove must be adjusted. This is shown in figure 11 where the

sequence of focal positions are sketched.

!

Figure 11: Schematic diagram of the

procedure for laser ma- chining screw threads

Figure 11 shows a laser machined piece of Si AT ON which originally had the square cross section remaining in the middle of the piece. The maximum peak to valley non-uniformity of

the smooth surface is 7.5 ym which confirms the potential of laser machining Si Al ON to desired

surface quality specifications. Figure 12 also shows the 1/2" x 13 screw thread which we laser machined on the other end of the laser turned SiAlON rod.

Figure 12: Laser turned and threaded

SiAlON

6. Summary

1. A high power CW carbon dioxide laser can be employed to shape workpieces of SiC, Si<jN 4 and SiAlON.

291
2. The material removal rate increases and specific cutting energy decreases with increasing;*

cutting velocity.

3. For a specific cutting velocity, an analysis is presented, which shows that the maximum

feed and the effective material removal rate are limited by the allowable surface roughness.

4. Several shaping operations including grooving, threading and the forming of concave and convex surfaces have been demonstrated in the turning configuration.

I Acknowledgments

This research is supported by the Advanced Research Projects Agency under Contract No.

00014-77-C-0478, ARPA Order No. 3421.

The authors wish to thank Dr. E. R. Thompson at United Technologies Research Laboratory

and Dr. S. Musicant, General Electric Co. for providing materials used in this investigation.

References

[1] Duley, W. W., C0 2 Lasers, Effects and Applications (Academic Press, New York, 1976).

[2] Bass, M., Copley, S. M., and Beck, D., Laser Assisted Machining, in Proceedings of the

4th European electrooptics conference , Ultrecht, Netherlands, October 1978.

[3] Boothroyd, G., Fundamentals of Metal Machining and Machine Tools , p. 134, (McGraw-Hill

Book Co. , New York, 1975).

[4] Kennedy, J. B. and Neville, A. M., Basic Statistical Methods for Engineers and

Scientists , (Dun-Donnelley Publishing Co., New York, 1976.

Discussion

COOPER

Have you used supplementary lenses to spread out, elongate, or otherwise modify the

shape of the spot during machining? If the efficiency is increased at short dwell times,

why not spread out the beam normal to the transversing direction?

COPLEY

No, because of diffraction effects you cannot focus the beam to a narrow line perpendicular to the transverse direction that would have as large an average flux density as the

Gaussian spot.

RICE

Your work on laser machining of ceramics has apparently been concerned with materials

that mainly vaporize without melting. Have you investigated materials that melt such as glass, AI2O3, etc., and is the associated use of a diamond tool also not effective with

them?

COPLEY

We have investigated laser assisting machining of AI2O3 using an Al203 -TiC cutting tool.

The tool was not effective in removing material. However, without the tool material, removal was observed by what appeared to be microspalling for conditions such that the surface temperature was less than the melting point.

292
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

I R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

OPTICAL AND SURFACE PHYSICAL CHARACTERISTICS OF DIAMOND-MACHINED INFRARED WINDOWS 1

D. L. Decker, D. J. Grandjean, and J. M. Bennett

Michelson Laboratories, Physics Division

Naval Weapons Center, China Lake, California 93555

This paper presents the results of the diamond turning of surfaces of a wide range of single-crystal and polycrystal 1 ine infrared window materials,

including Ge, CaF2, MgF2, SrF2, KC1 , and GaAs. These surfaces were obtained

from an advanced air-bearing, diamond-turning machine recently brought to an operational state at the Naval Weapons Center. The surface characterization

performed includes a wide range of optical and surface physical examination

including an analysis of scattered and absorbed light, and a surface microtopographic examination by diamond profilometry, phase contrast interference

(Nomarski), and scanning electron microscopy. These data are compared with

earlier results and are related to the characteristics of the Naval Weapons

Center diamond-turning machine and to the machining parameters. A companion

paper at this conference discusses the laser damage resistance of these

surfaces with implications for their use in laser applications.

Key Words: Diamond single-point machining; dielectrics; infrared optical

components; laser damage; optical properties.

Introduction

Diamond single-point turning or milling as a manufacturing method for precision

optics has been demonstrated at a number of laboratories, notably the Energy Research and

Development Administration complexes at Livermore, Calif., and at Oak Ridge, Tenn. [1,2]. 2 The advantages of this technology over conventional methods include significant cost savings

in many cases but, more importantly, the ability to readily generate aspheric surfaces

which would be difficult or impossible otherwise. The materials which are amenable to

diamond single-point machining are not limited to metals such as aluminum and copper and

many of their alloys, but also include an impressive list of infrared window materials,

useful as domes, windows, and lenses (see table 1). Much recent interest in diamond singlepoint machining is a direct result of the superior optical components for high energy laser

application that can reliably be produced by this method. However, there are many low

power laser or imaging optics applications for which the method offers significant system

performance improvement over what can be done with conventional components [3].

This paper presents surface physical and optical characterization results for dielectric surfaces machined on the Naval Weapons Center (NWC) diamond-turning machine. The

optical fabrication criteria outlined in table 2 are given as a reminder of the finishing

requirements for a component. This paper is primarily concerned with surface microtopography and the directly related topic of scattered light [4,5]. Most of the previously

reported work has been of a qualitative sort, and has not been followed up by detailed

optical characterization or optimization of results [6]. This paper presents for the first

time data relating to the mechanisms of material removal, and the relationship of these

mechanisms to the crystalline orientation of the work surface to the machining geometry.

x Work supported by Naval Weapons Center Independent Research Funds.

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

293
Table 1. List of dielectric materials which have been successfully

diamond single-point machined as of 1979.

Ca ?2 NaCl ZnS

iNa r 7n C0

SrF

2 Si LiNb0

3

KCl Ge GaAs

acrylic plastics

Table 2. Summary of the fabrication requirements

which apply to most optical components.

Optical figure

Surface absorption

Scattered light

Damage threshold

Machining Technique and the Material Removal Process

The basic diamond-machining process discussed here is that of single-point turning or milling of a surface in the fashion of ordinary machine shop fabrication, but using a very

precise machine and special diamond tools (fig. 1). The diamonds used are of gem quality

and are mounted and polished to a supersharp, highly accurate, radiused cutting edge. The

overall view in figure 2 shows the tool and spindle slides of the NWC machine which are arranged in a T-base configuration [7]. The part being turned in this picture is a 6-inchdiameter copper flat which is held on the spindle nose by a vacuum chuck.

PLANE MIRROR ON INSIDE OF CHUCK

Figure 1. Isometric schematic of the NWC two-axis diamond-turning

machine showing the placement of the spindle and tool slides in

a T-base configuration. Also shown is a simple interferometer

layout for measuring the primary motion of the slides.

294
Figure 2. Photograph of the NWC machine showing the tool and spindle

slides, the tool post assembly including precision height adjuster,

chip/coolant guard, and spindle housing. A 6-inch-diameter copper

flat is being machined.

In the case of a metal workpiece, material is removed by the tool in machining by

shearing a chip from the surface. The idealized surface profile in a diametral plane

perpendicular to the work surface for a turned surface is shown in figure 3. The tool nose

radius r is typically 1/8 inch, and the tool feedrate is in the range of 10 _ 4 inches per

spindle revolution. The theoretical peak-to-valley roughness in this case is 10"8 inches.

In actual practice, as the feedrate is made smaller, the surface finish does improve as predicted, but approaches an asymptotic roughness, controlled by material imperfection and

irregularity in the machining process. The latter effects include very high frequency (10^

Hz) chip vibration/surface interaction effects which are possibly excited by chip/tool

friction [7,8], Tools and feedrates typically used to machine dielectric surfaces are similar to those used to machine metal surfaces. In contrast with metal machining, there is no strong tool/surface speed dependence to the resulting dielectric surface characteristics.

ASSUMING THAT f/r « 1

OR RMS VALUES:

rmi mix

Figure 3. Ideal finish achieved by machining with

a tool of nose radius r and a feedrate f. Both

peak-to-valley (pv) and rms values are given.

295
Material removal from all dielectrics examined in this paper occurs apparently as a consequence of brittle fracture. There is no evidence either on the resulting surface or

on the machining debris of plastic deformation. This statement is based upon scanning

electron microscopy of the machined surfaces and machining debris. No subsurface examination has yet been made. A pitting effect occurs which is very sensitive to tool geometry

and to crystallographic orientation. The location of "clear" and "cloudy" areas turned surfaces of several materials is shown in figure 4. The visual effect of

area is a consequence of light scattering from microscopic pits. When turned on single-crystal materials develop a clear/cloudy pattern which will usually have

tional symmetry of the crystallographic plane exposed by machining

cubic lattice is tipped only approximately 20° away from the [111]

has threefold rotational symmetry. Samples a) and b) of CaF2 have

orientation, with sample b) showing the expected symmetry, but a)

on diamonda cloudy

center,

the rotaThe [331] axis in a axis which, of course,

virtually the same

shows curious skewed

twofold symmetry. The origin of the effect discovered in sample a) is unknown, and this

anomaly was observed only in this case out of several dozen samples of different materials

and a variety of orientations which were turned in this study.

CALCIUM FLUORIDE

[110]

GERMANIUM

[001

POTASSIUM CHLORIDE

[116] [116] [010]

[110]

Figure 4. Schematic showing the locations of "clear" and "cloudy"

structures resulting from diamond single-point machining relative

to crystallographic directions for several different materials

and orientations.

It is useful to examine these surfaces microscopically both with a light microscope

and by electron microscopy at much higher magnification. Figures 5 and 6 are Nomarski

interference contrast micrographs of an Irtran MgF2 sample and single-crystal CaF2 sample

b). In all cases, the samples were originally polished and were subsequently diamond

turned, leaving a central polished spot for direct comparison. As shown in figure 5, the

quality of the Irtran surface is substantially limited by voids remaining from processing

of the material, both in the polished and diamond-turned areas. In the case of singlecrystal CaF2 (fig- 6), the clear diamond-turned area is indistinguishable from the polished

area. The scanning electron micrographs of figure 7 compare CaF? &) polished and diamondturned clear areas at much higher magnification. The polished surface is defined by a superposition of scratches from the polishing grains. Most of the clear diamond-turned area is an extremely smooth, almost featureless, surface. The localized pitted areas often show

the effects of multiple tool passes, and a very interesting repetitive nature with a spatial

period which varies from slightly less than 1 um down to 0.5 ym (fig. 7). This variation

is seen to occur even over a very small area and shows no correlation with position on the

sample. Figure 8 shows the localized nature of the cloudy areas and a high magnification

view of the pits in a cloudy area. There is no -qualitative shape or size difference of the

pits occurring in clear or cloudy areas —only concentration differences. Often a decrease

296
POLISHED

J

1000 urn

DIAMOND TURNED

Figure 5. Nomarski phase contrast interference

micrograph of a polished Irtran surface which

has been partially diamond machined leaving a polished area for comparison.

POLISHED POLISHED

L

1000 Mm

DIAMOND TURNED 'CLOUDY'

1000 nm

DIAMOND TURNED

Figure 6. Nomarski phase contrast interference micrograph of a polished single-crystal

CaF2 surface which has been partially diamond machined leaving a polished area for

comparison. On the left side is a representative "cloudy" area and on the right a "clear" area.

297
V

mam

5 (Jim 5 ptm

POLISHED AREA CLEAR AREA

Figure 7. Scanning electron micrograph of the same single-crystal CaF2 surface

shown in figure 6. On the left side is a polished area and on the right an area selected from the clear diamond-machined portion showing localized pitting

and some debris from turning.

I I I 1 100 titn 5 urn CLEAR AND CLOUDY AREAS CLOUDY AREA

Figure 8. Scanning electron micrograph of the same single-crystal CaF surface

shown in figures 6 and 7. On the left side is a low magnification view of the

localized "cloudy" spoke-shaped area as it is seen near the boundary between

the polished and diamond-turned portions. On the right, at much higher magnification, is shown a representative "cloudy" area.

298
in spatial period can be correlated with a decrease in pit volume, in a sequence of pits.

The pit shape is, of course, strongly suggestive of an indentation mark made by the tool

tip. However, the relative motion between the surface being turned and the tool tip is a factor of at least 40 smaller than the typical 1-ym-deep pit. A probable explanation is

that the pits are produced by spall ing as a result of a localized tensile load at the

surface which exceeds the material strength following the release of the compressive loading at the cutting edge as the tool passes over the surface. The repetitive nature of the

spall pit formation comes from a constructive interaction of the tool loading and a standing

acoustic wave in the material which has been excited by the tool. The effect is similar to

chatter often observed in an ordinary machine tool, except that it is not a machine structure resonant effect, but an effect in the material itself. The effect is similar to that

observed in diamond machining of metal surfaces [7,8], The standing wave may or may not be

a consequence of a resonant effect. However, the temporal frequency of the material "chatter" is approximately 10^ Hz which is on the order of the resonant frequency of a mass on a spring consisting of a spall fragment connected into the lattice of the work material. In this simple model, the resonant frequency would scale as the square root of the reciprocal

of the volume of the spall pit. This relationship is at least qualitatively observed in

the decrease in spall pit size and accompanying decrease in spacing. The spall pit sequence

on the left of the superposition shown in figure 7 is a clear example of this effect. A

resonant explanation thus seems likely.

As the cutting radius increases at constant rotational rate, it has been observed in a number of cases that the spall tendency precipitously can decrease beyond a specific radius

(surface speed). This may simply be due to the lack of constructive interaction of the

tool and the resonant wave in the material. If so, the tendency to spall should occur in

annular bands, depending upon surface speed. This effect has not been observed. Another

possible explanation is that spall formation is dependent upon strain rate and hence surface

speed. This possibility is further corroborated by the curious fact that although the

concentration of spall pits is strongly dependent upon crytallographic orientation, the

surfaces defining the shape of the pits are formed by brittle fracture along arbitrary

crystallographic planes. There is no evidence either in the pits or on the fragment sur- faces of cleavage. Fracture produced in brittle materials at very high strain rates in

shock experiments often does not recognize the plane of easy division (cleavage). The

tendency towards spall formation can be decreased by two known approaches: 1) higher or

lower surface speeds and 2) tool rake angle. By using a tool with a very large negative

rake, in the range of 20 to 45°, very uniform surfaces have been turned on several dielectric materials including KC1 and CaF2- However, the success with both surface speed variation and tool geometry are very dependent upon the cutting edge characteristics of the tool

being used. Edge sharpness and orientation, and perhaps even the diamond bulk imperfection

is involved in the resonant chatter effect.

Optical and Surface Physical Characterization

A diamond stylus profilometer has been used to examine the topography of the several

diamond-turned surfaces reported here. The results shown in figures 9 through 12 were

obtained from a Talystep profilometer with a 1-ym-radius conical stylus. This instrument

yields surface profile data along the track of the stylus with a lateral resolution on the

order of a micrometer for the smoother, low slope surfaces examined here, and with a height

resolution of 3 A\ The output of the stylus transducer amplifier is digitized and subsequently processed to yield distribution functions for height, slope, and correlation [9].

In figures 9 and 10 unprocessed scans of polished and diamond-turned single-crystal CaF2

are compared for two different scan lengths, 15 and 600 ym. The diamond-turned clear areas

on the CaF£ sample of figures 9 and 10 are a factor of three smoother than the corresponding polished surface on the same sample, and have much less steeply sloped structure as well. The Irtran MgF2 surface scans shown in figures 11 and 12 indicate a much rougher

surface, the surface being substantially determined by voids still present in the hot

pressed material. The diamond-turned surface is uniform and slightly rougher than the

polished surface.

299
POLISHED

SCAN PARALLEL TO

GROOVES. CLEAR AREA

15 jum

Figure 9. Unprocessed Talystep surface scan of 15-ym length on the CaF2 sample previously shown in figures 6 through 8. Both

polished and diamond-turned "clear" areas are shown.

POLISHED

85 £ RMS ROUGHNESS

S

o<S

SCAN PARALLEL TO

GROOVES. CLEAR AREA

600 iim

Figure 10. Unprocessed Talystep surface scan of 600-ym length

on the CaF2 sample previously shown in figures 6 through 9.

Both polished and diamond-turned "clear" areas are shown.

The surface profiles of the polished and clear diamond-turned areas thus are very

similar in distribution of height and slope. Consequently, the light scattering from the

two surface types will be similar- The magnitude of the total integrated scattered (TIS

)

light compared to that light specularly reflected, AR/R, can be related to the rms roughnes

6

AR/R £ (4^6/A)'

300
15 nm

Figure 11. Unprocessed Talystep surface scan of length 15-pm

on the polycrystal 1 ine MgF2 previously shown in figure 5.

Both polished and diamond-turned areas are shown.

POLISHED

43 X RMS

ROUGHNESS

WITHOUT PITS

DIAMOND TURNED

inrinr^

.60 & RMS

TIOUGHNESS WITHOUT PITS SCAN PARALLEL TO GROOVES

600 (jm

Figure 12. Unprocessed Talystep surface scan of length 600-wn

on the polycrystal 1 ine MgF2 previously shown in figures 5 and

11. Both polished and diamond-turned areas are shown.

where x is the wavelength of the incident light [10]. Table 3 gives values of TIS for

Irtran MgF2> CaF2» and SrF2 for light in the visible (0.647 urn) and in the near infrared

(3.39 ym). There is little difference in the visible light scatter from the clear diamondturned polished areas, with the cloudy diamond-turned areas on CaF2 and SrF2 scattering a factor of five more light than the clear areas. In the near infrared the differences

between polished and clear diamond-turned scattered light is even smaller, with still a factor of roughly five difference between clear and cloudy. From an examination of the

CaF2 diamond-turned surfaces shown in figures 10 and 11, it is obvious that much debris

from machining is present even after "cleaning." It is certainly true that much of the

scattered light originates from these particles, and that more complete cleaning would

materially reduce the TIS, especially in the infrared. The presence of machining debris

301
also has a negative impact upon the laser damage resistance of the machined surface, especially for short pulse lengths where dielectric breakdown is a predominant damage mechanism

[11]. However, in a companion paper at this conference [12], no significant difference in

laser damage thresholds for the polished and diamond-turned surfaces described in this

paper were discovered for 100-ysec pulses of 2.7- and 3.8-ym wavelength laser radiation.

Table 3. Values of TIS at two different wavelengths, 0.647 and 3.39 ym,

for three different materials which were diamond single-point machined

in this study. Included is a comparison of polished and diamond-turned

"clear" and "cloudy" areas on the same sample.

Material

Wavelength, pm Pol ished

Diamond

Cloudy

turned

Clear

MgF

2 0.647 .02 .03 .03

( Irtran) 3.39 .006 .01 .01

Cat

2 0.647 .0007 .04 .001

3.39 .00004 .004 .00003

SrF

2

0.647 .0008 .02 .004

3.39 .0001 .0006 .0001

Conclusion

The diamond stylus profilometer results from diamond-turned dielectric surfaces are statistically not much different than surfaces produced by conventional grinding and polishing by loose abrasive. As a consequence, the TIS in the visible from a diamond-turned

surface can be quite comparable to that of a high quality polished surface, and in some cases even lower. For the materials examined in this paper, Irtran MgF2> SrF2, and CaF2,

the diamond-turned surface scattered light at 3.39-ym wavelength is on the order of that

from the corresponding polished surfaces. As was shown, the diamond-turned areas are covered with a significant concentration of debris from turning, even after "cleaning."

Much of the measured scattered light certainly originates from these particles. More

elaborate cleaning must be accomplished. The scattering from the "cloudy" diamond-turned

surfaces is much larger, as a consequence of a much larger concentration of localized spall

pitting. A full understanding of the latter effect and its elimination are important to

the fully successful application of diamond turning to infrared optics. However, the

Irtran MgF2 results are very interesting and encouraging since this polycrystal 1 ine material is often used for low cost infrared optics. There are no preferential cloudy diamondturned areas on this surface and the diamond-turned surface scatters only slightly more

light than the polished area. It is thus demonstrated that under some conditions diamondturned dielectric surfaces can be quite comparable in optical quality to corresponding

polished surfaces. The advantages that diamond single-point machining brings to the production of infrared optical components, high throughput, and ready generation of aspheric

surfaces promise to revolutionize infrared optics production.

Acknowledgment

The authors of this paper acknowledge the contributions of many co-workers at NWC,

notably Phil Archibald for scattering measurements, Cliff Fountain for x-ray orientation of

samples, and Bev Reymore for editorial assistance.

302
References

[1] Arnold, J. B., and Steger, P. J., Diamond turning on advanced machine tool prototypes

(SME Clinic on Industrial Application of Precision Machining and Gaging, 19-21 November 1974, San Francisco, Calif.) Report Y-DA-5975 (Union Carbide, Inc., Oak Ridge,

Tenn. , 1975).

[2] Bryan, J. B. , et al., A practical solution to the thermal stability problems in

machine tools, Technical Paper MR72-138 (Society of Manufacturing Engineers, Dearborn,

Mich., 1972).

[3] Decker, D. L., Soileau, M. J., Porteus, J. 0., and Bennett, J. M., Surface, optical,

and laser damage characteristics of diamond turned metal mirrors, in Laser Induced

Damage in Optical Materials: 1976 , A. J. Glass and A. H. Guenther, eds., pp. 153-

164. NBS Spec. Publ. 462 (National Bureau of Standards, Washington, DC, 1976).

[4] Decker, D. L., Bennett, J. M. , Soileau, M. J., Porteus, J. 0., and Bennett, H. E.

,

Surface and optical studies of diamond-turned and other metal mirrors, Opt. Engr. j_7,

[2], 160-166 (Mar/Apr 1978).

[5] Decker, D. L., and Bennett, J. M., Surface evaluation techniques for optical components

(Proc. Society of Photo-Optical Instrumentation Engineers) Optical Coatings II, 140 , 32-41 (1978).

[6] Arnold, J. B., Morris, T. 0., Sladky, R. E., Steger, P. J., Machinabil ity studies of

infrared window materials and metals (Proc. Society of Photo-Optical Instrumentation

Engineers) Advances in Precision Machining of Optics 93_, 96-103 (1976).

[7] Decker, D. L., and Grandjean, D. J., Physical and optical properties of surfaces

generated by diamond-turning on an advanced machine, in Laser Induced Damage in

Optical Materials: 1978 , A. J. Glass and A. H. Guenther, eds. , pp. 122-1317 NBS

Spec. Publ. 541 (National Bureau of Standards, Washington, DC, 1978).

[8] Decker, D. L. , Presented in a paper read before the Society of Manufacturing Engineers, Precision Machining Workshop, 27-29 June 1978, Williamsburg, Va.

[9] Elson, J. M., and Bennett, J. M., Relation between the angular dependence of scattering and the statistical properties of optical surfaces, J. Opt. Soc. Am. 69_, 31

(1979).

[10] Bennett, H. E. , and Porteus, J. 0., Relation between surface roughness and specular

reflectance at normal incidence, J. Opt. Soc. Am. 5]_, 123-129 (1961 ).

[11] Temple, P. A., and Soileau, M. J., Resonant defect enhancement of the laser electric

field, in Laser Induced Damage in Optical Materials: 1976 , A. J. Glass and A. H. Guenther, eds., pp. 371-378. NBS Spec. Publ. 462 (National Bureau of Standards,

Washington, DC, 1976).

[12] Porteus, J. 0., Soileau, M. J., Faith, W. N., and Jernigan, J. L., Surface characteristics of diamond-machined window surfaces related to 2.7- and 3.8-ym laser damage

(this conference).

Discussion

BOWLES

In view of the similarity of the cracks shown in the diamond point surfaced to Hertzian

cracks generated by a sliding indenter, might not the "frequency" of spall distribution arise

from the flaw population in the surface being worked?

303
DECKER

It is likely that a spall is located by a flaw, but once initiated it appears that the

spatial frequency of additional spall sites is controlled by intrinsic material parameters.

Extensive repetitive spalling is observed in the high susceptibility regions which is too

regular to be accounted for by extrinsic material properties (flaws).

LANDINGHAM

I have examined a single-point machined fl uorphosphate glass which shows spalling in some regions of surface but other regions are very smooth. This coarser spalling seemed to

be related to rake angle and tool holder vibrations. Smooth surfaces were examined in an SEM and no submicon spalling cracking could be detected.

DECKER

The tendency to spall in the crystalline materials reported on here is very strongly

related to rake angle, tool sharpness, and surface speed. However, I don't believe that

tool holder vibration plays a significant role, since the observed temporal frequency of

the spall is in the megahertz range and thus is orders of magnitude higher than any resonant

frequencies associated with the tool holder assembly.

RICE

1. You noted larger, apparently sporatic surface pits in MgF2 from processing defects

(mainly pores?)? Can you estimate the surface density and size range of these to indicate

the character and distribution of such processing defects?

2. You stated that the removal mechanism appeared to be by brittle fracture rather

than by a deformation process. This is very surprising, especially for several of the

materials you have studied. What is your evidence for this statement?

3. Can you elaborate on cost factors, or ititial investment and operator skill vs. polishing? Can you comment on possible applicability of single point diamond machining

to harder ceramics such as AI2O3 or MgAl 2 0it?

DECKER

1 The hot pressed MgF 2 used in this study has a defect density of approximately

lOVMM2

. These defects are primarily pores of from 0.5 - 3 ym in diameter.

2. It is quite possible that plastic deformation plays a role in spall function, but

there is no superficial evidence of plastic flow either associated with the spall pits or with the fragments so produced. For example, there is no upturned edge on any spall pit

examined and all fragment edges are very sharp, right to the limit of resolution of the SEM.

3. There is no question that precision single point diamond machining requires a large capital investment, but for high volume production, especially of aspheric or non- standard components, significant cost savings have been demonstrated. The operator skill

level required for precision diamond machining is very similar to that required for conventional polishing, involving to a substantial degree technique in handling of optical surfaces. Hardness of the work outlined, per se, is not a criteria for determining the diamond

machinabil ity of a material. I have been told by several people that diamond tools have

been used to finish AT 2O3 parts mechanically. It may thus be entirely possible to produce

an optical finish by use of a suitable tool and machine.

304
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

ELECTROCHEMICAL GRINDING AND SURFACE DAMAGE FOR CEMENTED CARBIDES

S. Mai kin and R. Levinger

Materials Processing and Machine Tool Center

Faculty of Mechanical Engineering

Technion-Israel Institute of Technology

Haifa, Israel

The electrochemical grinding process (ECG) for WC-Co cemented

carbides is shown to be governed by differences in electrochemical

dissolution rates into the two phases. An electrochemical analysis

indicates that the initial electrochemical etching rate into the

cobalt phase is much faster than into the WC. This facilitates

mechanical removal during ECG but produces a damaged surface weakened by partial cobalt depletion. The degree of surface damage

is quantitatively measured as the reduction in mechanical grinding

resistance as a function of depth below the finished surface.

Key Words: Cemented carbides; electrochemical grinding; grinding;

tungsten carbide.

1. Introduction

Electrochemical grinding (ECG) has often been proposed as an alternative to con- ventional mechanical grinding for machining of cemented carbides. Combined

electrochemical and mechanical action in ECG provides faster and more efficient

material removal than conventional grinding. However, the resulting surface quality

is considered to be inferior, and this has restricted the more widespread use of

ECG for machining of cemented carbides.

The conventional analysis of ECG is based upon the assumption that electrochemical

and mechanical removal proceed comDletely indeDendently of each other. While this

assumption may be satisfactory for single phase materials, it is unlikely to be valid

for cemented carbides which consist of hard carbide particles with a ductile metal

binder. With WC-Co cermets, it will be shown that the electrochemical etching rate

into the cobalt phase should be much faster than into the WC phase. Selective removal

of the binder metal in this way should weaken the cermet and thereby facilitate the

mechanical grinding action. However, this same effect may be the cause of poor surface quality after grinding, as the surface produced is likely to be weaker than the

bulk material due to a deficiency in cobalt binder selectively removed by electrochemical dissolution.

- The present paper reviews the results of a recent investigation of the ECG process

for WC-Co cemented carbides with particular consideration of the role of heterogeneous electrochemical dissolution [1, 2] 1

. Differential electrochemical etching

into the carbide and cobalt phase is analyzed and related to the mechanical grinding

energy requirements of the process. Surface damage after grinding is quantitatively

measured as the reduction in mechanical grinding resistance of the material in the

finished surface layer.

2. Electrochemical Considerations

The electrochemical dissolution process is governed by Faraday's Laws. For a single

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

305
phase material, the volumetric electrochemical dissolution rate Z is proportional to the

applied current I: e

with the proportionality constant

where M is the molecular weight, F is the Faraday constant (F=96000 amp), u is the valency

of the electrochemical reaction, and y is the density. For one-dimensional electrochemical

removal normal to a surface of cross sectional area A, the linear removal rate (etching velocity) v is obtained by dividing both sides of eq. (1) by A:

where i is the current density.

Consider now what happens during electrochemical dissolution of WC-Co cermet, a two-phase

material. Let us assume that the WC and cobalt Dhases behave as parallel electrical conductors near the surface. A portion of the total current normal to the surface will flow

through the WC phase and the remainder through the cobalt, the partition of current flow between the two phases depending on the cermet composition and electrical resistivities of

each phase. Taking this into account, the corresponding etching velocities into each phase,

V Co

and V

V!C' Can be obtained by applying eq. (3) to the cobalt and to the WC [1]:

'Co

WC

1 +

*Co

x 1 + e

1 t e

fC

( 1 + eg

)i

(4a)

(4b)

where a- and are the values of a for each material, e is the ratio of electrical resistivity ot cobart to that of WC, and e is the volume ratio of tungsten carbide to cobalt in

the cermet. The relavent properties for WC and cobalt are given in table 1 [3]. The ratio

of v Cot0 V: is:

Table 1. Properties of WC and Cobalt.

WC Cobalt

Molecular weight, M 195.92 58.94

Density, y (-gjjr-) 15.7 8.85

Resistivity, p (ohm»m) 18.5-19.2xl0" 8

5.8xl0" 8

Electrochemical constant,

ct (x „ . ) amp»mm

0.97 2.08

Valence of electrochemical

reaction, u>

8 2

306
v

Gq = _^Cp_ (5)

VWC BaWC

which depends only on the properties of the two phases and is independent of cermet composition. Putting in the appropriate properties of WC and cobalt from table 1 gives a velocity

ratio v Cq

/ vMC

= 7.2.

It can therefore be expected that the etching velocity into the cobalt phase will be

much faster than into the WC phase, thereby causing selective etching of the cobalt. This is

the situation which prevails at the beginning of the electrochemical process. As this pro- cess continues, however, the current flow through the cobalt phase will decrease because of

the high electrical resistivity of electrolyte filling up the gaps left behind by selective

cobalt removal, with a corresponding increase in current through the WC in order to maintain

the total applied current. The decreasing current in the cobalt phase will reduce the etching velocity into this phase from its initial value given by eq. (4a), and the increasing

current in the WC phase will raise its etching velocity above its initial value given by eq.

(4b). Eventually a steady state situation will be reached beyond which the etching velocities into both phases are the same with the cobalt level a fixed depth ahead of the WC

level at the surface. The steady state etching velocity is given by [1]:

v* = a*i (6)

where

*

aCo aWC (1 *e)

a* = (7) m aWC

+ aCo

3. Experimental

ECG experiments were conducted using the face grinding arrangement illustrated in figure

1. This tyDe of grinding allows for observing the influence of infeed velocity v- relative

the etching velocities obtained from the electrochemical analysis. The ECG process parameters were the infeed velocity v. and current density i. During ECG, the force components

307
1

normal to the wheel surface, F , and tangential to the wheel surface parallel to the wheel

velocity, F , were measured wi¥h a semiconductor strain gauge dynamometer. All experiments

were performed using the same grinding wheel (Eleform D100) which contained 100 mesh diamond

of 100 concentration (4.4 carats/cm 3

) in a copper binder. The rotational speed of the wheel was fixed at 5000 r/min corresponding to a velocity V=29.5 m/s at the mid-point of the grinding area which was located 66/mm down from the wheel axis. During ECG, a commercial electrolyte (ELY-S10) consisting mainly of sodium nitrate and nitrite was supplied to the grinding area at a nominal flow rate of 2 liters/min and temperature of 25°C. Eight commercial

grades of WC-Co cemented carbides were used as listed in table 2 together with their composition and properties furnished by their manufacturer.

The weakening of finished surfaces as a consequence of ECG was quantitatively evaluated

for some cases from an analysis of the forces obtained during a grinding transient after shutting off the electrical current at the end of the ECG and continuing with mechanical grinding

at the same overall infeed velocity v. [2], During the transient, the normal grinding force

generally increases because the process is no longer electrochemically assisted, thereby

causing more of the overall infeed to be taken up by elastic deflection of the grinding system

so that the actual infeed velocity drops. Eventually the actual infeed velocity rises again

and approaches the steady state value v.. at the end of the transient. From an analysis of

the elastic deflection of the system during the transient, the actual infeed velocity and the

total infeed depth after switching off the current can be calculated at any time during the

transient. Selective electrochemical removal of the cobalt binder prior to shutting off the

current should weaken the material below the surface, thereby reducing the grinding resistance

of the material removed during the transient. The weakening of the surface layer is expressed:

as the ratio of the mechanical grinding energy at the calculated depth below the surface after

ECG to the normal grinding energy for this material, the comparisons being made at the same actual infeed velocities.

4. ECG Results

ECG experiments were performed on all eight cemented carbides listed in table 2. Illustrative results for two of these materials are presented in figures 2-4 for the supply voltage V to maintain the required current, and the grinding force components F and F as a

s l n function of infeed velocity v. for various values of current density i.

Table 2. Composition and Properties of Cemented Carbides

Material

Number

wt. %

WC

wt. %

Co

wt. %

additions

WC . qram

Vicker's

hardness

(kq/mrn2

)

TRS

a

(N/mm2

)

e

1 93.5 5 1.5 TaC 1.2 1850 2050 in. 71

2 91.7 5.8 2.5 Tac 1.8 1700 2200 9.15

3 85 13 2 TaC 1.5 1330 3150 3.77

4 94 6 2.0 1610 2430 8.83

5 91 9 2.0 1460 2750 5.70

6 87 13 2.5 1250 3560 3.77

7 89.7 10 0.2 TaC

0.1 TiC

3.5 1250 2780 5.07

8 87 13 4.0 1090 3100 3.77

a TRS = Transverse rupture strength

308
The results for the supply voltage V in figure 2 show a linear decrease with faster infeed velocity up to a certain point, beyond which the drop in supply voltage becomes steeper.

Since each of these curves is for a fixed current, this means that the electrical resistance

across the grinding zone follows a similar relationship. The discontinuity at the start of

the steep voltage drop occurs at a faster infeed velocity with a bigger current density.

0 02 0-4 0.6 0.8 1.0 1.2 0 02 0.4 0.6 08 1.0 1.2 vj,(mm/min) Vj,(mm/min)

Figure 2. Supply voltage versus infeed velocity at various current densities for

Materials 4 and 6.

Both the normal and tangential force components in figures 3 and 4 increase with faster

infeed velocity. Discontinuities can be seen in the results in figure 4 for which cor- respond to those discontinuities in figure 2. Since the tangential force component F. is

proportional to the mechanical grinding power requirements, the decrease in with bigger

current densities in figure 4 indicates the effectiveness of electrochemical assistance in

reducing the mechanical grinding power. By contrast, the normal force component F in figure

3 is much less affected by current density during ECG and in some cases exceeds the normal

force for purely mechanical grinding (i=o), probably because of pressure from gasses evolved

from the electrochemical reactions. Extrapolation of the force components to zero in figures

3 and 4 generally corresponds to a positive infeed velocity except with i=o.

One measure of grinding behavior for conventional mechanical grinding is the specific

grinding energy input per unit volume of material removal. For ECG, an analogous parameter

is the specific mechanical energy u , which is the mechanical energy input per unit volume

of material removal

m hi

"i

A (8)

where the numerator is the mechanical power input and the denominator is the volumetric re- moval rate. The specific mechanical energy results, corresponding to those in figures 2-4,

are presented in figure 5. Here it can be seen that for mechanical grinding (i=o), the

specific mechanical energy decreases as the removal rate is raised by increasing v., which

is typical for mechanical grinding. With ECG, the opposite trend is observed, the specific

mechanical energy being very small at slow infeed velocities and increasing at faster velocities. It is evident from these results that the ECG action is relatively more effective

at slower infeed rates. Discontinuities in the results for u are-.observed which correspond

to those seen in figures 2 and 4. At infeed velocities beyona the discontinuities, the specific mechanical energy increases at a faster rate coming closer to the results for mechanical

grinding (i=o).

309
MATERIAL 4

vj, (mm/min) Vj j

( mm/min )

Figure 3. Normal force comDonent versus infeed velocity at various current densities

for Materials 4 and 6.

Figure 4. Tangential force component versus infeed velocity at various current densities

for Materials 4 and 6.

310
vj , (mm/min) Vj,(mm/min)

Figure 5. Specific mechanical energy versus infeed velocity at various current densities

for Materials 4 and 6.

5. Surface Damage

Weakening of the surface due to ECG was determined from the transient grinding behavior

after switching off the electrical current, as described above. An example of the force

transient obtained after switching off the current is shown in figure 6a for one carbide at

two current densities. Both the normal and tangential force components can be seen to in- crease with time after the start of the transient. Taking into account the elastic deflection

of the grinding system as described above, the actual infeed velocity v during the transient

and the actual depth removed t at any time t during the transient were

T computed. The re- sults corresponding to the transient in figure 6a are shown in figure 6b as a plot of v versus t. Here it can be seen that the actual infeed velocity drops abruptly from its initial

steady state value due to increasing elastic deflection from the growing normal force. This

is followed by a subsequent increase in v towards its steady state value v.. The tangential

grinding force results from figure 6a are

T also included in figure 6b.

Surface damage caused by ECG is evaluated in terms of the specific grinding energy of the

material in the surface layer. From results such as given in figure 6b, the specific grinding

energy at depth t can be calculated as:

v A

T

(9)

where the numerator is the grinding power and the denominator is the instantaneous removal

rate. The relative weakness is defined as the ratio <|> of the specific energy in the surface

311
(N) 901-32

•0.11 i •0J7 22P

•A- mm1

• i-CU7

MATERIAL 8

20 24 28 32

T(sec)

(a)

MATERIAL 8

i.0.20 A— mm*

(b)

Figure 6. Results of grinding transient: (a) force components versus time after switching

off current; (b) actual infeed velocity and tangential force component versus depth progressed.

layer after ECG to the normal specific energy for mechanical grinding of the same material

with the same wheel in a seperate series of tests. Damage to the surface by selective etching of cobalt should weaken the material and reduce its specific grinding energy leading to

values of <j> less than unity. The smaller the value of <f>, the greater the damage. Results

obtained in this way for <}> are shown in figure 7 which includes those for the grinding transient in figure 6. It can be seen in figure 7a that <{> increases from a very small value at

(a)

60 100

t (/xm)

.6-

.2-

MATERIAL 7

MATERIAL 8

20 40

±

Figure 7. Weakening factor <J> versus depth below finished surface:

60

(b)

(a) v.

80 100

t (/xm)

0. 30mm/min

;

(b) v. = 0.48 mm/min with 1=0.20 A/mm2 for Material 7 and 1=0. 17 A/mm2 for

Material 8.

312
the surface to unity at a depth of approximately 50-100um, which is the thickness of the sur- face layer affected by ECG. With the larger current density, the weakness factor at the surface is smaller and the material is damaged to a greater depth. Similar results were obtained

with other cemented carbides. With a faster infeed velocity in figure 7b, there is much less

weakening at the surface and the ECG affected depth is shallower.

6. Discussion

The electrochemical analysis for WC-Co cermets shows that the initial etching velocity

into the cobalt phase is much faster than into the WC phase. It is proposed that selective

etching of the cobalt in this way v/eakens the cermet and facilitates mechanical removal

during ECG.

In order for this selective etching phenomenon to occur, it is apparent that the infeed

velocity of the ECG process must be slower than the initial etching velocity into the cobalt.

If the infeed velocity exceeds this etching velocity, the electrochemical assistance should

be much less effective. It may therefore be expected that the steep rise in mechanical power

requirements beyond the discontinuties , such as seen in figures 4 and 5, may be due to the

infeed velocity exceeding the initial etching velocity into the cobalt. In order to test

this hypothesis, the infeed velocity v. at the discontinuity was plotted versus the initial

etching velocity v~ given by eq. (4a) for all eight workpiece materials with various current

densities. The resalts in figure 8 show these two velocities to be almost identical.

The particular effectiveness of ECG on UC-Co cemented carbides can be appreciated by

cE

1.4 O 1 0 5 A 2 X 6

V 3 + 7

1.2 D 4 * 8

MATERIALS:

Vj - VCo V

Figure 8. Infeed velocity at discontinuity

versus initial etching velocity into cobalt.

1.0 1.2 1.4

Vco f (mm/min)

applying a conventional type of analysis to the present results [4-6]. The conventional ECG

analysis is based upon the assumption that electrochemical and mechanical removal proceed independently, so the total volumetric removal rate is given by:

Z = Ze + m

(10)

According to this assumption, the relationship between the mechanical power (or force) and

mechanical removal rate is the same for both ECG and mechanical grinding with i=o. Therefore,

for a given relationship between the power and the removal rate for mechanical grinding, the

mechanical contribution Z during ECG can be estimated by measuring the mechanical power.

With this result, the electrochemical efficiency during ECG can be obtained as:

313
(11)

where the numerator is the electrochemical removal rate obtained by substracting the estimated

Zm

from the tota1 removal rate, and the denominator is the theoretical electrochemical rate

according to Faraday's Laws (eqs. 2 and 3). Representative results for the efficiency are shown in figure 9 which were obtained from the results in figures 4 and 5 with a=a*. In all

cases, the computed efficiencies were found to exceed the theoretical limit n = 1.

MATERIAL 4

/ / A \

'.II /// V

\ \N

o * i=i7£!!!E

ri mm*

.8

_L

1.0 1.2 v;,(mm/min)

4.0

3.0

2.0

1.0-

0

0

MATERIAL 6

mm /

r

r * A.UL

J A - f O /

;.I? \ v

\ r

i=.23 i=.26

8 1.0 1.2 vj,(mm/min)

Figure 9. Calculated electrochemical efficiency versus infeed velocity at various current

densities for Materials 4 and 6.

Such large calculated values for n are apparently due to selective electrochemical etching

of cobalt which weakens the material, thereby reducing the mechanical grinding resistance of

the cermets. The degree of weakening obtained in this way can be estimated by applying the

conventional ECG analysis setting n = 1. The value of Z is now obtained in a different way

by subtracting Z , given by eqs. (2) arfd (3) with a=a*,

mfrom the total Z. The specific en- ergy for only mechanical removal during ECG can be obtained as:

where the numerator is the Dower and the denominator is the mechanical removal rate. The

effectiveness of electrochemical action in reducing the mechanical grinding resistance can

be seen by comparing the value of u with specific energy u for grinding with i = 0,. the com- parison being made at the same mechanical removal rate. The specific energy u is the same as the specific mechanical energy um

when i = Ch Representative results for u/u are shown in figure 10 corresDonding to those in figure 9. It can be seen that u/u is much less

than unity for those materials, especially with slower infeed velocities and larger cur- rent densities. Beyond the discontinuities, which correspond to those previously discussed,

this ratio increases at a much faster rate toward unity as selective etching should no longer occur.

314
Figure 10. Ratio of specific grinding energy durina ECG to specific grinding energy for i=0 versus infeed velocity at various current densities for Materials 4 and 6.

In the surface damage experiments, weakening in the finished surface layer was expressed

by the parameter Weakening of the material during ECG can be expressed by the ratio u/u

Both $ and u/u0 measure the same thing, namely, the reduction in mechanical grinding resistance due to selective electrochemical etching. Therefore, the value of <f> at the surface

(t=o) should be eguivalent to u/u0 for the same ECG conditions, as the state of the material

below the surface immediately after switching off the electrical current should be the same

as that of the material being mechanically ground away during ECG. Indeed, values of <$> at

t=o were found to be very close to corresponding values of u/uQ

.

In summary, the effectiveness of ECG as a material removal process for WC-Co cermets

can be attributed to selective electrochemical etching of the cobalt binder. However, this

same effect is responsible for poor surface quality. One approach to practically utilize

the ECG process would be to begin with a high current density for rough machining followed'

by a decrease in current density, even to zero, for final grinding to remove the damaged

layer and to achieve the reauired surface integrity. The results of the present study provide a basis for designing appropriate ECG cycles.

References

[1] Levinger, R. and Malkin, S., Electrochemical grinding of WC-Co cemented carbides,

ASME Paper Mo. 78/WA/PR0D.-26 , J. Eng. Ind. , Trans. ASME, to appear.

[2] Malkin, S. and Levinger, R. , Quantitative characterization of surface damage to WCCobalt composites by electrochemical grinding, Wear 51^, [1]> 157-157 (1978).

[3] Kreimer, G.S., Strength of Hard Alloys , p. 2 (Consultants Bureau, New York, 1968).

[4] Backer, W.R. and Dahlin, R.A., The effects of electrolytic assistance in peripheral

grinding , ASME, Paper No. 60 WA-4, 1960.

[5] Cole, R.R., An experimental investigation of the electrolytic grinding process , J. Eng. Ind., Trans. ASME, 83, 194-201 (1961).

[6] Kaczmarek, J. and Zachwieja, T. , Investigation of the material removal rate by

electrochemical grinding of cutting tool materials, Int. J. Mach. Tool Des. Res.6_,

[1], 1-13 (1966).

315

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

A NEW ULTRA—FINE POLISHING TECHNIQUE FOR SAPPHIRE

Tadaaki Sugita, Seiho Kawakami and Osamu Imanaka

Department of Precision Engineering, Kanazawa University

Kodatsuno, Kanazawa, Ishikawa 920, Japan

A new method of ultra-fine polishing for sapphire oriented with the polishing

surface parallel to the (0001) plane is proposed, and the nature of the polished

surface is analyzed. Experimental polishing is performed in water using a pin-ondisk machine with a tool disk made of sintered alumina. The sapphire is remarkably

polished by friction action causing removal of the surface asperities; the polishing

rate is significantly decreased as the smoothed surface is formed. Roughness of

the polished surface is less than Rmax of 200 A. Behavior of friction is very

stable in comparison with specimen worked in air. A high crystal lographic integrity

of the surface polished in water is confirmed from electron diffraction images.

Analyses of the surface with x-ray photoelectron and ion microprobe mass spectrometers indicate the presence of a very thin hydrated layer which acts as an anti-friction agent preventing the damage of the nonhydrated substrate.

Key Words: Crystal lographic integrity; electron diffraction images; hydrated

layer; ion microprobe mass spectrometer; roughness; sapphire; ultra-fine polishing;

water; x-ray photoelectron spectrometer.

1. Introduction

Exacting requirements must be considered in the machining of ceramic materials which are

to be used as high performance elements in the field of electronics, especially in the production

of highly integrated semiconductor circuits. These are high geometrical accuracy, including an ultra-fine surface finish; crystal lographic perfection of the worked surface without any

damage such as lattice disorder; scratches and cracks; and freedom from chemical contamination

which could be caused by the tool material, machining environment and/or chemical etchant.

In general, most ceramic materials have been mechanically or chemically polished using a fine

hard abrasive, like diamond, or chemical etchants. It is inevitable, however, that

mechanical polishing is apt to leave scratches and damages on the polished surface and that

chemical contamination occurs in chemical polishing.

In order to obtain such extremely superior surfaces, elastic emission machining using

fine hard powder particles [1,2]*, and mechanochemical polishing with chemically reactive

powder particles [3,4] have been developed. These techniques should be considered as novel

ultra-fine machining methods. In the elastic emission machining, the intent of the collisions

of the fine powder particles with the material surface is to make the size of the stress

field smaller than the distance between the pre-existing defects and to verify the feasibility

of elastic fracture of the order of atomic size. In mechanochemical polishing, the machining

mechanism is based on mechanochemical reactions at real contact points between the work piece

and abrasive powder which are softer than the work piece, and easy removal of the chemical

reaction compounds from the interface without any surface scratches and damages. These

machining methods are expected to be useful for finishing hard electronic crystals.

On the other hand, Imanaka et al. [5,6] have newly developed a hydration polishing

technique which is based on the chemical characteristics of a hydration phenomenon occurring

by the action of water vapor molecules on the surface of the crystal to be polished. This

method uses neither polishing powders nor chemical etchants. Hydrates formed on the activated

surface are rubbed off by the friction of the polisher. If the hydration occurs only at the

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

317
true points of contact between the work piece and the polisher, and if only the hydrates are removed from the work piece surface, this polishing method may fulfill the above mentioned

exacting and severe requirements for a polished surface.

A new polishing technique to be presented in this paper also utilizes hydration polishing

It is based on the concept that the polished surface would be free from chemical contamination

by the polisher material if the same material used as the work piece is selected for the

polisher. Also, this polishing technique which removeds only the hydrates formed at the real

points of contact may be effective in avoiding various surface damages and making an ultrafine polishing surface because the hydrates are relatively easily removed.

2. Experimental Procedure

Polishing equipment is shown in figure 1. 8

In this figure, the (0001) plane of a A1 2 0 3 single crystal (sapphire) is rubbed on a 80 mm diameter rotary disk made of high

purity sintered alumina2

. The sapphire

specimen is held at the end of a reciprocating lever with a normal load applied on top of the specimen. The reciprocating

lever is attached to a square sliding shaft

with a horizontal connecting pin which allows freedom for vertical movement. The

rotary disk rotates at 140 r.p.m and the

reciprocating lever moves back and forth

at 43 cycles per min. Water is always

directed on the rotary disk which is in a water tank. Rubbing tests in air are also

performed with this machine without using

water and keeping the surface of the disk

to be dry. Friction forces due to the test

environment are measured with strain gauges

mounted on the reciprocating lever which

remains stationary during such measurements.

The sapphire specimen had a diameter

of 3 mm and was oriented with the (0001)

basal plane parallel to the polishing face.

It was prepared for the polishing test by

grinding to a surface roughness of 1 ym (maximum height of irregularities). The removed

volume and the removal rate by polishing was obtained from a change of diameter at the top

of another conical specimen, which was only prepared for this test, with an apical angle of

120°. The variable contact pressure in the process of polishing was calculated from a change

of the contact area as an average value.

The profile of the polished surface was obtained with a profilometer. The crystalline

structure of the polished surface layer was estimated by comparing it with that of the fracture surface using electron diffraction equipment with an accelerating voltage of 100 KV.

XPS data which provide information in regard to chemical shifts due to a property change of

the material or to chemical contamination of the surface were obtained using an ESCA 650B

photoelectron spectrometer, Du Pont, employing Mg Ka exciting radiation. Binding energies were corrected to Ci S = 285 eV for the test samples. The instrument was calibrated by the

Au4#7/2 Deak with a binding energy of 84 eV„ The measured binding energies are considered

to be accurate to + 0.3 eV. Elements and compounds in the extreme surface layer of one or

two atomic order were detected with an ion microprobe mass spectrometer (Hitachi IMA-2

system, Japan). From these analyses the contaminant elements and compounds on the surface

are further identified.

2 Si<30, Mg<10, Na<5, Ca<2, Fe<10, Cu<l ppm

Figure 1. Polishing machine-. 1, rotary tool

disk(sintered alumina); 2, sapphire specimen;

3, specimen holder; 4, weight; 5, reciprocating

lever; 6, sliding shaft; 7, slide bearing; 8,

eccentric lever; 9, eccentric cam; 10, water

tank; 11, inlet pipe of water; 12, outlet pipe

of water.

318
a.

3. Results and Discussion

3.1 Polishing characteristics

As well known, adhesion occurs easily at the contact interface of two sliding solids

when the material in the solids is the same. Therefore, in this study, it is necessary to

take this behavior into account in the polishing experiments because the polisher and the

work piece have the same structure and com- position in order to avoid chemical contamination from the tool material „ Figure 2 shows

typical friction behaviors in the test environments of water and air. Although the friction

obtained by sliding in water is very stable,

the friction force in air on a dry disk is

changeable and large because of adhesion,,

The crystalline structure of the surface

rubbed in air, as described in following

section, becomes disordered . The stable

friction in water may be caused by the

presence of a water film or hydrate; in the

environment of water this friction may be due

to shearing the layered aluminum hydroxides

at the interface- Thus, the hydrates play the

role of anti -friction agents between the solids

in moving contact and, moreover, are swept away

by the mechanical forces developed in polishing.

T IM E

3

a.

Z3O

MM

TIM E

(a) (b)

Figure 3 shows an exponential relationship between the contact pressure calculated

from the change of contact area and the rate

of removal. This result was derived from the

changes of the removed volume and the average

contact pressure during rubbing in water under

a load of 1 Kg. The contact pressure is con- tinuously decreased with the increase of wear

because the conical specimen with an apical

angle of 120° was used in order to measure

the removed volume. Sliding distances in Km are shown in parentheses on the figure. It

is evident that the removal rate of specimen

depends on the change of contact pressure.

From the above result, it is suggested that

in the polishing test, the polishing rate is

relatively high until the asperities formed

in the preliminary machining of the specimen

are leveled. Also, because of a corresponding

increase of contact area, a decrease of the

local normal stress must occur which then

remains constant. It is expected that the

presence of the hydroxide film on the polished

surface, because of the continually added

water, protects the surface from the damages

that would be introduced by the polisher under

dry conditions.

Figure 3. Relationship between rate of

removal and contact pressure in rubbing

test using a conical specimen with an apical

angle of 120° under a load of 1 Kg.

Sliding distances in Km are shown in parentheses ,

Figure 2. Behavior of friction (a) in water

and (b) in air, under a noma! load of 0.75

Kg.

I 2 4 6 8 10

AVERAGE CONTACT PRESSURE (Kg/mm)

319
3.2 Analysis of polished surfaces

The profile of a surface polished in water under a normal load of 0.5 Kg is shown

in figure 4. The polishing roughness in water is less than Rmax of 200 A. In this c§se,

if the measured length is limited to 5ym, the roughness is approximately Rmax of 100 A or

less. This result suggests that the thickness of the layered chemical reaction products

at the contact interface is less than 100A. However, a wide area of the polished surface

is difficult to realize because a very heavy load has to be applied to the machine for

polishing. This problem will be solved by an improvement of the machine system. The

roughness of surface polished in air is about 400 & because of active adhesion.

Figure 5 shows electron diffraction images for three sapphire surfaces prepared by

fracturing and polishing in water and air. The penetration depth of electrons is about

50 & under the accelerated voltage of 100 KV. Kikuchi lines are clearly observed in the

diffraction patterns for the fractured surface and the surface polished in water. This

crystallograph evaluation indicates that surface damage, like lattice disorder and/or

plastic deformation, were not introduced into the surface polished in water.

(a) (b) (c)

Figure 5„ Electron diffraction images: (a) fractured surface, (b) polished surface

in water, (c) rubbed surface in air.

Figure 6 shows wide scanned x-ray photoelectron spectra (XPS) for the three sapphire

samples. The absolute intensities cannot be used for comparison because the amount of hydrocarbon deposited on the surface varied from sample to sample. The carbon(Cls ) electron

energy of 285 eV is used to calibrate the spectra. C ls is probably due to the hydrocarbon

derived from the vacuum system in the x-ray photoelectron spectrometer. The test environment

can easily introduce Na.

Figure 7 shows changes in the photoelectron spectra of Al

2

P for the fractured surface

and the surface polished in water before and after argon ion bombardment: 1 is before ion

bombardment, and 2, 3, 4 are after ion bombardment for 10, 30 and 50 seconds, respectively.

The corrected Al

2

P binding energies for the treated sapphire samples and commercial aluminum hydrates are given in table 1. The table data indicate that these specimens are similar suggesting that the exposed fracture surface becomes hydrated and t^at the hydration

layer is very thin since the rate of ion bombardment is approximately lOA/min. The chemical shift of about 1 eV in the specimen polished in air can be attributed to a disordered

lattice near the surface which is assumed to exist according to the electron diffraction

image shown in figure 6(c). The binding energy (75.1 eV) for Al 2p in the fractured

sample is different from that (76.3 eV) measured by lindsay et al. This difference may be due to the use of different calibration lines, such as the C ls lines

320
700 600 900 400 900 200 100

BINDING ENERGY (eV)

Figure 6„ Wide scanned photoel ectron

spectra for the three sapphire samples

used in figure 5: (a) fractured surface,

(b) surface polished in water, (c)

surface rubbed in air.

Table 1. Binding energies of Al 2p

electrons obtained from experimental

surfaces of sapphire, and aluminum

hydroxides

.specimen tsincnng energy^ev;

11 Fractured surface 2

3

4

74 7 74 0 6

74„8

75 0 1

1 Polished surface, 2

in water, (in air) 3

4

74.3 (73.6)

74.8 (74.7)

74.8 (74.1)

75.0 (74.0)

a-Al(0H) 3 74.5

a-AlO(OH) 74.4

•0 75 70

BINDING ENERGY

(Q)

80 75 70

BINDING ENERGY (eV)

(b)

Figure J. Changes in the photoel ectron

spectra of Al 2p for (a) fractured surface,

and (b) the surface polished in water,

before and after argon ion bombardment.

1 is before ion bombardment. 2, 3, 4 are after ion bombardment for 10, 30 and

50 seconds, respectively.

in this test and the Au 4 f line in their test, or to different test materials. It has been

shown that the Cis electron line from a hydrocarbon contamination may shift as a function

of the type of sample being examined [7], but the data in this analysis are comparable and

can be used effectively^to compare the test specimens. The hydroxides a-Al (OH) 3 and a-AlO(OH)

have layered structures with hydrogen bonds joining the layers with a difference in the

degree of hydrogen bonding. From the data in table 1 it appears that the indicated hydrates

on the surfaces are thin and removed in about 30 seconds of argon ion bombardment. Their

structures could not be determined from the photoel ectron spectra.

321
The ion microprobe mass spectrometer( IMA) is useful for analyzing the chemical composition of a surface., Figure 8 is an IMA spectrum of a surface polished in water. While

argon gas ions are impinging on the surface at an extremely low current density and sputtering off surface atoms and molecules, the ionized fractions „are analyzed with a mass

spectrometer,, The rate of removal by ion bombardment is 0,2 A/s„ This spectrum was obtained

within 2 min, after the analysis was begun. It indicates that the detected elements and

compound units are within about 24 A of the surface layer. It is possible that the source

of the chemical contaminants K, Na and Ca was the water

3 used in polishing. However, the

contamination markedly decreased 10 min,

after the beginning of the analysis. The

amount of aluminum hydroxides may be small

as already indicated, but their existence

is supported by the presence of the species

A1H+ and A10H+,

4. Conclusions

An ultra-fine polishing technique for

sapphire with the use of water^developed,

and the nature of the polished surface was analyzed with various surface analyses techniques. The experimental results are sum- marized as follows:

(1) The polishing mechanism is due to hydration of the a-Al 2 0 3 surfaces and the

mechanical removal of these hydrates

without damaging the bulk structure

resulting in the removal of the asperities which formed the initial contacts

at the interface. The roughness of the

poli§hed surface is less than Rmax of

200 A,

(2) An undisturbed crystalline structure in

the water polished surface layer and a damaged structure in the dry polished

surface layer are confirmed by electron

diffraction images.

(3) Analyses of the surface of sapphire

polished with water with x-ray photoelectron and ion microprobe mass spectrometers support the presence of a hydrated layer on the surface in the

range of about 5 - 24 A„

20 30 40 50

m/e

100

Figure 8. A spectrum for sapphire surface

polished in water obtained by ion microprobe analysis; sputtering time, 2 min.

Acknowl edgments

The authors express their appreciation to Professor J.A,Pask(L)niv. of Calif.) for

helpful discussions and advice and to K,Nishikawa for constructing the apparatus. The

authors also are grateful to Dr. Y.Namba(Osaka Univ, Japan) for the ion microprobe analysis and to T, Wakabayashi for the x-ray photoelectron analyses.

3 Si , Mg, Na, Ca and K were detected from the deposit after vaporizing the water.

322
References

[1] Tsuwa.H.and Aketa^Y,, Elastic emission maching as an ultra-fine finishing method, Proc.

1st, Inter. Conf. Prod. Engrq. (symposium) , pp.33, Tokyo, August, 1974.

[2] Mori ,Y„ Ikawa,N„Okuda,T. and Yamagata,K„ Numerical ly controlled elastic emission

maching, Tech. Repts. Osaka Univ. 26, [1301], 283 (1976).

[3] Yasunaga,N„Obara,A.and Imanaka,0„Mechanochemical polishing of single crystals with

soft powders, Proc. 1st. Inter. Conf. Prod. Engrg,, pp.32, Tokyo, August, 1974.

[4] NambajJ. and Tsuwa,H„ Ul tra-fine finishing of sapphire single crystal, Proc. 27 General

Assembly of CIRP, pp.1, New Delhi, August, 1977.

[5] Okutomi ,M.and lmanaka,0„ Hydration polishing of sapphire— A novel finishing technique

using hydration phenomenon--, Bull. Japan Soc. Prec. Engrg., 10, [4], 156 (1976).

[6] Okutomi ,M. and lmanaka,0„ A novel finishing technique using hydration phenomenon -- Hydration polishing of sapphire --, Proc. 27 General Assembly CIRP, pp. 79, New Delhi,

August, 1977.

[7] Lindsay, J. R, Rose, H. J., Swartz, Jr., W.E., Watt, Jr., P.H. and Rayburn, K.A., X-ray

photoelectron spectra of aluminum oxides: Structural effects on the "Chemical shift",

Appl. Spectroscopy, 27, [1], 1 (1973).

323

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

FINE POLISHING OF SAPPHIRE CRYSTALS

M. IKEDA, A. YAMADA, Y. KOKAJI, S. KIMURA

Electrotechnical Laboratory, MITI

Mukodaimachi, Tanashi, Tokyo 188, Japan

A new method for polishing sapphire crystals has been developed which is based

on a chemical reaction at the contact area between surfaces of solids to be polished

and a polishing disc. Special features of this method are as follows: (1) A scratchfree and smooth surface is produced on the basis of reaction between silicon oxide

and alumina. (2) The worked layer left in a subsurface polished is so thin as not to

be detected with a Berg-Barrett method. (3) Use of a rigid polishing disc results in

a highly flat finished surface. (4) Since no abrasive particles are used, a clean

working environment is maintained.

Key Words: Damaged layer; fine polishing; sapphire; silicon; solid state reaction; wear debris.

1. Introduction

Sapphire single crystal is one of important inorganic electronic materials, as substrates of semiconductor devices. Generally, sufaces of these materials to be used are polished with high geometrical accuracy without residual damage and strain. Many polishing

techniques which produce the suitable surface have been studied mainly for semiconductor

crystals

These techniques are concerned with mechanical polishing by using fine, hard abrasives

such as diamond, silicon carbide, alumina and plastic or elastic polishing sheets, which

leave damaged layers on the finished surfaces. So a final treatments i.e. chemical or electrochemical etching are needed in order to remove such a worked layer at the expense

of decreasing geometrical flatness of the surfaces.

The present polishing method is based on the chemical reactions at real contact areas between two surfaces of solids, that is, sapphire crystals worked and a polishing disc.

This method can produce fine surfaces which are free of scratches and any damage and have

high dimensional accuracy. This paper describes experimental results of the new surface

finishing method of sapphire crystal and evaluations of polished surfaces.

2. Background of the polishing method

The concept of this method was derived from the fundamental experiment observing the

wearing properties of alumina abrasives against silicon crystals. When a sapphire cone scratches a silicon crystal surface, it was found that the tip of the cone was worn away,

and chemical reaction was proved with the X-ray diffraction analysis of the wear debris in

sliding tests.

2.1. Scratching Test

Silicon-sapphire contact was examined by scratching test which was performed with a light load scratching equipment specially designed (fig. 1), on a clean surface of silicon

single crystal wafer (111) with a sapphire cone whose tip radius was about 5 vm. Then

tracks on the wafer and the tip of the cone were observed with a SEM.

325
Fig. #1. Scratching apparatus.

Intersecting tracks shown in figure 2 have been formed by scratching in two directions

at a speed of 30 mm/min under 7 gram weight load. A series of numerals noted on the diagram in the figure present the order of scratching. Inside of tracks are pellicular and

squeezing out scarcely occurs, unlike the tracks scratched with a diamond stylus for ex- ample. Broadening of track width with the advancement of scratching length means wear of

the sapphire cone. The tip of the cone worn away by scratching test extending through 140

mm (fig. 2) has been flattened to a relatively smooth plane leaving some shallow grooves

oriented in scratching direction.

1 2 3 4 5 6

Fig. #2. Tracks on Si crystal surface scratched with a sapphire cone.

326
Such wear of sapphire seems not to be anticipated because of hardness difference between sapphire and silicon. The mean contact pressure at the interface between a silicon

surface and a sapphire cone is estimated to be several million kg/cm^, and the temperature

at the contact point is considered reasonably high enough because of adiabatic compression

and frictional heat occurrence. Thus, the condition is considered to be enough for solid

state reaction between two materials.

Anyhow, it is apparently expected that sapphire crystals are able to be polished

by silicon without any abrasives.

2.2 Sliding Test

In order to obtain some explanation about the wear of the sapphire cone mentioned

above, a sliding test was carried out.

2.2.1 Experimental Set-up

A sapphire specimen (the apparent contacting area 5x4 mm ) was slid on the end face of

a silicon crystal rod (diameter 35 mm, length 70 mm) mounted to a chuck of a lathe. Mean

sliding speed was about 7790 mm/min at which chipping of both materials didn't occur. Loads of 250, 600, 1000 g were selected, and they were set dynamically by compression of a coil spring.

The orientation of the silicon rod was [III] and the sapphire specimen surface was about 5° from (0001), which is known to wear most easily (fig. 3). Structure of the wear

debris, collected on a sheet of paper, was investigated by X-ray diffraction analysis.

Initial feature. After scratching.

Fig. #3. Tip of sapphire cone.

2.2.2 An analysis of the debris

Figure 4 shows a part of an X-ray diffraction chart (Cu-Ka

) of the debris produced

under 1000 g load. Within the angle of diffraction from 28° to 29°, apparently two peaks

are seen. The right one is known as a diffraction peak from a silicon (111) surface, but

the left one is a diffraction neither from silicon nor sapphire. It is considered that tiny

silicon chips fractured during sliding gave the diffraction peak from silicon [III], since

both the occurrence of fragments and the height of the diffraction peak decreases with decreasing load. .

The diffraction curve plotted by a scaler method in the angle from 28° to 29° is shown

in figure 5. In this case, the debris produced under 600 g was used. Intensity of X-ray

diffracted is extremely weak, so the measured values fluctuate a little. No peak is observed at the angle of diffraction from silicon (111), but a peak is seen at the lower diffraction angle which is considered to be diffraction from Kyanite (Al20g.Si02) . This means

that the chemical reaction occurs between Si02 covering the silicon crystal surface and

AI2O3, and another material; that is, Kyanite is produced during a sliding experiment.

Kyanite is generally known as an intermediate reaction product between Si02 and AI2O3 under

relatively gentle conditions.

327
28<

I 30°

Fig. #4.

32° 34°

Angle of diffraction

X-ray diffraction chart of wear debris

36c 38(

100 -

>>

-p

•i

inc

a>

4->

Angle of diffraction

Fig. #5. X-ray diffraction curve of wear debris.

The X-ray diffraction chart also suggests that the amount of crystalline reaction prod

uct in the wear debris is extremely small; the major portion of wear debris appears to be

amorphous

.

328
3. Polishing Experiment

3.1 Experimental Set-up

Two kinds of polishing discs using silicon were provided for polishing experiments.

One of them has been formed as follows, silicon crystals sliced in random orientations were plastered like a mosaic on a steel plate and ground into a flat surface. The other is a similar disc of bonded silicon powder which was crushed from single crystals into about 80

mesh. The disc is set on a polishing machine of particular design and is rotated (fig. 6).

Three sapphire speciments (10x10x1.5 mm) mounted to a holder are pressed to the disc and

loaded with weight; the holder is forced to rotate on the polishing disc through a synthetic rubber belt driven by a motor. Water (5 cm^/min) was poured into the disc in order to

eliminate the refuse formed during polishing. Surfaces of sapphire specimens to be polished were sorted into two classes, one of which was lapped with GC #2000 abrasives and

another was ground with a wheel composed of diamond pellets of 1200 mesh. The experimental

conditions are described in figure 7.

Fig. #6. Polishing apparatus.

j Specimen

i Sapphire crystal (1102) 10x10x1.5TM

, Surface to be polished Lapped (GC#2000), Ground with diamond pellets (4utn)

Polishing disc Si crystal sliced, Si powder

: Polishing time K10 hours

Polishing pressure 150^800 9 /cm2

j Polishing fluid

3 Water, 2^5cm /min

i Rotating speed Disc; lOOrpm, Specimens; I8rpm

Fig. 7. Experimental conditions.

329
3.2 Experimental Results and Discussions

3.2.1 Stock Removal

Stock removal of sapphire crystal increases naturally with polishing time as seen in

figure 8. The progress rate of the removal is relatively large at early stage of the polishing and it decreases gradually with the lapse of time. It is supposed that the descent

of the removal rate results from reduction of the activity at contacting interface with ad- vance of solid state reaction. It must be also taken into consideration that the polishing

load at the interface is mitigated owing to increase of real contact area.

0.4

load(kg/cm2),preceding process

diamond pellets

GC *2000

: 0-6

: 0-1

O : 0.6 •

:

0-1

Fig. 8. Stock removal variation with time.

4 6

Time (hour)

0.2

<NE

I 0-1

>

o

B

0)

S-o

o

4-> X

•: Si crystal

o: Si powder

Fig. 9. Pressure dependence of stock removal

> 200 400 600 800 1000

Polishing pressure (g/cm2

)

The difference of stock removal among data from different conditions confirms this effect, that is, the surface roughness of specimens ground previously with diamond pellets is

about five times as large as those lapped with GC#2000 abrasives. Though the influence of

polishing pressure is comparatively large for specimens processed with diamond pellets, it

is not so large for the specimens lapped with GC abrasives. This difference is due to sur330
face properties depending on the processes. When the disc made of sliced silicon crystals

is used, the stock removal of the sapphire crystal increases in proportion to the polishing

pressure below about 500 g/cm^ and the influence of the pressure lessens gradually at

higher pressure (fig. 9). In the case of polishing with the disc formed of silicon powder,

stock removal is very small in the lower range of polishing pressure. Under the conditions

that the pressure exceeds about 200 g/cm , the influence of the pressure appears to be similar to the mosaic disc. Therefore, it is possible to polish sapphire crystals with any

form of silicon. The reason why the stock removal concerning the powder disc is small at

low pressure may be that the binder prevents the silicon particles from sufficiently con- tacting the specimen surfaces.

3.2.2 Characteristics of Polished Surfaces

The surface roughness of sapphire specimens polished for 10 hours under a pressure of

600 g/cm^ on the sliced-silicon disc was at most, 0.02 urn R^x- On the polished surface

(fig. 10), there are relatively few scratch marks, in contrast to surfaces polished with

diamond paste. This contrast shows that the polishing mechanism of the present method is

quite different from that of the conventional one with diamond paste that is based on mechanically scraping off. This inference is supported by X-ray Berg-Barrett photographs

shown in figure 11 where serious defects are not observed on the polished surface, that is

there is little strain present in the crystallographic lattice beneath the surface of the

specimen.

Fig. #11 Berg-Barrett photograph of sapphire crystal polished;

(a) with pellon cloth and diamond paste; (b) by present method.

331
Geometrical flatness of the surface polished by present method is excellent and surface curvature at the edge of samples is very small relative to that of samples polished

with diamond paste and pellon cloth, as seen in figure 12. The ESCA spectra of polished

surfaces prove that any reaction product or diffused silicon are not a bit left (fig 13)

by comparison with an etched surface.

Fig. #12. Interferometric pattern of specimen fringe; (a) with pellon cloth and diamond;

(b) by present method.

eV

Fig. #13. ESCA spectra on sapphire crystal surface.

4. Conclusions

Experiments show that sapphire crystals are able to be polished with only silicon and

water and the method produces sufficiently good surfaces.

(1) By this method which is fundamentally based on the solid state reaction between two

surfaces of SiC^ and A^Og, a scratch- free and smooth surface (R^x 0.02 um) is produced.

332
(2) The worked layer left in the subsurface is hardly detected by X-ray Berg-Barrett

method.

(3) The use of rigid polishing discs produces a flat surface without curvature at the

edge of specimens.

(4) No employment of any abrasives keeps the working environment clean and is convenient

to automatic polishing.

In order to apply this method to practical use several problems should be solved, that

is, to find out the most suitable surface state to be polished, to keep up the activity of

the polishing disc, to prevent the occurrence of cracks at the edge of surface during

machining and to increase the stock removal.

Authors wish to thank Mr. M. Okutomi for help of surface observation by ESCA.

References

Main articles presented recently,

Ikeda, M. , et al., Polishing of silicon crystals with an elastic wheel, Annals of CIRP 23

(1) 103-104 (1974).

Nhar van Huyne, et al., Studies on electrolytic buffing on single crystal silicon,

Seimitsu-kikai, 43 (6) 650-655 (1977).

Miyake, S.„ et al., Ferrites precision grinding, ibid, 43 (12) 1415-1420 (1977).

Nhar van Huyne, et al. Studies on electrolytic buffing of ferrite single crystal, ibid,

43 (4) 421-426 (1977).

Mori, Y., et al., Elastic emission machining, High Accuracy, 6 (1) 32-39 (1975), (2)

55-67 (1978).

Yasunaga, N., et al., Mechanochemical polishing of single crystals with soft powders,

Proceedings 1st ICPE, 32-37 (1974).

333

Session III

Characterization of Machined Surfaces

Session Chairman

B. J. Hockey, National Bureau of Standards
:
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

SURFACE CHARACTERIZATION OF OPTICAL MATERIALS 1

H. E. Bennett

Michelson Laboratory, Physics Division

Naval Weapons Center, China Lake, CA 93555

A complete description of the surfaces of optical materials involves their topography, presence of particulates, crystallinity and

degree of disorder at the surface, impurities in the vicinity of the

surface, stoichiometry at the surface, and surface films. The most

direct method for observing small surface imperfections is by using

electron or optical microscopy. Often transmission electron microscopy using replication has too high a lateral resolution and too low

a depth resolution to investigate imperfections of greatest optical

interest. Surprisingly, optical interference microscopy has one to

two orders of magnitude higher depth resolution than electron microscopy, and techniques such as differential interference contrast,

sometimes called Nomarski microscopy, are extremely useful in evaluating surfaces qualitatively. For quantitative results, stylus techniques and interometry using fringes of equal chromatic order are very

useful and give results in good agreement with those obtained by scattered light measurements. Total integrated scattering provides a simple noncontact method for determining the rms height of surface irregularities and may be applied in various surface roughness determinnations. Evaluating the roughness of ball bearings is one example.

Surface impurities and stoichiometry at the surface also affect

the performance of optical materials. Impurity atoms in concentrations as low as 0.001 monolayer can be detected using Auger spectroscopy. There is a danger however, of disturbing the surface with the

high energy electrons required to produce the Auger electrons unless

very low electron currents are used. This problem is alleviated by

using electron spectroscopy for chemical analysis (ESCA) , which depends on photoemission induced by x-ray or vacuum ultraviolet irradiation rather than on electron bombardment. ESCA is also more sensitive to molecular forces than is Auger spectroscopy, so that compounds

as well as the atomic constituents can often be detected. Even more

sensitive is secondary ion mass spectroscopy. In many cases, surface

impurity levels in the 10" 6 range can be detected using this technique.

Surface films also affect the optical properties of materials, often

significantly. Their growth can be monitored using ellipsometry

,

which is insensitive to roughness introduced by the film and is sensitive to changes in average film thickness of a fraction of a monolayer. A new surface evaluation technique being developed is high

fluence diagnostics and multithreshold analysis. Most of the techniques discussed are applicable to nonoptical surfaces as well as optical surfaces, and should be used more widely in analyses of surfaces produced by ceramic machining.

Key Words: Auger spectroscopy; corrosion films; ESCA; impurities; microroughness ; optical microscopy; polishing defects;

scattered light; SIMS; stylus instruments.

x Work supported by the Naval Air Systems Command, Office of Naval Research,

and the Department of Energy.

337
1. Introduction

Optical surfaces when viewed from the vantage point of common machine

shop tolerances are virtually perfect. For example, the flatness of a good

A/4 optical surface is typically within 0.15 urn (6 yin) and with more care

0.015 ym surfaces can be achieved. The surface roughness of typical optical

surfaces is 0.0025 ym (0.1 yin) rms and values as low as 0.0005 ym have been , achieved. These values are orders of magnitude smaller than those usually

dealt with in machine shop practice, although with the advent of diamond- turn-j

ing procedures utilizing air bearings they can now be approached using precision machining techniques. Historically optical surfaces have been generated!;

by pitch polishing on machines which are not nearly as accurate as the finishe 1 component. The fact that optical surfaces can be produced at all is thus most

remarkable and a tribute to the ingenuity of the optician.

A complete description of the surface of an optical material involves its

|

topography, crystallinity and degree of disorder at the surface, impurities ini the vicinity of the surface, stoichiometry at the surface, and surface films.

Topography of optical surfaces may be divided into long-range deviations from i the desired surface, which is referred to as optical figure and is often deter

mined interferometrically , and shorter range imperfections. These short-range

imperfections may again be divided into (1) macroscopic imperfections such as scratches, cracks, voids, or bumps and (2) micro-irregularities only a few ten

of angstroms in height which cover all well -polished optical surfaces. Optical figure is important since it determines the resolving power of an optical

train and also the intensity of light which can be focused at the focal spot

of an optical system [l] 2. The short-range surface imperfections, together

with surface particulates, determine the fraction of the incident light which

is scattered or, in some cases, absorbed by the surface [2]. They may also

important in nonoptical applications. For example, macroirregularities such

as cracks can have a significant effect on the strength of brittle materials.

Surface irregularities can also affect the wear properties of surfaces in

mechanical contact such as bearings and ways.

In this paper we will first discuss techniques for measuring the surface

topography of optical surfaces and implications of these techniques for non

optical applications. We will then discuss the detection of surface impurities, stoichiometry, and lattice disorder induced by the surface finishing

process. Finally, we will discuss the presence of thin films at the surface

and their detection.

2. Surface topography

The topography of an optical surface may be divided into two categories:

short-range and long-range. The long-range deviations from an ideal surface,

i.e., surface figure, are usually determined either interferometrically , as illustrated in figure 1(a), or by measuring diffraction effects. They can be

established with uncertainties as low as A/80 to A/100 in the best cases

(A is the visible wavelength used in the tests, usually 0.6 ym or 24 yin).

Figure 1(a) shows an interference measurement of a good diamond-turned optical

surface. Shorter range imperfections in the surface are less commonly

measured quantitatively. For optical surfaces used in the visible and ultraviolet regions of the spectrum microirregularities usually contribute most to

scattered light, whereas macroscopic irregularities and particulates are the

determining topographic features causing scattering in the infrared. All can

be readily detected using optical or electron microscopy. Figure 1(b) shows

the copper surface whose optical figure was shown in figure 1(a) as viewed

using a differential interference contrast (DIC) , i.e., Nomarski, microscope.

The grooves in this machined surface are only a few hundredths of a micro2 F igures in brackets indicate the literature references at the end of this

paper

.

338
FIGURE (1/20th WAVE) EFFECTIVE ROUGHNESS 21.0 X RMS

(a)

Figure 1. Zygo (Fizeau) interferogram

and Normarski micrograph of a diamond-turned copper sample having an

excellent figure and low scattering

(from ref . 2)

100pm

meter in height and would be invisible in an ordinary microscope. The fact

that they can be readily observed using an interference microscope causes a curious problem. Such grooves do not appear on conventionally polished optical surfaces and might be cause to reject diamond- turned surfaces for optical

applications. A closer investigation, however, shows that these grooves have

such a wide spacing that light scattered by them is very close to the specular

direction, as illustrated in figure 2, and thus does not affect the optical

performance of the surface significantly. At the other extreme, irregularities separated by less than a few tenths of a micrometer (i.e., 20 yin or

less) which are regularly spaced will not contribute to scattered light either,

since the scattering angle, as seen in figure 2, would be more than 90°. The

light averages over such closely spaced irregularities and the sputtered silver

film whose micrography is shown on the left in figure 3 will scatter no larger

a fraction of the reflected light than the uncoated quartz surface seen on

the right. Similarly, the scratch is too narrow to affect optical performance significantly, although it is easily seen using the electron microscope.

Electron microscopy, although a very powerful tool in many areas of science,

often has too high a lateral resolution and inadequate height sensitivity

Figure 2. Nomogram giving separation

of surface features (spatial wavelengths) that produce scattering at

a particular angle and wavelength.

The dashed lines are an illustrative

example (from ref. 3).

10 r 1—i i l l l ill

SCATTERING ANGLE, 0

"Mami 1 1

NORMAL INCIDENCE ILLUMINATION

d - X/«in0

J I I I I I III

SEPARATION, d (urn)

to be of great use in optical surface analysis. A comparison of the lateral

resolution and height sensitivity of stereo electron microscopy and that

obtainable by stylus and interferometric techniques [3] is given in table 1.

Scanning electron microscopy has much lower sensitivity than conventional

electron microscopy, which has poorer sensitivity to heights than interferometric or stylus techniques. In most investigations of optical surfaces

these latter techniques are thus found to be the most useful. Descriptions

of interferometric techniques for surface evaluation using fringes of equal

chromatic order (FECO) are found in the literature [5,6]. Stylus techniques

339
SPUTTERED SILVER ON FUSED QUARTZ BARE FUSED QUARTZ

AV. HEIGHT 90130A

1

AV. WIDTH 475+84 X

AV. HEIGHT 44H6A

1

AV. WIDTH 4S8+160A'

Figure 3. Transmission electron micrographs

of surface replicas used in stereo electron

microscopy [4]. (a) Sputtered silver film on

fused quartz, and (b) bare fused quartz.

Average heights and widths of surface features are given.

Table 1. Comparison of statistical methods (from ref. 3).

Instrument Height

sensitivity

Lateral

resolution

Maximum

length

FECO interferometer

Talystep profilometer

Stereo microscopy

~3 A rms

~1 A rms

±20 A

~2 . 0 ym

~0 . 1 ym

~5 A

1 mm

2 mm

~1 ym

are equally powerful but are often abused. In order to prevent inelastic

deformation of the surface by the stylus, extremely light loadings are

required [7]. If loadings in the range of a milligram or less are used, it

is possible to study the surfaces of materials even as soft as KC1 without

introducing scratches [8]

.

Both microirregularities and more localized nonuniformities such as

scratches or cracks can be studied interferometrically or by stylus techniques. Figure 4 shows cracks observed in forged KC1 windows when grain

boundaries were not removed prior to final forging, as reported by Anderson

and Bennett [9] . Such imperfections introduce non-Gaussian height distribution statistics for optical surfaces. In most cases, however, the primary sources of roughness are the microirregularities, and the height distribution function is Gaussian to an excellent approximation [6] . In such cases

it is possible to unambiguously specify the microtopography of an optical

surface by giving its rms roughness. The range of rms roughnesses found at

Michelson Laboratory over the years for various optically polished materials

[3] is given in figure 5.

A complete description of an optically polished surface contains both

its height distribution function and its autocovariance function [2]. The

fraction of the incident light which is scattered into all angles by an optical surface depends almost entirely on its rms roughness, whereas the angle

into which it is scattered depends on the autucovariance function [7] . A

340
2 534 #537

25 jjffl

' "

50 /itn ~~l

54.7 A RMS

50 "m

83.7 X RMS

6° -8° SLOPES OF CRACKS

GRAIN BOUNDARIES REMOVED

BEFORE FINAL FORGING

100 X

50 tim

3.S A RMS

QUARTZ DIE

Figure 4. Surface finish on areas of forged KC1

windows with and without grain boundary removal

before final forging. The bottom trace is the

surface finish of the die used to forge the two

windows (from ref. 9).

FIIKFIl QUARTZ

SILICON CARRIflF*

5 *^*K3I^k^i^^HB 30

7 MM"« 'LIMITED NUMBER OF SAMPLES

11 HHHB^ — 339^ COPPER (DIAMOND TURNED)

Fl ECTROI FSS NICKEI

TITANIUM u w^b^mh^^^b 39

mPPFR

T7M

MOI Y8DENUM

ALUMINUM

BERYLLIUM COPPER

INVAR 21 HBBSBBKSSHi^i^i^i^B 70

MONEL* 21 g^^^^^^^^^^^ 72 ^

POTASSIUM CHLORIDE

STAINLESS STEEL*

A. 110 —

1 1 1 1 1 1 1 1 1 1

RMS ROUGHNESS (i)

Figure 5. Average rms roughnesses and range of

values measured for optical surfaces, as determined from total integrated scattering, interferometric, and stylus measurements (from ref . 3) .

simple noncontact method for evaluating surface roughness is thus to measure

the total amount of light scattered from it. Such a measurement is often

called total integrated scatter CTIS) measurement. Angular scattering

measurements, which make it possible to determine the bidirectional reflectance distribution function CBRDF) of the optical surface, are also useful

for detecting particulates or other contaminants on an optical surface [10]

If the height distribution function of the surface irregularities is

Gaussian and the rms roughness <5 is much smaller than the wavelength x, the

relation between the TIS and 5 is [2]

TIS m C4ir6/X) 2. (1)

341
For typical optical surfaces this relationship holds accurately in the visible and ultraviolet region of the spectrum and begins to deviate only in the infra- red [2], as illustrated in figure 6. The deviation is thought to occur because scattering from the microirregularities becomes so small in the infrared that particulates and macroscopic irregularities, whose scattering properties are independent of wavelength, begin to dominate. If the surface microirregular- ities are somewhat larger, eq. (1) holds in the infrared also, as has been demonstrated for ground surfaces [11].

i—i—i i i

io-«

10-6

AV-MIN SCATTERING

(103 particte/mm2

)

THEORY FOR 29.5 m -x

RMS ROUGHNESS \

9.5 A \

0.4 J I I I I L_ 1.0 2.0 4.0 WAVELENGTH Ipml

J I ' I I 10.0

Figure 6. Scattering from an aluminized polished

dense flint glass mirror. The diagonal line

gives the contribution predicted for microirregularity scattering (eq. (1)) by a 29.5 A rms

roughness surface. Circles indicate the minimum scattering observed, and the bars and squares

the difference between the average and minimum

scattering levels observed at several points on

the surface (from ref. 2).

Since for surfaces with Gaussian height distribution functions there is a unique relation between the rms roughness of the surface and the amount of

light scattered from it, light scattering provides a very convenient tool for

evaluating surface roughness quantitatively. Figure 7 shows an instrument for

making scattering measurments in either reflection or transmission [12] ; the

optical diagram is given in figure 8. Scattering plots in reflection and

transmission made with this instrument in the visible and in the infrared are

shown in figure 9 [12]. If a single reflecting surface is measured, such

data can be used to determine the surface microroughness

.

If average microroughness is the primary object of the scattering measurement, a much simpler arrangement than that shown in figure 8 can be used and

the light itself can do the averaging. A possible schematic diagram for such

an instrument [14] is shown in figure 10. The light beam from a point source

or a coherent source such as a laser is expanded to cover a significant area

of the surface to be tested. Light which is specularly reflected falls on a detector; the variation in signal is a measure of the scattered light and

hence of the surface roughness. To improve sensitivity, if required, both

the specular signal and the scattered light could be recorded. Such a system

could be used to evaluate cylindrical or spherical surfaces such as roller

bearings or ball bearings as well as flat surfaces, and should be particularly

useful as an evaluation technique for production samples. For example, one

may visualize a setup for checking the surface roughness of ball bearings in

which balls roll down a track and come to rest with the lower part of the ball

projecting through a hole. The light reflected from this projecting part of

the ball to the detector causes a signal which either indicates accept or

342
Figure 7. Photograph of the Optical Evaluation Facility

used for making total integrated scattering measurements in either reflection or transmission.

CO2

9.5-10.6/1 m

I HF/PF | -jr

2.5-3.0, 3.5-4.0/1 m

1.15, 3.39/zm

M8'

M10 0.351-0.364

0.4880, 0.5145/im

Kr

^D3

M2 N

* /13Hz \ ^ Ml

Ch F, 8p r- SF

M11 0.3507 0.5208 0.7525

0.3564 0.5309 0.7931

0.4762 0.5682 0.7993/im

0.4825 0.6471

0.4880(Ar) 0.6764

Figure 8. Optical diagram of the Optical Evaluation Facility showing

lasers and available laser lines, mirrors M, lens L, spatial filter

SF, attenuation filters F, chopper Ch, diaphragm d, Coblentz spheres

C, pyroelectric detectors D, and samples S (from ref. 12).

343
CYLINDRICAL MIRROR

(OR SPHERICAL MIRROR)

TEST OBJECT CYLINDRICAL ROD (OR SPHERICAL BALL) CYLINDRICAL LENS

(OR SPHERICAL LENS) FLAT TEST OR REFERENCE

SURFACE

L,

CENTERING

MICROSCOPE

REMOVABLE

PLANE MIRROR

/

THERMISTOR BOLOMETER

FIELD STOP BEAM

SPLITTER

CaF, LENS

POINT SOURCE

INTERFERENCE

FILTERS FOR 2 COLOR SYSTEM

Figure 9. Three-dimensional

plots of the scattering in

reflection and transmission

from a forged KC1 window in

the visible and infrared

(from ref . 13)

.

Figure 10. Schematic diagram of an instrument to measure roughness in the 0.4 to

to yin range (from ref. 14).

reject. An air jet cou

The ball is then ejecte

is repeated. The whole

tive. For small roughn

larger roughnesses (^0.

the roughness would not

optics effects would oc

does hold in the infrar

surfaces and that the s solving what is otherwi

Id rotate the ball if complete coverage was desired,

d and a new ball drops into the hole where the process

operation is simple, rapid, accurate, and nondestrucesses C^O-Ol v> m ) visible light could be used. For

1 jam) infrared light would be necessary since otherwise

be small compared to the wavelength and geometrical

cur. It has been shown experimentally that eq. (1)

ed as well as in the visible for appropriately rough

cattered light technique provides a simple method for

se a difficult measurement problem [14]

.

3. Impurities and surface defects

Other important parameters to characterize on optical surfaces are the

presence of impurities and surface stoichiometry . Various techniques using

electron spectroscopy are most useful in this regard. For example, figure 11

records a scan made of Auger electrons from a copper surface which was

polished using AI2O3 powder 115] . The powder has been imbedded in the copper

surface; its presence is indicated by the feature at 1400 eV. No feature is

seen near 1600 eV, showing that contamination from compounds containing silicon is absent. The sensitivity of Auger spectroscopy varies depending on

the sample and experimental conditions, but it is approximately 0.001 monolayer and thus gives us a very sensitive technique for detecting surface

In most cases, it is insensitive to chemical bonding and

rimarily for identifying atomic constituents. Also, the excieam can easily damage dielectric samples through electron

that care must be taken to maintain low irradiance levels.

contaminants

.

hence is used

ting electron

bombardment so

An alternate technique of comparable sensitivity is electron spectroscopy

for chemical analysis (ESCA) . The surface to be tested is irradiated either

with x-rays or with radiation in the vacuum ultraviolet, and the photoemitted

electrons are monitored. Since there is no intense electron beam irradiating

the surface, the probability of surface damage is minimized by using ESCA

instead of Auger spectroscopy. Also, chemical bonding information may sometimes be deduced from ESCA traces. Figure 12 shows an ESCA scan of an evaporated silicon oxide film made using x-ray irradiation. Peaks from Auger

344
AUGER ELECTRON ENERGY (eV) 600 800 1000 1200 1400

I I 1 1

Cu (Al2 03 POLISH) —mM/^

1 1 1 1

1

1

r T\l

Cu

"""^aT" s'i

1 1 1 1 1400 1600 AUGER ELECTRON ENERGY <«V) Figure 11. Auger electron energy Scan showing

the presence of A1 2 0 3 polishing compound on

the surface of polished copper (from ref. 15).

ENERGY (eV)

Figure 12. ESCA scan of a Si0 2 film. The energy

region around the Si 2 p peak is expanded in

figure 13.

electrons are seen as well as photoemitted electrons [16] . Of particular

interest is the Si 2 p peak, which is shown on an expanded scale in figure 13.

The particular sample investigated here, which was produced by reactively

sputtering silicon in an oxygen atmosphere, yields a trace which is identical

to the trace produced by a fused quartz sample. It is thus stoichiometric,

but this sample has a peak which is intermediate between the Si0 2 peak and

that produced by pure silicon and is thus not stoichiometric. This example

illustrates the advantage ESCA has in making it possible to investigate not

only the presence of contaminants but also to some extent the chemical composition and stoichiometry of the surface.

o Both ESCA and Auger spectroscopy probe as little a 2-20 A below the surface, since they depend on the escape depth of the emitted electrons [16].

They can be used to profile the chemical composition of surfaces to much

greater depths than this, however, by sputter-etching down through the surface

and monitoring one or more emission peaks as a function of depth. This technique makes it possible, for example, to investigate the composition of sur345
EVAPORATED Si + 0,

Figure 13. ESCA scan of the Si 2 peak on an expanded

scale. The upper traces show the difference in

energy between silicon in the form of Si0 2 and as free silicon, while the lower traces are for two

silicon oxide films prepared under different conditions . face films, evaluate multilayer coatings, and determine the depth of disturbed surface layers in some cases.

A drawback of both ESCA and Auger spectroscopy is their reduced sensitivity to impurities as compared with bulk analytical techniques. No bulk

technique can sample only a few angstroms below the surface, but typical sensitivities of bulk techniques are often in parts per million, as compared to

parts per thousand for ESCA and Auger spectroscopy. A solution to this problem is secondary ion mass spectroscopy (SIMS) . The surface is bombarded with

energetic ions, usually of argon or other rare gases, and the ejected secondary ions are mass analyzed. Both positively and negatively charged ions may

be measured with sensitivities approaching those of bulk techniques, i.e.,

ppm. SIMS is thus a very powerful technique, and although at present it is

difficult to extract quantitative data from SIMS measurements, it makes a very useful companion technique to ESCA or Auger spectroscopy. A SIMS trace

of a very pure NaF coating [16] is shown in figure 14.

Ellipsometry furnishes another tool for surface analysis. It can be

used either in a vacuum chamber or in air, and is particularly useful for

studying the growth of thin films or the presence of a polish layer on prepared surfaces. Figure 15 shows the growth of a silver sulfide tarnish layer

on a silver film prepared in our laboratory and exposed to air [17]. Note

that the film growth is only a few angstroms in average thickness per day,

but differences in that growth rate at different times in the day are easily

detected ellipsometrically . The arrows indicate mornings. The rate of film

growth was significantly larger during the day, when the laboratory was open

and being used, than at night. When the film was enclosed in an atmosphere

of dry nitrogen, growth virtually stopped. Since ellipsometry depends on

phase measurements rather than intensity measurements, it is relatively insensitive to surface roughness, scattered light, or reduction in reflectivity,

and hence can be used to monitor film growth which significantly reduces the

optical quality of the surface. Silver sulfide is a good example of such a film.

There are many other surface evaluation techniques which could be mentioned. One promising new technique now under active development in our

346
OH_

CH

_

X10

\

X1

i u/

C2 H2

C|—

llLilliilLjl

XI 00

( till inn 1 11 mini iimi iiiiiiiiiiiiiiniiiiiiiiiiiiiiiiiiiii

Na /C2 H

2 ci-

\ 1 i X10

NaF-SIMS

_i 1 1 i_

X10

H+

Na

Na,

UK.

Na2 F

+ 5LJlJ!£L

xio

Si +

_l L.

NaF

+SIMS

X100

20 40 . 60 MASS (a.m.u.l 80 100

Figure 14. SIMS scan of a very pure NaF film

showing the presence of impurities in the

parts per million range.

40 1

1 1 1 1 1 1

1 r

TIME OF FILM GROWTH (DAYS)

Figure 15. Growth of a silver sulfide tarnish

layer on a silver film in room air as measured ellipsometrically . The vertical arrows

indicate mornings. After the dashed line on

day 8 the sample was enclosed in a nitrogen

atmosphere (from ref. 17).

laboratory by Dr. J. 0. Porteus and colleagues is high fluence diagnostics

[18,19] . The sample is mounted on the ultrahigh vacuum chamber shown in

figure 16, which contains various evaluation instruments including an ion col

lector, visual microscope, scanning electron imager, and Auger analyzer. The

sample is irradiated with a high intensity pulsed laser beam, and the thresholds for various processes such as ion emission, plasma formation, change in

work function, melting, intergranular slip, pitting, and so on are determined

Significant differences between these various thresholds are found depending

on the surface treatment, purity, and grain structure of the samples.

347
Figure 16. Ultrahigh vacuum laser damage test chamber

containing instruments to evaluate effects occurring

on the sample during laser irradiation.

4. Conclusion

A complete description of a surface must include its surface topography,

structural imperfections, and surface impurities, including the presence of

surface films. Although optical surfaces are much more nearly perfect than

most surfaces in common experience, they still have the same types of impertions as other surfaces. Hence, techniques used to explore surfaces used in

optical applications may be useful in characterizing other surfaces as well.

A prime example as far as surface topography goes is a bearing surface, which

for ceramic bearings of materials such as silicon nitride must be finished to

nearly optical tolerances. By using techniques such as analysis of scattered

light, FECO interferometry , stylus^measurements , and DIC microscopy, surface

irregularities as small as 5 x 10 ym in height have been detected on optical

surfaces. With some modifications ground surfaces can also be evaluated using

some of these techniques. Surface impurities can be detected in the range of

parts per thousand in thicknesses of only a few monolayers using ESCA and Auger spectroscopy, and in parts per million using SIMS. The variation in impurity concentration with depth can be determined by sputter-profiling. Nonstoichiometry can also be detected, and by using ellipsometry the growth of

thin surface films in air or vacuum can be monitored with a sensitivity of a fraction of an angstrom in average film thickness. New evaluation techniques

such as high fluence diagnostics are also under development. These techniques

are very powerful and should be more widely used in investigations of ceramic

surfaces

.

348
References

[1] Bennett, H. E., "Thermal distortion thresholds for optical trains handling high pulse powers", in Laser Induced Damage in Optical Materials : 1976, A. J. Glass and A. H. Guenther , eds . , National Bureau of Standards,

Washington, DC (NBS spec. Publ. 462), pp. 11-24; Bennett, H. E. , "Factors

limiting the energy transmitted by an optical train", in Laser Damage

of Optical Materials , A. J. Glass and A. H. Guenther, eds~. (Academic

Press, New York , 1979) , in press.

[2] Bennett, H. E., "Scattering characteristics of optical materials", Opt.

Engr. 17, 480 (1978).

[3] Bennett, J. M. , and Elson, J. M. , "Surface statistics of selected optical materials", in Laser Induced Damage in Optical Materials : 1977 , A.

J. Glass and A. H. Guenther , eds . , National Bureau of Standards,

Washington, DC (NBS Spec. Publ. 509), pp. 142-155.

[4] Butler, D. W. , "A stereo electron microscope technique for microtopographic measurements", Micron 4_, 410 (1973) .

[5] Tolanski, S., Multiple-Beam Interferometry of Surfaces and Films (Clarendon Press, Oxford, England, 1948).

[6] Bennett, J. M. , "Measurement of the rms roughness, autocovariance function and other statistical properties of optical surfaces using a FECO

scanning interferometer", Appl. Opt. 1_5, 2705 (1976).

[7] Elson, J. M. , and Bennett, J. M. , "Relation between the angular dependence of scattering and the statistical properties of optical surfaces",

J. Opt. Soc. Am., in press.

[8] Bennett, J. M. , "Surface evaluation techniques for optical components",

in Laser Damage of Optical Materials , A. J. Glass and A. H. Guenther,

eds"! (Academic Press, New York, 1979) , in press.

[9] Anderson, R. H. , and Bennett, J. M. , "Optical properties of KC1 forged

between optically polished dies", in Laser Induced Damage in Optical

Materials : 1978 , A. J. Glass and A. H~ Guenther , eds . , National Bureau

of Standards, Washington, DC, in press.

[10] Williams, V. L., and Lockie, R. T., "Optical contamination assessment

by BRDF measurement", Opt. Engr. 1_8 (1979), in press.

[11] Bennett, H. E., "Specular reflectance of aluminized ground glass and the

height distribution of surface irregularities", J. Opt. Soc. Am. 53, 1389

(1963).

[12] Bennett, H. E., and Stanford, J. L., "Structure-related optical characteristics of thin metallic films in the visble and ultraviolet", J. Res.

Natl. Bur. Stands. 80A, 643 0-976)

[13] Anderson, R. H. , and Bennett, J. M. , "Properties of KC1 optical elements

forged between optically polished dies", in Proceedings of the High

Power Laser Optical Components and Component Materials Meeting , j" S.

Harris and C. L, Strecker, eds." [Defense Advanced Research Projects

Agency, Arlington, Va., December 1977), pp. 177-186.

[14] Bennett, H. E., "Measurement of smooth surface finishes", Industrial

Quality Control 20 [8], 18 CFebruary 1964).

349
[15] Bennett, H. E . , and Archibald, P. C. , "Optical requirements for laser

mirrors", in Laser Induced Damage in Optical Materials : 1974 , A. J.

Glass and A. H. Guenther, eds ., National Bureau of Standards , Washington.

DC, (NBS Spec. Publ. 414), pp. 23-29.

[16] Green, A., Naval Weapons Center, private communication, 1978.

[17] Bennett, H. E., Peck, R. L., Burge, D. K. , and Bennett, J. M. , "Formation

and growth of tarnish on evaporated silver films". J. Appl. Phys . 40.

3351 (1969). —

[18] Porteus, J. 0., Jernigan, J. L . , and Faith, W. N., "Multithreshold measurement and analysis of pulsed laser damage on optical surfaces", in

Laser Induced Damage in Optical Materials : 1977 , A. J. Glass and A. H.

Guenther, eds., National Bureau of Standards, Washington, DC (NBS Spec.

Publ. 509), pp. 507-515.

[19] Porteus, J. 0., Decker, D. L., Jernigan, J. L. , Faith, W. N. , and Bass,

M. , "Evaluation of metal mirrors for high power applications by multithreshold damage analysis", IEEE J. Quantum Electron. QE-14 , 776 (1978).

Discussion

VORBURGER

The optical technique for roughness measurement that you discussed is useful for highly

polished surfaces only. Are there any possibilities for an optical technique that will be

useful for measuring roughnesses on the order of a micrometer?

BENNETT

To measure roughness of the order of a micrometer rms in height one should go to wavelengths of 30-40 ym, i.e., the far infrared. In most applications of which I am aware,

however, the roughness values to be measured are much smaller than a micrometer. For

example, Dr. Sibley, in his talk on ceramic bearings, concluded that the maximum roughness

they could tolerate was about 0.01 ym AA and some balls had roughness as low as 0.004 ym.

For evaluating such roughnesses easily and accurately, light scattering is almost ideal and

visible wavelengths can be used. For roughnesses of 0.1 or 0.2 ym, one should use infrared

wavelengths of 5-10 ym. This is the roughness range which has been of most interest in the

past and it is too large to evaluate accurately by light scattering techniques using visible

wavelengths. Measurement of light scattering using infrared wavelengths for roughnesses of ^ 0.1 ym or visible wavelengths for roughness ^ 0.01 ym provides us with a very convenient,

rapid, noncontact method of determining surface roughness which is particularly attractive

for production applications.

350
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C. , 1979).

THE NATURE OF STRENGTH CONTROLLING

MACHINING FLAWS IN CERAMICS

R. W. Rice and J. J. Mecholsky, Jr.

Naval Research Laboratory

Washington, D. C. 20375

The nature of machining flaws introduced in a variety of

glass, single crystal, and especially polycrystal bodies are reported, based on fractographic determination of such flaws as

the source of mechanical failure. Particular attention is given

to grinding where the anisotropy of strength due to the direction of grinding relative to the tensile axis is shown to be due

primarily to a dual population of flaws of differing shapes. One

set of flaws form ~ perpendicular and another ~ parallel to the

grinding direction. The latter flaws are typically substantially

more elongated and often larger than the former and thus give

lower strengths for stressing perpendicular to the direction

of grinding. Sawing, sanding, and polishing are also shown to

result in similar dual flaw populations. The character of flaws

are examined as a function of the type of material and machining, as well as specimen size and shape. The latter suggests a

limited strength-size effect due to machining flaws. However,

the key findings are that machining falws do not vary greatly

with typical variations of machining or material parameters, e.g.

composition, grain size, or hardness.

Key Words: Brittle failure; flaws, flaw populations; fracture

origins; f ractography ; grinding; machining, polishing; sanding;

strength-size effects.

1. Introduction

Machining of most glasses, single crystals, and dense polycrystalline

specimens is typically assumed to introduce flaws which control their mechanical strength. However, in discussing the evidence for such flaw controlled failure versus dislocation controlled failure, Rice Ll] , at the

previous Machining Conference, pointed out that there had been little or no

direct verification of flaw failure. This paper summarizes and significantly

extends recent work on the nature of flaws introduced in a variety of ceramic materials, as a function of various machining and specimen parameters [2-

9]. The nature of the two flaw populations that result in the common anisotropy of strength associated with different grinding directions are treated

first. Then, more limited studies of other modes of machining are compared

with these grinding studies, followed by some discussion of flaw variations

and distributions.

Present mailing address: Sandia Laboratories, Albuquerque, New Mexico 87185.

Figures in brackets indicate the literature references at the end of this

paper.

351
2. Experimental Procedure

The grinding observations summarized here cover several years involving

several machine operators as well as a variety of materials. Grinding was

done predominantly, but not exclusively, with 320 grit diamond wheels, most

often either with a vitreous bonded wheel (~ 10 cm dia. operating at ~

7900 rpm) or a metal bonded wheel (~ 20 cm dia. operating at ~ 1725 rpm) . A

range of feed rates, commonly of the order of 1 cm/sec, and depths of cut,

commonly ~ 0.05 mm, per pass and at least two different sizes and styles of

grinding machines have been used [3,10,11], In addition, some observations

are from specimens or components of different materials, sizes, and configurations, ground and tested by at least three other organizations. Thus,

the results should have broad applicability.

Diamond sawing was with typical diamond saws, e.g. as described elsewhere ClO] with cutting rates of the order of 2 cm^/min for the MgF2 of this

study. Hand sanding was done with 320 or 600 grit SiC paper under dry conditions with finger pressure only. Diamond polishing was typically done to ~ 1 jtim grit size on a maple lap with kerosene as a fluid. Some specimens or

components were polished and tested in other laboratories.

Mechanical testing was typically by 3 point flexure at ~ 22°C on bars ~ 1.3 x 2.6 mm in cross section on a span of ~ 1.3 cm with a head travel of ~ 1.3 mm/min, but other size and test results are included. Fracture origins

were determined on tested samples by optical and scanning electron microscopy

using the techniques described by Rice [12],

3. Strength Anisotropy and Flaw Shape as

a Function of Grinding Direction

At the previous ceramic machining conference, Rice reported studies

showing that an anisotropy of strength resulted in a variety of glasses,

single crystal, and polycrystalline flexure bars as a result of stressing

relative to the grinding or sanding direction and the bar axis [11]. Characteristically, tensile stresses parallel with the grinding direction, i.e.

parallel to the bar axis, gave the highest strength, while tensile stresses

perpendicular to the grinding direction gave a lower strength, often by up

to 50%. Kirchner also independently reported similar anisotropy of strength

relative to the grinding direction of round rods of fine grain A12

C>3 [13].

Subsequently, other investigators have reported similar anisotropies of

strengths especially in hot pressed Si3N4 [14] which was also included in

Rice's original study.

Rice originally proposed that the cause of this anisotropy of strength

was stress concentrations associated with resultant grinding striations or

grooves that could be viewed as two facing surface steps, each having a

stress concentration associated with it of ~ 1 + 0.7 (h/r) where h is the

height of the steps and r is the radius of curvature at the base of the

step. Thus, for stressing parallel with the grinding striations there would

be no stress concentration while for stressing perpendicular to the direction of grinding the stress being normal to the grinding striations would

provide the maximum stress concentrations at the surface steps or grooves.

Several subsequent investigators have also generally associated the strength

anisotropy with these grinding striations. However, there are two factors

which show that these grinding striations are typically not a major factor

in the anisotropy of strength associated with stressing relative to the

grinding direction.

The first factor showing that the grinding grooves thermselves are

probably not the predominant factor in the strength anisotropy as a function

of the stress direction relative to the grinding direction is a closer eval352
uation of the nature of the grooves themselves and their possible stress

concentration. If a grit particle gouged out a simple groove defined by the

average upper limit of its protrusion from the grinding wheel, h ~ r, it

would give an ~ 70% increase in the component of stress normal to the groove.

Thus, if h were of the order of the flaw sizes in ceramics, the ~ 70%

stress concentration would suggest the possibility that the grooves themselves are the source of failure. However, examination of more severe grinding grooves which are typically associated with the flaws causing failure

shows that h/r < 1, e.g. \, which is also indicated in the single particle

machining experiments of Gielisse et al [15]. This reduces the stress concentration to ~ 30% or less. While local irregularities in the grooves

would locally increase the stress concentrations, these are not likely to

extend over a sufficient range to change the situation. Further, the grooves

often do no appear to be as deep as typical flaws. Another aspect of the

grooves themselves that indicates they are not the predominant cause of

strength anisotropy is the fact that they are often not continuous [ll],

and failure is observed to sometimes occur where grooves have not formed

or have been removed by spalling. Note also that correlations between grit

particle size and flaw sizes [16-19] do not have to arise from the groove

being the flaws. It is quite likely that the different sized grit particles

act as different size indentors with larger grit sizes introducing cracks

that extend deeper below the surface. Thus, cracks extending in from the

surface due to the grooving by the particle may well correlate with particle

size. However, because of the increased distance of the crack periphery

from the surface steps, the effect of the step stress concentration on the

crack would be greatly reduced.

The second, and major factor, showing that the grinding striations

themselves are typically not a major factor in the anisotropy of strength as

a result of stressing relative to the grinding direction is direct fractographic identification of the nature of flaws from which mechanical failure

initiates in specimens. The net result of a substantial amount of study is

that there are basically two different sets of flaws extending into the body

from the surface generated by grinding (figure 1) . Sample fractographs illustrating the difference in shape between the two sets of flaws are shown

for fine grain (figure2), large grain (figure 3), and single crystal bodies

(figure 4). Other examples will be shown and discussed later. Additional

examples of the two sets of flaws can be seen in references 3-9 and 20-24.

These studies show that one set of flaws form essentially parallel with,

and approximately at, the bottom of the grinding groove. The other set of

flaws typically form ~ perpendicular to the grinding grooves, apparently due

to stick-slip or other dynamic variations in local tensile stresses generated by the machining action ~ parallel with the directions of machining, i.e.

of grinding particle motion. Typically, both sets of flaws extend to similar depths. However, the flaws forming along the grinding grooves are characteristically more severe than those forming perpendicular to the grinding

grooves consistent with the observed strength anisotropy. The greater severity of the flaws along the grinding grooves is generally due

more to their greater length rather than their depth (figures 1-4).

Analysis of extensive observations of the anisotropy in flaw shape between the two sets of flaws on strength is given in table 1. Here the observed anisotropy of strength is compared with the strength anisotropy predicted from the Griffith equation using the observed flaw size and shape

parameters. This clearly shows that the anisotropy in flaw shape is typically the overwhelming factor in the strength anisotropy.

In addition to the above more detailed, quantitative studies, a variety

of qualitative observations further extend the scope of the observation of

strength anisotropy being due to the differences in shape of the two flaw

populations associated with machining as indicated in table 2. Thus, for

example, in studies of hot pressed Si3N^, limited determinations of machining flaws at fracture origins of bars machined parallel with the tensile (bar)

353
TRANSVERSE

Figure 1. Schematic representation of flaws introduced in scratching a ceramic surface. While the main focus of the schematic is the effect of

an abrasive grit particle on the surface of the ceramic, it is also reasonably representative of what commonly happens in a scratch hardness

type test. One of the main differences between these two types of scratching operations is that in the abrasive machining operation extensive spalling often occurs along the path of the grit particle which is not shown

in order to keep the schematic simple. Such spalling may or may not occur

in the scratch hardness type test depending for example on the load and

material. This sketch schematically summarizes the basic findings of extensive observations on machining flaws which are also generally consistent

with observations made from scratch testing; namely that the motion of

particle or an indentor over the surface introduces two sets of flaws extending into the body from the surface. One set of flaws is essentially

parallel with the groove. Such flaws are typically reasonably elongated

either in a fairly continuous fashion or in a more irregular fashion due

to the partial or complete overlap of a number of shorter flaws. The

second set of flaws are typically approximately perpendicular to the motion of the particle or point and their periphery often approaches a semicircle. While some of the second type of flaws may be relatively planar

they often have some concavity towards the direction from which the particle or point approached the point where the flaw initiates. The insert

to the right schematically illustrates the approximate analogue for a static indentation. While for simplicity, a round particle has been shown

from which more than two sets of flaws might well form in a simple indentation test only two have been shown here since this would tend to be more

354
representative of the angular nature of particles similar to what one obtains with an angular, e.g. Vickers, indentor.

Figure 2. Example of machining flaws at fracture origins in a fine grain

ceramic machined parallel and perpendicular to the tensile axis of the

bar. A and B represent lower and higher magnification of the fracture

origin of a dense, fine grain mullite bar machined parallel with the

tensile axis; fracture stress ~ 325 MPa (47,000 psi) . Note the nearly

semicircular, i.e. half-penny shape of the flaw periphery, but that the

flaw has some curvature to it, e.g. its shape is similar to half a clam

shell. Figure C and D show fracture surfaces and the fracture origin of

a test bar of the same material ground perpendicular to the tensile

axis; fracture stress ~ 207 MPa (30,000 psi). Note some difference in

depth of the flaws in A-D, the more elongated character of that in C and

D, and the fact that the latter consists of at least two partially overlapping segments.

355
Figure 3. Examples of machining flaws at the fracture origins for parallel

and perpendicular machining of a large grain ceramic. A and B show

the fracture origin area and the specific failure causing flaw (arrows)

as a result of grinding a dense, sintered Y2O3 (Yttralox) bar parallel

with the tensile axis. Note the nearly semicircular nature of the final

flaw boundary and that the flaw is substantially less than the size of

the grain in which it formed. C and D show the fracture surface and

fracture origin (arrows) of a specimen of the same material ground perpendicular to the tensile axis. Note the more elongated character of

the flaw but that it is again substantially smaller than the size of

the grain in which it is located.

356
Figure 4. Examples of flaws at the fracture origins in parallel and perpendicular ground MgAl2C>4 single crystal. A and B show the fracture

surface and fracture origin (arrows) of a stoichiometric MKAI2O4 crystal

ground parallel with the <100> tensile axis and having a Ilio} tensile

surface, fracture stress ~ 145 MPa (21,000 psi). C and D show the

fracture surface and origin (arrows) of a sample of the same material

and orientation ground perpendicular to the tensile axis; fracture

stress ~ 117 MPa (17,000 psi). While the flaw shapes shown here are

somewhat more irregular than often observed, the flaw causing failure

from machining parallel with the tensile axis is definitely of similar

depth but less elongated than that causing failure from machining perpendicular to the tensile axis.

axis show the flaws to commonly approach ~ half penny flaws, while more extensive observations on specimens ground perpendicular to the tensile axis

both in the form of flat bars and round tensile specimens clearly show much

more elongated flaws [20,2lJ (see also fig. 18). Similarly, large round

tensile specimens of AI2O3 which were ground circumferentially [22], i.e.

perpendicular to the tensile axis characteristically show more elongated

flaws at fracture origins than observed in AI2O3 or other materials ground

parallel with their tensile axis. Also, limited studies of P-A1203 bars

ground perpendicular to the tensile (bar) axis show more elongated flaws at

fracture origins than the ~ penny shaped machining flaws at fracture origins

of bars ground parallel with the tensile axis [23 J. There are also preliminary indications of strength anisotropy in a cermet (table 2). Finally,

Richerson has shown evidence of grinding direction dependent strengths in

some denser reaction sintered Si^N4 [24] consistent with other observations

of machining flaws in such material [20,21].

The observations of this section on two sets of flaws from grinding,

one set forming generally parallel with and essentially ~ along the bottom

of the grooves generated by the individual grinding particles and the other

357
TABLE 1

EFFECT OF GRINDING ON FLAW GEOMETRY OF CERAMICS

Material

A. Glasses:

Silica

A luminosilicate

Borosilicate

Lead Silicate

Soda Lime

Glassy Carbon

B. Crystals:

Ti02

Ti02

Ti02

Ti02

Ti02

MgAl20

4

C. Polycrystals

Keatite Glass

^ . ,1 Orient

.

or kj • o . (urn) G.D. 2

No. 3

Fracture

(MPa)

4

( um) /i,4,5 a/b '

II 7 73±10 45±16 1.3±.4

X 4 62±8 75±37 1.0±.4

II11 12 117±10 39±9 1.6±.2

j. 5 80±7 71±10 0.5±.l

II 3 99±8 35±6 0.9±.3

x 4 92±9 39±9 1.1±.3

II 7 69±3 43±17 1.35±.6

X 7 59±17 64±38 1.26±.8

II 25 97±5 28±9 1.6±,2

X 16 68±8 67±5 0.5±.2

II 10 4u±y u . y± .

o

X 9 87±18 35±15 1.2±.6

<110> II 6 113±7 25 0.7

(110) X 6 120±4 30 0.7

<iTo> II 5 113±3 30 1.5

(001) X 5 77±3 70 1 .

5

<010> IIII 6 288±19

(100) X 5 239±12 25 1.3

<010> II 5 108±6 60 1.3

(001) X 5 103±6 70 1.3

<001> II 8 285±86 10 0.4

(100) X 7 183±36 35 0.18

<100> II 3 217±28 15±0 0.4±0.6

(110) X 3 165±28 14±7 0.4±0.6

Ceramic

MgF2

Mullite

B4C

B4C

CaF,

Yttralox

< 1

< 1

1-3

2-10

100-200

50-150

100-200

II 6 180±39 16±5 0.9±0.4

X 7 102±8 36±15 0.5±0.2

II 10 87±2 54±21 1.1±.07

X 10 53 ±2 89 ±20 0 . 5± . 08

II 7 319±35 41±23 1.0±0

X 6 259±54 24±5 0.4±0.3

II 5 374±69 19±3 1.0±0

X 9 154±24 26±10 0.2±0.07

II 4 282±228 19±9 0.7±0.5

X 6 250±55 23 ±3 0.3±0.2

II 3 50±13 23±7 0.7±0.2

X 3 40±5 33±7 0.6±0.6

II 6 99±6 36±13 0.8±0.3

X 7 77±8 44±20 0.7±0.3

ql/g||

0.84

0.68

0.9

0.86

0.65

0.93

1.1

0.7

0.83

1.0

0.6

0.8

0.6

0.6

0.8

0.6

0.9

0.8

0.8

B

0.86

0.69

0.9

0.82

0.65

1.03

0.9

0.7

0.9

0.5

0.8

0.5

0.6

1.0

0.8

0.7

0.8

0.9

"Orientation of single crystals < > gives tensile axis, and ( ) tensile sur358
face. G.S. = grain size.

'G.D. = grinding direction relative to tensile axis.

'no. = number of specimens.

^c = smallest flaw dimension (smallest of a or b below).

'b = flaw half length along surface; a = flaw depth.

'Ratio of fracture stress for perpendicular to parallel grinding.

B = $

*a El — 1~c~ wnere c

n and c x are tne smallest flaw dimension for paral11 x lei and perpendicular grinding respectively and and 3>± are respective elliptic integrals for the flaw shapes.

TABLE 2

OTHER MATERIALS IN WHICH GRINDING DIRECTION

EFFECT HAS BEEN INDICATED1 BY

Material Material

2 _£ a , b^

Sapphire X X

H.P. A1203 X X

S. A1203 X

H.P. Mullite X X

H.P. P-A120

3 X X

''"In studies by the authors

2

S = sintered, H.P. = hot pressed.

Zr02~fine grain

partially stabilized

H.P. SiC

H.P. Si3N4

75% TaC-25 Inconel5 Cermet

x

X

X

a ,b

x

x

x

cr = strength, i.e. strength anisotropy.

a,b = flaw depth and half length, i.e. flaw geometry.

*From selection of bars with ~ circular pattern milling or grinding marks ~ || to tensile axis (173 ± 33 ksi) and ~ x to tensile axis (140 ± 38 ksi)

for 3 spec. each.

set generally forming ~ perpendicular to those grooves, are consistent with

cracking patterns observed over a period of years by a variety of investigators for both static hardness impressions and especially dynamic, i.e.

scratch, hardness grooves [25-29] including earlier work, Bruche and

Poppa [25, 26]. Note also that the flaws ~ perpendicular to the grinding

direction have a curvature consistent with that expected for the motion of

the grinding particles. Thus, while there are variations in these two sets

of flaws, they are a basic aspect of grinding as well as other machining

processes as will be discussed later.

Before proceeding with a more detailed description of machining flaws

and variables effecting them, it is appropriate to briefly consider the detection of these flaws and their meaning. Typically, the flaws observed

are believed to be the original machining flaws for two reasons. The first

359
of these is based on the fact that cracks typically propagate normal to the

tensile stress causing their propagation., Thus, since the orientation of

the stress forming the crack during machining can almost never be duplicated

by the applied stress causing failure, the crack will be propagated on a

somewhat different surface than that on which it originally formed. This

change of surface results in a demarcation, which is generally detectable

unless failure is nearly or completely intergranular , in which case the demarcation is generally obscured by the grain to grain variation of the

crack topography. Since other possible changes in local stress around a

crack due to crack velocity effects are much more subtle, if they exist at

all, the demarcation between the original flaws and its first subsequent

motion for failure are clearly predominant and typically what is observed.

This is a major reason why the purely qualitative speculation of Doremus [30l

that similar flaws they observed at the fracture origins of machined glass

bars represent not the original flaws but some intermediate stage of failure

should be rejected.

The second reason that the demarcations seen are attributed to the original flaws rather than some stage of subsequent propagation is due to

changes in fracture mode that can occur during some stage of propagation.

Mecholsky et al have shown that environmentally induced slow crack growth

often proceeds intergranularly in polycrystals , e.g. MgFg [31]. Thus, since

both the original machining flaws and the subsequent final fast fracture are

mainly, or exclusively, transgranular in nature, the slow crack growth region is clearly defined showing the flaws such as those in this paper to

typically be the original flaw and not involving some subsequent subcritical

growth.

4. Other Aspects of Machining Induced Flaws

4.1 Microstructure , Composition and

Machining Flaws from Grinding

Extensive studies of interaction of machining flaws with porosity have

not been conducted. However, a number of observations indicate that isolated pores near the surface can be connected to the surface by machining

flaws, e.g. figure 5. See also figure 13C and D.

Substantially, more extensive observations have been made on the effect of grain size on machining flaws. The essential result is that over

the typical range of grain sizes encountered there is at best a limited

decrease in the depth of flaws introduced by machining with decreasing

grain size. Thus, for example, note the similarity in flaw sizes of specimens of comparable size for single crystals and polycrystals of MgAl204 of

different grain sizes (compare table 1 and figure 4 with figure 3 of

reference 22), different grain sizes of B4C (table 1) and single and polycrystalline flaws in AI2O3 [22]. A variety of other more limited, e.g.

non-quantitative observations, support this. For example, note the machining flaw at a possible, but uncertain, fracture origin in large grain CVD

in figure 6 has ~ the same depth as the typical flaws observed in fine grain

hot pressed Si3N4 (see references 20 and 21 as well as figure 18)

.

Further, note in general that there are limited differences in flaw

depths and no systemmatic pattern between materials of different compositions

and hence of differing properties such as hardness. Thus, note the typical

flaw sizes for the variety of materials of similar specimen size and a common machining direction, do not vary over a significant range (table 1).

This further reinforces the limited effect of grain size on the size of machining flaws that control strenght (e.g. compare figures 2-4). The ramifications of the observation that there is little or no dependence of the

size of machining flaws on grain size of the body being machined, are discussed in another paper [22].

360
PZT cr

f

~15 KSI

Figure 5. Example of pore-machining flaw interaction as a source of failure.

On the fracture surface of this commercial lead zirconate titanate sample,

three similar size machining flaws can be seen along the surface. Flaw 1

has connected with a pore, P, somewhat below the surface, to form the

primary source of failure. Flaws 2 and 3 indicate two other machining

flaws

Figure 6. Grinding flaw in large grained Si3N4. A and B show respectively

lower and higher magnification of a probable fracture origin and probable

flaw (arrows) causing failure in this large, columnar grain, CVD Si3N4

. Note also the two or three nearly concentric markings suggesting progressive stages of flaw development. Such markings with varying spacings

are observed fairly frequently with flaws forming approximately perpendicular to the direction of machining. Photo courtesy of Dr. Carl Cm. Wu.

361
4.2. Machining Flaws from Other Modes of Abrasive Machining

Three other major modes of machining, diamond sawing, hand sanding with

abrasive paper, and polishing (with diamond paste) have been studied to varying extent. Consider first diamond sawing. While extensive studies have not

been conducted, direct comparisons have been made using optical grade, hot

pressed, MgF2 as shown in table 3. The circular nature of the blade leaves

TABLE 3

* FLEXURAL STRENGTH OF AS-DIAMOND SAWN MgFg

Strength

Bars No. of Tests MPa ksi

0.36 x 0.65 cm

1.9 cm span

0.23 x 0.55 cm

1.27 cm spann

10 71 ± 21 10 ± 3

3 100 ±23 15 ± 3

Kodak IR window grade tested in 3 point flexure with a 1.27 mm/min head

travel rate.

an arc of "sawing" grooves with, in turn, presumably result in a directionality of the strength of diamond sawn specimens if they were tested as a

function of such orientation. This has not been explored, most specimens

had the sawing grooves at intermediate angles (e.g. ~ 45°) to the tensile

axis. The strengths of typical laboratory size bars as a result of such

diamond sawing are in the same range but somewhat less than those from diamond grinding. Fractographic studies in this material, which is excellent

for such studies, show definitive flaws at fracture origins from diamond

sawing. These flaws are very clearly associated with more extreme surface

grooves. Again both sets of flaws, i.e. those essentially ~ parallel and ~

perpendicular to the grooves are seen (figs. 7 and 8) with the former tending

to dominate failure with the grooves at ~ 45° to the tensile axis.* Both

the strengths (table 3) and actual flaw observations (figs. 7 and 8) show

similar size flaws from sawing as from the comparable direction of grinding.

Previous studies have shown that stressing relative to the direction of

hand sanding results in a similar anisotropy of strength relative to the

direction of abrasive particle motion Cll], Thus, hand sanding (dry) with

SiC abrasive paper clearly suggests that there are again two populations of

flaws, one population forming parallel and the other forming perpendicular

to the motion of the abrasive particles. The difference in the shapes

of the two sets of such sanding flaws is thus apparently very similar to

those of the two sets of flaws from grinding. Again limited, but definitive,

observation support this conclusion, e.g. note the character of flaws associated with hand sanding (fig. 9) of very fine grain, high strength SiC resulting from the CNTD+ process [32]. It should be noted that this similarity

suggests that the differences between grinding and sanding, e.g. orders of

magnitude difference in abrasive particle velocity have a limited effect

Such results indicate that much study of mixed mode failure can be done

directly with machining flaws by varying the angle of machining relative

to the tensile axis, rather than using artificial indent induced flaws.

Essentially a modified chemical vapor deposition process.

362
Figure 7. Examples of machining flaws in MgF2 from diamond sawing. A and B

show the fracture surface in the vicinity of the fracture origin (upper

part of photos) and part of the machined tensile surface (lower portion

of photos) of dense hot pressed optical grade MgF2- Note the deeper

sawing grooves at ~ 45° to the tensile axis and that each of these appears to have a machining flaw associated with it on the fracture surface

e.g. note flaws 1 and 2 associated respectively with sawing scratches

a and b in both photos. The larger flaw, 1, is clearly the origin of

failure. Note that this flaw is somewhat less elongated than average

for being ~ parallel with the machining groove. Fracture stress ~ 59 MPa

(8,600 psi)

363
!

i

i

1

Figure 8. Further examples of flaws from diamond sawing at fracture origins

of MgF2° A and B show another example of an elongated flaw along one of

the deeper sawing grooves at ~ 45° to the tensile axis as the source of

failure. C and D show another sawn sample of the same hot pressed optical grade MgF2 failing from sawing flaws where the sawing grooves aire ~

parallel with the tensile axis. This illustrates the more complex character of flaws that sometimes can be encountered. Failure appears to

have occurred from two flaws; i.e. from part of a flaw parallel to the

grinding groove (1 in D) and part, or all, of a flaw ~ perpendicular to

the grinding groove. The apparent intersection of two such flaws at

the fracture origins is rare but illustrates some of the complications

that can arise.

364
4 -2

Figure 9. Examples of flaws from sanding CNTD SiC. A and B represent lower

and higher magnifications of the fracture origin in a sample sanded ~ parallel with the tensile axis. Note the nearly semi-circular nature of the

periphery of the flaw causing failure and its curved character. Note also

that it is associated with a somewhat discontinuous sanding groove, (1).

Note also that each of the other sanding grooves marked 1 through 5 in A

has a flaw associated with them, but smaller than the flaw, F, which

caused failure. Note also that some of the other flaws are parallel with

the grinding groove, e.g. the flaw associated with scratch 3, and hence

results in a short vertical step on the fracture surface. A number of

cases of this type of behavior are observed for machining falws, further

reinforcing the concept of the two flaw population observed here. C, D,

E, and F are photos of flaws at the origin of different specimens. Note

the generally progressively more elongated character of the flaw or portion of the flaw exposed on the fracture surface as the angle of the

scratch forming the flaw increases relative to the tensile axis. The

much smaller flaw sizes in this material may reflect in part its extremely

fine grain size, e.g. 0.2 fjjo., but may also be significantly affected by

the apparently fairly high residual stress that may exist in these bodies.

They, however, provide an excellent opportunity for observing such scratches

365
and the resultant flaws not only because the fine grain size allows more

definitive fractography , but also because they have a near pristine surface to begin with, i„e. producing strengths of the order of up to or

exceeding 3,500 MPa (500,000 psi). The failure stress of the specimens

were as follows: A, B - 1170 MPa (170,000 psi); C - 986 MPa (143,000

psi); D - 778 MPa (113,000 psi); E - 606 MPa (88,000 psi); F - 220 MPa

(32,000 psi)

.

which is predominately of degree rather than of a kind on the flaws generated

The third other form of abrasive machining, diamond polishing, has been

investigated some. Observations have been made on materials polished in a

typical fashion, i.e. with the specimen mounted on a disk which is allowed to

rotate while the polishing wheel rotates. This results in a net pattern of

abrasive particle motion over the surfaces which is essentially random as

sketched in figure 10. Polishing flaws observed at fracture origins of

single and polycrystalline materials studied to date are highly elongated

flaws, e.g. figures 11 and 12, consistent with all of the above results in

the following fashion. As a result of the essentially random motion of the

abrasive particle over the surface, one should always have some of the most

severe flaws, i.e. those forming parallel with the direction of motion of

abrasive particles at or near perpendicularity to the tensile axis of the

resultant test bar so that these elongated flaws would dominate failure as

observed. It is also consistent with the fact that one has to do rather

extensive polishing to obtain significantly higher strengths than can be obtained by diamond grinding parallel with the tensile axis. This results because

of differences in the consistency of abrasive particle motion. Although the

flaws from the typical polishing operation are somewhat shallower than those

from grinding, strengths of polished specimens are typically being controlled

by the more severe set of flaws generated by abrasive motion because of the

essentially random direction of abrasive across the specimen. On the other

hand, specimens ground parallel to the tensile axis have their strength controlled by the less severe population of flaws generated by the consistent

motion of abrasive particles across the specimen.

Two factors should be noted about the polishing flaws observed. First,

their periphery tends to be much smoother in contrast to the more irregular

POLISHING WHEEL SPECIMEN

Figure 10. Schematic representation of the common polishing operation and

resultant particle and scratch patterns on the surface of the specimen.

Typically, these will approach a random pattern; hence the most severe

flaws, i.e. those formed essentially parallel with the particles will

dominate failure, even in uniaxial stressing.

366
ligure 11. Examples of ceramic polishing flaws. A and B show the area of

fracture and the fracture origin (arrows) of a fine grain polycrystalline

body (hot pressed, optical grade MgF2) as a result of polishing. This

sample was tested in biaxial flexure, failure stress ~ 91.6 MPa (13,300

psi). Note the elongated character of the flaws as well as their relatively smooth periphere, in contrast to the commonly much more irregular

shape of such elongated flaws from grinding. See figure 17 for another

similar example. C and D show the fracture origin area and fracture

origin (arrows) of a polished specimen of A^C^-rich MgAl2C>4 (Verneuil)

single crystal.

nature of elongated flaws from grinding. As shown in Sections 3 and 4.3,

the elongated flaws from grinding are often more irregular, e.g. due to

their being made up of a series of smaller flaws of varying degrees of overlapping character. This smoother character of polishing flaws may well represent either, or both of, two aspects of the polishing operations relative

to the grinding operation. The much slower motion of the abrasive particles

may directly be the cause of the much smoother nature of the flaw periphery

due for example to less variation in friction. On the other hand, the slower

speed of the abrasive particles may indirectly be effectively allowing more

opportunity for stress corrosion or other phenomena to result in a smoother

crack periphery during machining flaw formation. Secondly, it should be

noted that while there are some differences between grinding and polishing

flaws, overall the general nature of the flaws introduced appears to be similar. This again suggests that the velocity and forces associated with different abrasive processes are not a major factor in a character of the flaws

that are introduced, but only result in varying degrees of the flaw parameters, e.g. their depth and smoothness.

4.3. Machining Flaw Distributions and Strength-Size Effects

An important set of questions is: 1) how variable are the flaws from

machining in terms of size, shape, and orientation; and 2) how do such variations compare with the observed dependence of strength of specimens on their

size and shape due to the differences in the amount of surface or volume

under significant stress. While much yet needs to be done in terms of detailed studies of the variation and statistical distribution of the different

367
Figure 12 „ Examples of variation of grinding flaws in hot pressed, optical

grade MgF2. A and B are fracture photos of the area of origin and the

failure causing flaw (arrows) from grinding parallel with the tensile axis,

fracture stress ~ 64 MPa (9,300 psi). Note the unusually elongated character for a flaw forming perpendicular to the direction of grinding. This

may be in part due to the fact that this flaw is next to the edge of the

sample. C and D show the fracture origin and failure causing flaw (arrows) in a sample of the same material ground perpendicular to the tensile

axis; failure stress ~ 50 MPa (7,200 psi) . This shows a somewhat more

exaggerated, but not extreme, example of the type of variation that one

can have in flaws forming essentially parallel with the grinding groove.

types of machining flaws, some important observations can be made at this

time o

A variety of statistical variations of flaws are observed. One of the

most frequent and obviously expected variations of flaws forming parallel

and ~ along the bottom of the grinding grooves is their extent and regularity.

Basically, these often appear to be not one flaw, but a series of flaws that

partially or totally overlap, e.g. see figures 2,3,4,8,9, and 12. For flaws

of the type characteristically perpendicular to the grinding grooves, a limited but definite number are observed forming at a variety of angles to the

grooves, e.g. figure 13. In finer grain bodies, these angular variations

represent sporatic variations of the material or variations of the local machining conditions. However, in larger grain bodies, and single crystals,

especially non-cubic ones where planes of the same cleavage system may not

be orthogonal, orientation of preferred fracture planes can also be important, figure 13C, D.

Also, some flaws of the type typically forming perpendicular to the

grinding grooves may also be wider, i.e. more eccentric than ~ half penny

flaws. Again local variations in grinding forces, e.g. a deeper grinding

groove or flaw forming between grinding grooves (e.g. see reference 20, 21),

as well as local variations in material properties. Another variation, especially with flaws of the type forming perpendicular to the grinding grooves

are two to several ~ concentric markings before the final flaw boundaries,

368
Figure 13. Examples of variation in flaws forming approximately perpendicular to the direction of grinding. A and B are fracture photos of the area

of origin and the fracture origin arrows in a specimen of fine grain,

dense mullite, fracture stress ~ 206 MPa (30,000 psi). Note that the flaw

is at a substantial angle and has substantial curvature relative to the

fracture plane and the tensile axis. While some such flaws may actually

be the continuation, or a curved end, of a flaw forming parallel with the

grinding direction, clearly many of them are not and are simply variations

of the angle of flaws forming perpendicular to the direction of grinding.

This sample represents an extreme misorientation of such flaws. Note also

that there is no obvious deep grinding groove associated with this flaw

and that there is a definitive grinding groove to the left of it in B.

C and D are photos of the area of fracture origin in dense, hot pressed

SiC, failure stress ~ 372 MPa (54,000 psi). While the specific fracture

origin and the machining flaws are both less well-defined, 1 is a fairly

probable machining flaw and 2 is a possible machining flaw in D. Note

that these are ~ equal to the grain size and that neither is associated

with obvious grinding grooves. Note also the substantial angle of 1

relative to the direction of grinding and the tensile axis may be as

much or more due to the orientation of the grain in which it formed,

e.g. due to the orientation of a preferred cleavage plane in that grain,

rather than a statistical variation in the machining or material parameters governing the formation of cracks.

369
e.g. figure 6. Such observations, which are more common than shown in the

examples in this paper indicate that the flaws are often formed in stages by

some degree of oscillating stresses, or with other vibration.

Another important variation in machining flaws is the formation of edge

flaws, such as along the edges of rectangular test bars or edges of components, e.g. trailing edges of turbine blades. While fracture origins from

sample edges are generally readily identified, specific flaws are not always

evident, e.g. figure 14. However, a substantial number have been identified,

100/im C

I 1

mm

Figure 14. Examples of ill-defined edge flaws at fracture origins in fine

grained ceramics. A and B are fracture photos showing the origin of a fine grain Zr02 body partially stabilized with 8 w/o Y2O3, fracture

stresses 276 MPa (40,000 psi) . C and D are similar photos of fine grain,

dense mullite, failure stress ~ 296 MPa (43,000 psi). Note that the speci

mens failed from one of their edges are quite definite and usually fairly

readily determined but the specific flaw size and shape often cannot be

defined, though first approximation would simply be to use the approximate dimensions of the edge chip in A and B, or lip in C and D as a flaw

dimension. Such fracture origins are quite common for specimens failing

from edges.

370
some of which may be close to a quarter ellipse or circle, but there can be

a variety of shapes as indicated in figure 15 . An example of an edge flaw

in hot pressed Si3N4 is shown in reference 21 . Some indications of the frequency with which such flaws can occur are given elsewhere [4], It has also

been shown that while such edge flaws may be more severe, and hence giving

lower strengths, this is clearly not universally so.

An important question is what are the spatial and size distributions of

machining flaws. At present, only partial indications of these distributions

exist. Thus, for example, fracture surfaces have been examined for machining flaws other than the one from which failure initiated. Such examinations

are most definitive in the smooth, fracture mirror, area since machining

flaws, especiallly smaller ones, will generally be partly or totally obscured

in the hackle and crack branching regions of fracture [12], Thus, more area

is seen on weaker specimens which could bias the results. However, such observations do give a preliminary idea of flaw distributions, e.g. other machining flaws 10-50% the size of the one causing failure have been reported

within ~ 100-300 jum along ground surfaces of fine grain MgF2 and AS2S3

glass [4], Similarly note other machining flaws near the fracture origins

of specimens in figures 7-9. Finally, one other preliminary indication of

flaw distributions are slow crack growth observations in hot pressed Si3N4

(HPSN). At elevated temperatures, flaws begin to open up on tensile surfaces

of HPSN due to local deformation, presumably predominantely by grain boundFigure 15. Examples of specific failure causing flaws at specimen edges.

A and B are photos of the area of origin and specific fracture origin

(arrows) in a single crystal of CaO stabilized Zr02, failure stress 105 MPa

(15,300 psi). C and D are fracture photos of the area of origin and actual origin (arrows) of a failure causing flaw at the edge of a hot pressed MgF2

.

371
ary sliding. These opening cracks appear to often originate from larger machining flaws (some cases of this have been verified [21]). Therefore, the

distribution of these opening cracks gives some idea of the surface density

of flaws, e.g. figure 16. While controlled, e.g. side, lighting, or etching

Figure 16. Example of slow crack growth opening of machining flaws. A and

B show hot pressed Si3N4 (HS-130) tensile creep rupture tested at ~ 87 MPa

(12,500 psi) at 1300OC (failure strain ~ 13.8 x 10"3

). Note the ~ linear

features developing on the surface, shown at higher magnification in B.

Machining flaws such as those identified as the source of fracture at

lower temperatures (e.g. fig, 18A, B) have been identified as a major

source of such flaw growth L2l],

of machined transparent bodies, especially of glasses and crystals can be

important tools in better defining machining flaw populations, opening of

flaws due to slow crack growth may have fairly broad applicability. Thus, a

variety of AI2O3 and other commercial ceramic bodies, such as crystallized

glasses [33] are known, or expected, to exhibit slow crack growth that could

be exploited for further definition of machining flaw populations.

The studies thus far of this paper have been made mainly on bars ~ 1.3 x 2.6 x 13 mm. While much strength-specimen size dependence arises

from processing defects [5,6,34,35], the observed size dependence of the

strength of ceramic material such as glasses and dense polycrystalline bodies

where machining flaws should dominate would suggest that there is a size

dependence to the maximum flaw size introduced by machining samples of different sizes. In order to specifically test this, large bars of optical

grade, fully dense, hot pressed MgF2 were ground on one side parallel to the

tensile axis and on the other perpendicular, then tested. Sections of these

large MgF2 bars after flexural testing were then cut up into smaller test

bars whose tensile surfaces were the original ground surfaces of the large

bars so we were testing the same flaw population. The results of these

tests along with some data on different sets of B4C specimens are shown in

table 4.

The results of table 4 suggest a possible effect of the specimen size

on the size of machining flaws controlling strengths, but several complications occur. First, since the same head travel rate was used, the strain

rates were lower for larger bars allowing more slow crack growth to occur;

thus strain rate is being held constant in other tests of MgF2 and S1O2

based glasses now underway. Second, processing defects (and possibly variations in B4C) play a larger role in the failure of larger bars, limiting the

machining results. Third, both MgF2 and B4C have internal stresses due to

their non-cubic crystal structures which limit the increases in strength

with decreasing flaw size [36], Thus, while the strength results do not

372
strongly show a size effect, the flaw observations more strongly indicate

such an effect „ In order to minimize these complications, additional tests

are now under way on SiC>2 based glasses. However, initial results have been

complicated by different types of flaws being common sources of failure in

many of the large bars. These flaws represent either handling damage (which

could be much more severe in larger, i.e. more massive, bars), or a new, much

more limited population of machining induced flaws, e.g. from such possible

effects such as damage from unusually large swarf particles trapped between

the specimen and grinding wheel.* Other indications that a size effect for

machining flaws may exist, but may generally be limited, are shown by fracture origins of components or prototypes, e.g. figures 17, 18, and references 5,16,18, and 19. These suggest somewhat deeper flaws than in test bars

but specimen shape may also be a factor. Thus, the data of Bansal and Duckworth on Pyroceram 9606 [34] reporting no size effect of machining is not

necessarily at conflict with the preliminary studies. Either their material

or machining may have produced limited changes in machining flaw sizes not

discernable in the limited number of observations made. (Again, possible

effects of internal stresses could limit strength differences between different size bars with different size flaws.)

Figure 17. Photos of the area of fracture origin and actual origin (arrows)

of a polished, IR dome, MgF2 failed in a simulated aero-thermal heating

test. The failure stress of ~ 69 MPa (10,000 psi) estimated from the observed flaw size and the measured fracture energy and Youngs modulus is

in excellent agreement with the failure stress predicted from thermal

analysis. Note the similarity in size and shape of this flaw in a polished component to that shown earlier in figure 11 in a polished MgF2 disk

Another possible source of different flaws are those formed at the edge of

the grinding wheel on larger bars that are wider than t he wheel in contrast

to small bars which are narrower than the wheel.

373
Figure 18. Examples of machining flaws at fracture origins in hot pressed

Si3N4 (NC-132). A and B are photos of the area of fracture origin and

the fracture origin (arrows) of specimen ground circumferentially for

true tensile testing; fracture stress 312 MPa (45,300 psi). B shows

a portion of the highly elongated flaw which is about average in definition for such flaws in Si3N4» Other clearer examples of these types of

flaws can be found in references 20 and 21. C and D show the fracture

origin area and fracture origin (arrows) in a simple prototype SigN4

turbine blade spun tested to failure at room temperature. The specimen

(courtesy of Mr. Dave Richerson, Garrett Airesearch) failed in the attachment area where stresses are highest and machining is perpendicular to

the blade (tensile) axis; failure stress was ~ 483 MPa (70,000 psi). Note

the Richerson and Yonushonis [3 7] have shown similar, clearer examples of

such fracture origins in failed turbine blades.

5. Summary and Conclusions

Fractography has been used to identify machining flaws controlling the

strengths of a variety of ceramic bodies. This shows that sawing, grinding,

sanding, and polishing all introduce predominately two populations of flaws

extending in from the surface. One set, those giving the lowest strengths

are formed essentially parallel with the machining grooves formed by the grit

particles and are generally fairly elongated. The other set of flaws, those

typically giving higher strengths, are typically formed perpendicular to the

machining grooves and are generally less elongated, often approaching a semicircular periphery. The differences in size and shape in these two sets of

flaws combined with their orientation relative to the direction of machining

are the predominate cause of the anisotropy of strength as a function of the

stressing direction relative to the direction of machining.

Variations in the depth, shape, and orientation of the above sets of

flaws have been observed, as well as effects of composition and microstructure. While some variations in flaw orientation appear to be sporatic in

larger grain bodies where flaws are similar to or less than the grain in

size, and especially in single crystals, the orientation of preferred cleavage planes can be an important factor. The set of generally elongated flaws

that typically form parallel with the machining grooves vary from a continuous flaw to a series of partially overlapping flaws. Flaws forming generally

perpendicular to machining grooves often have a fan and curved character,

e.g. like half a clam shell. No obvious effect of body composition is seen

on the size, i.e. depth, of either set of flaw introduced, and decreasing

grain size has at best a limited effect on reducing the size of the strength

controlling flaws introduced. Thus, the typical variety of machining processes commonly introduce flaws in the range of 20-50 /jm depth for both sets

374
of flaws in glasses, single crystals, and typical grain size dense polycrystals in most laboratory size test bars. Limited evidence, as well as the

statistical nature of machining flaws, both suggest that the size of machining

flaws controlling strength increases some with specimen or component size.

However, increasing competition from other flaws, e.g. processing flaws or

flaws introduced by handling processes limit such machining flaw sizespecimen size evaluation.

TABLE 4

SPECIMEN SIZE-STRENGTH1 AND FLAW SIZE EFFECTS IN GROUND BARS

Ground Ground 2

Strength Flaw3 Strength Flaw3

ksi MPa a (urn) b/a ksi MPa a(um) b/a

MgF? (KODAK. HOT PRESSED. ~ 0 POROSITY. GRAIN SIZE ~ 0.5 urn)

Span 4 cm, cross section 1 x 1.8 cm:

11±2 76±14 75±20 4±2 12±2 80±14

(4) (4) (3)

Span 1.27 cm, cross section 0.51 x 0.25cm:

12±1 80±7 42±3 ~ 4 16±3 110±20 57±11 0.8±0.3

(4) (3) (9) (3)

B4C (NORTON. HOT PRESSED. < 1% POROSITY. GRAIN SIZE ~ 10 urn)

Span 2.54 cm, cross section 1.27 x 0.60 cm:

35±4 240±27 100±80 1.1±0.2 44±6 304±40 60±40 1±0.5

(4) (3) (4) (3)

Span 1.27 cm, cross section 0.62 x 0.24 cm:

34±12 235±80 18±3 3±1 49±6 338±40

(7) (3) (7)

B4 C (FIBER MATERIALS, HOT PRESSED, < 1% POROSITY. GRAIN SIZE ~ 2 jjm)

Span 1.9 cm, cross section 0.51 x 0.25 cm:

19±2 130±14 29±11 4±1 47±12 325±80 13±8 2±1.8

(5) (4) (5) (3)

Strengths measured at 23°C in 3 point flexure with a head travel rate of ~ 1.27 mm/min.

1 and II refer respectively to the direction of grinding relative to the

bars' tensile axis. Lower numbers in parentheses give the number of values.

The fewer number of values for flaw parameters than strengths reflect

failure from processing defects or specimen edges.

a = flaw depth, b = flaw length along surface.

375

ft
References

[ l] Rice, R. W„, Panel Discussion, in The Science of Ceramic Machining and

Surface Finishing , NBS Special Pub. 348, 407-408, S„ J„ Schneider

and R. W. Rice, eds. (Govt. Printing Office, Wash., D.C. 20402, 1972).

[ 2] Mecholsky, J. Jo, Freiman, S. W., and Rice, R. W., Effect of Grinding

on Flaw Geometry and Fracture of Glass, J. Am. Ceram. Soc. [60] 3-4

(1977)

.

[ 3] Mecholsky, J. J e , Freiman, S. W., and Rice, R. W., Effect of Surface

Finish on the Strength and Fracture of Glass, Proc. of the Xlth Int.

Congress on Glass, Prague, Czechoslovakia (1977).

[ 4] Rice, R„ W., Machining of Ceramics, in Ceramics for High-Performance

Applications , 287-243, Proc. of 2nd Army Materials Technology Conf.,

Hyannis, Mass (1973).

C 5] Rice, Ro W„, Mecholsky, J. J., Freiman, S. W. and Morey, S. M., to be

published in Proc. in Quantative NDE Conf. , Cornell Univ., Air Force

Report #AFML-TR-77 (1977).

[ 6] Rice, Ro W., Processing Induced Sources of Mechanical Failure in Ceramics, in Proc. of Conf. on the Processing of Crystalline Materials,

H. Palmour III, R. F. Davis, and T. M. Hare, eds., 303-319 (Plenum

Press, 1978).

[ 7] Rice, Ro Wo, Microstructure Dependence of Mechanical Behavior of Ceramics, in Treatise on Materials Science and Technology, Vol. 11 , Properties and Microstructure, 199-381 (Academic Press, N.Y., 1977).

[ 8] Freiman, S. W., Mecholksy, J, J., Rice, R. W. and Wurst, J. C, Influence of Microstructure on Crack Propagation in ZnSe, J. Am. Ceram. Soc.

[58] 9-10 , 406=409 (1975).

[ 9] Rice, R. W. , Mecholsky, J. J. and Becher, P. F., The Effect of Grinding

Direction on Flaw Character and Strength of Single Crystal and Polycrystalline Ceramics, to be submitted to J 0 Am. Ceram. Soc.

[10] Rice, R. Wo, Machining, Surface Work Hardening, and Strength of MgO,

Jo Am. Ceram. Soc. [56] 10, 536-541 (1973).

[11] Rice, R. Wo, The Effect of Grinding Direction on the Strength of 1

Ceramics, in The Science of Ceramic Machining and Surface Finishing , NBS Special Pub. 348, 365-376, S. J. Schneider and R. W. Rice, eds.,

(Govt. Printing Office, Wash., D. C. 20402, 1972).

['12 ] Rice, Ro W. , Fracture Topography of Ceramics, in Surfaces and Inter - faces of Glass and Ceramics , 439-472, V. D. Frechett, W. C. LaCourse,

and V. L. Burdick, eds. (Plenum Press, N. Y. 1974).

[13] Kirchner, H. P., Gruver, R. M. and Walker, R. E. Strength Effects

Resulting from Simple Surface Treatments, in The Science of Ceramic

Machining and Surface Finishing , NBS Spec. Pub. 348, 353-363, S. J 0 Schneider and R. W. Rice, eds. „ (Govt. Printing Office, Wash., D. C.

20402, 1972).

[14] Richerson, D. W., Schuldies, J. J., Yonushonis, T. M. and Johansen,

K. M., ARPA/Navey Ceramic Engine Materials and Process Development

Summary, in Proc. of 5th Army Materials Technology Conf. on Ceramics

for High Performance Applications-II , J. J. Burke, E. N. Lenoe and

R. N. Katz, eds. (Brook Hill Pub. Co., Mass. 1978).

376
[15] Gielisse, P. J., Kim, T. J. and Choudry, A., Dynamic and Thermal Aspects

of Cemmic Processing, Naval Air Systems Command Contract Report,

15 Nov 1970-15 Nov 1971, Univ. of Rhode Island, Ingston (Dec. 1971,

AD 740 833).

[16] Tressler, R. E . , Langensiepen, R„ A. and Bradt, R. C, Surface-Finish

Effects on Strength-vs-Grain Size Relations in Polycrystalline Al 203,

J. Am. Ceram. Soc. [57] 5, 226-227 (1974).

[17] Bradt, R„ C. , Dulberg, J. L. and Tressler, R. E. , Surface Finish

Effects and the Strength-Grain Size Relationship in MgO, Acta Metallurgic [24] 529-534 (Pergamon Press, Great Britain, 1976).

[18] Cranmer, D„ C. , Tressler, R. E. and Bradt, R. C. , Surface Finish Effects and the Strength-Grain Size Relation in SiC, J. Am. Ceram. Soc.

[60] 5, 230-237.

[19] Bradt, R„ C. and Tressler, R. E., Micros tructure/Surface Finish Control of the Strength of Polycrystalline Ceramics, Proc. of 6th Int.

Mat. Symposium on Ceramic Microstructures , 76, R. M. Fulrath and

J. A„ Pask, eds., 785-795 (Westview Press, Boulder, Colorado, 1976).

[20] Rice, R. W., Freiman, S. W., Mecholsky, J. J., Ruh, R. and Harada, Y.,

Fractography of Si3N4 and SiC, in Ceramics for High Performance Appli- cations-II , Proc. of 5th Army Materials Technology Conf., J. J. Burke

E.N.Lenoe and R.N.Katz, eds. (Brook Hill Publishing Co., Chestnut

Hill, MA, 1978).

[21] Rice, R. W., Freiman, S 0 W. and Mecholsky, J. J 0 , Fracture Sources in

Si3N4 and SiC, Proc. of 1977 DARPA/NAVSEA Ceramic Gas Turbine Demonstration Engine Program Review, 665-688, J. I. Fairbanks and R. W.

Rice, eds. Maine Maritime Academy, Castine, Maine, 1-4 Aug 1977

(MCIC 78-36)

[22] Rice, R. W„, Machining Flaws and the Strength-Grain Size of Ceramics,

this conference.

[23] McDonough, W. J„, Flinn, D„ R. , Stern, K. H. and Rice, R. W„, Hot

Pressing and Physical Properties of Na Beta Alumina, J. Mat. Sci„, 13,

2403-2412 (1978).

[24] Richerson, D. and Yonushonis, T., in Ceramic Gas Turbine Engine Demon- stration Program , Interim Report #4, 4-8 to 4-10, Airesearch Report

#76-212188(4), Contract #N00024-76-C-5352 (1977).

[25] Bruche, E„ and Poppa, H„, Zur Kenntnis der Glasoberflache , Electron

Microscopy, 336-338, F. S. Sjostrand and J. Rhodin, eds. (Academic

Press, New York, 1956).

[26] Imanaka, 0., Lapping Mechanics of Glass-Especially on Roughness of

Lapped Surface, Annals of the C.I.R.P., Vol. XIII, 227-233 (Pergamon

Press, 1966).

[27] Allgeyer, G. H. and Colwell, L. V„, Aspects of Machining Glass-Ceramic

Materials, in The Science of Ceramic Machining and Surface Finishing , NBS Special Pub. 348, 53-57, S. J„ Schneider and R. W. Rice, eds

(Govt. Printing Co., Wash., D.C. 20402, 1972).

[28] Brian, L, and Wilshaw, R, , Review Indentation Fracture: Principles and

Applications, J. Mat. Sci. [10], 1049-1081 (1975).

[29] Evans, A. G., Abrasive Wear Damage in Brittle Materials, this

conference.

377
[30] Doremus, R. H. , Johnson, W. A., Depths of Fracture-Initiating Flaws

and Initial Stages of Crack Propagation in Glass, J. Mat. Sci„, 13,

855-858 (1978).

[31] Mecholsky, J. J. Fractographis Analysis of Delayed Failure in MgF£,

submitted to J. Am. Ceram. Soc.

[32] Dutta, S., Rice, R. , Graham, H. and Mendiratta, M., Characterization

and Properties of Controlled Nucleation Thermo Chemical (CNTD) Deposited Silicon Carbide, to be published.

[33] Govila, R. K., Kinsman, K. R. and Beardmore, P., Fracture Phenomenology of a Lithium-Aluminum-Silicate Glass-Ceramics, J. Mat. Sci„, 13,

2081-2091 (1978).

[34] Bansal, G, K. , Duckworth, W. H. and Niesz, D. E. , Strength-Size Relationships in Ceramic Materials: Investigation of a Commercial GlassCeramic, J. Am. Ceram. Soc. [55] 3, 289-308 (1976).

[35] Bansal, G. K., Duckworth, W„ H. and Niesz, D.E., Strength Size Relationships in Ceramic Materials: Investigation of an Alumina Ceramics,

J. Am. Ceram. Soc. [59], 11-12 , 472-478 (1976).

[36] Rice, R. W., Pohanka, R. C. and McDonough, W, J 0 , The Effect of Stresses

from Thermal Expansion Anisotropy, Phase Transformations and Second

Phases on the Strength of Ceramics, submitted for publication.

[37] Richerson, D, and Yonushonis, T. , in Ceramic Gas Turbine Engine Demon- stration Program, Interim Report #5 , 3-9 to 3-11, Airesearch Report

#76-212188(5), Contract #N00024-76-C-5352 (1977).

Note added in Proof

Two factors, slow crack growth and residual stresses, not addressed in

the text should be briefly considered. While slow crack growth might increase

flaw sizes by 50-100% in the extreme, this would not change the strength anisotropy unless the growth was significantly different for the two populations of flaws. The generally good agreement between measured strength anisotropy and that calculated from observed flaws argues against significant

differences in flaw growth. Further, fracture energies calculated from observed flaw sizes generally agree with measured values [9,20-23,36] unless

altered by other factors such as thermal expansion anisotropy [36] , indicating flaw growth in most of these materials is limited, e.g. < 50%. If

residual stresses in the surface from machining were compressive in nature,

they would compensate, at least in part, for slow crack growth. On the other

hand, if residual stresses are tensile, e.g. as indicated in recent indentation tests (private communication with Dr. Brian Lawn) then their effects

would be additive to those of slow crack growth. Again, the above results

indicate that any additive effects are limited. Thus, while further study

of slow crack growth and residual stress are important to refine our understanding of machining flaw-mechanical property relations, they do not appear

to require major changes in the results and conclusions of this study.

378
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

ACOUSTIC SURFACE WAVE MEASUREMENTS OF SURFACE CRACKS IN CERAMICS

B. T. Khuri-Yakub 1

, A. G. Evans 2, and G. S. Kino 2

Edward L. Ginzton Laboratory

Stanford University

Stanford, CA 94305

Materials Department

University of California

Berkeley, CA 94720

Acoustic surface wave techniques for detecting and characterizing surface

cracks in ceramics have been devised. The techniques have been demonstrated

to detect cracks at least as small as 60 ym in depth, on polished samples. The

crack detectability was diminished by surface grinding, because of a background

effect attributed to the backscattering from surface microcracks. Preliminary

experiments conducted on turbine blades indicated that the techniques can be

applied to components of complex shape.

Key words: Acoustic surface wave; ceramic; surface characterization; surface

crack; turbine blade.

1. Introduction

Surface cracks are an omnipresent source of fracture in structural ceramics [1-3]. The

cracks are generated during finishing operations (such as grinding and machining) and usually

consist of arrays of semi -ell iptical cracks with an essentially random inclination to the

surface but a preferred alignment parallel to the direction of motion of the abrading

particles [4]. The characterization of these cracks is an essential constituent in an effective failure prediction procedure. The characterization techniques that have been conven- tionally applied to surface cracks are dye penetrants and fracture statistics [5]. The dye

penetrant techniques are restricted to relatively large cracks because of inherent resolution

limitations and are only viable in structural ceramic systems as preliminary inspection procedures. Statistical methods exhibit superior prospects [1,5]. But the statistical approach

is not, by itself, capable of identifying components subject to anomalously severe surface

damage on an individual basis. Alternative, or additional, surface crack characterization

techniques must, therefore, be sought. An acoustic surface wave technique that has the basic

pre-requisites for satisfying the failure prediction requirements is described in the present

paper. The technique has been demonstrated to give good quantitative predictions of flaw

size and fracture stress in glass samples. The experiments on hot pressed silicon nitride

ceramics are at an earlier stage but indicate that surface cracks as small as 60 ym radius

can be detected.

The scattering theory that provides the basis for pertinent acoustic measurements is

firstly described, for both the short and long wavelength regimes. Then, the types of transducer designed to provide the requisite acoustic information are described. Thereafter,

specific acoustic techniques for detecting and characterizing surface cracks on both plate

samples and on turbine blades are presented: separate emphasis being afforded to the short

and long wavelength regimes. Finally, the implications of the results for ultrasonic failure

prediction are discussed.

II. Scattering Theory

An acoustic surface wave incident on a crack will be partially reflected. When the

wavelength x of the acoustic wave is comparable to the crack radius a, it will be demonstrated

(eq. 2) that the scattered power P received at the transducer should be proportional to a

2

. Whereas when a << x, P«a 6 A (eq. 7), and when a >> X, P should be independent of the wave379
length, (being roughly equal to its value when a ~\) . Thus for maximum detectability, acoustic surface wave measurements should be made at a wavelength comparable to or smaller than the

crack size. For example, on silicon nitride (surface wave velocity, 5.8 x 10 3 m/s), an acoustic surface wave with a frequency of the order of 100 MHz should optimize the detectability of

a 60 ym crack. Thus it has been necessary to develop entirely new techniques for excitation

of surface acoustic waves on ceramics, using frequencies much higher than those normally

employed in nondestructive testing. Additionally, important information concerning surface

cracks can be obtained in the long wavelength limit, a << x, whenever the cracks are sufficiently large to yield detectable scattering. It is also of interest, therefore, to obtain

information at the lower frequencies (10 - 20 MHz) pertinent to the long wavelength regimes.

Kino and Auld [6,7] have shown that the scattering coefficient Sn for a surface wave impinging at normal incidence onto a surface crack is;

s

ll " f /

Auj*ij"i dA

c O)

]

A

where S n is defined as the ratio of the amplitude of the scattered signal at the terminals

of the transducer to the incident signal; Ac is the crack area; ajj is the incident stress

associated with the acoustic wave in the region of the crack, prior to the introduction of

the crack; Auj is the displacement of the crack surface due to the incident acoustic surface

wave, P is the power injected into the transducer and u is the frequency. This relation will

be used as a basis for predicting scattering characteristics at both short and long wavelengths. Note that the ratio of the scattered power to the incident power, P , is proportional at |S n

2

| . s

II . 1 Short Wavelength Scattering

When a >> x and the width of the transducer w >> a (as would normally be the case), an approximate estimate of Sn can be made. It can be shown by analogy with plane wave scattering [6] that Auj has roughly twice the value of the displacement in the incident acoustic

wave. It follows that

S n

^(vn/wja (2a)

or,

P„ Myn/w) 2

a

2

(2b)

where x is a constant, equal to 2 for plane wave scattering, but < 2 for a crack; and n is an efficiency coefficient relating the input power at the transducer to the power incident on the crack. For a crack located on the central axis of the beam,

n = nTn

D (3)

where nj is the transducer power lost upon converting the input electrical signal to an acous- tic surface wave, and nn. is a diffraction loss term. For a flaw on the acoustic beam axis

located in the Fraunh'dfer region, (i.e., in the far field where z > w

2 /x) at a distance z from the transducer, it can be shown that [8]

n

D

= w

2

/Az (4)

Combining equations (2), (3), and (4), the power loss becomes;

p

s = i-^-i m (5)

A crude estimate of the crack detectability can be obtained from eq. (5). The reflection

coefficient of a 90° corner for a material with a Poisson's ratio of 0.25 is 12 dB [19]:

yielding a x value of ^0.3. Since tw is typically ^0.1 for our transducer (see Section III),

a frequency of 100 MHz and a 1 mm beam width (typical of high frequency transducers) requires

that z > 1.6 cm. Hence for a flaw size of 50 ym, the power loss predicted by eq. (5) is

^54 dB. A 50 ym crack should thus be easily detected using an acoustic surface wave with a

380
wavelength comparable to the crack dimensions. It should also be noted that crack resonances

can be excited, and observed, at short wavelengths [10]. The resonances of the cracks will

change the amplitude of its reflection coefficient in a periodic manner; a fact which has

been neglected in our rough calculation.

I I. 2 Long Wavelength Scattering

When a << x, we might expect that the displacements in the neighborhood of a crack

would be similar to the quasistatic displacement. The displacement Auj of the crack surface

can be estimated from the quasistatic solution for a penny-shaped cracK [7],

Au

ZZ

4(1 -v*)

nE

(a' r

2

)

1/2

(6)

where a Z2 is the stress normal to the crack, E is Young's modulus, r is the distance from the

center of the crack, and v is Poisson's ratio. Insertion of eq. (1) gives the scattering

coefficient at the crack as;

'11

ja)(l-v

2

)a

z2

2

3EP

(7)

Furthermore, by writing a zz

2 in terms of P

zz 8E

R

(7)

-, 2

(v*) TTWl(w)' (8)

and by invoking the transducer efficiency, n. it can be shown that the scattered signal at

the transducer is;

3.

'11

or,

n_

TT^7)w

a 6

/A

4

(9)

(10)

where v», v

s and vr are the longitudinal, shear and Rayleigh wave velocities respectively,

and f is a normalized parameter tabulated by Auld [11]: for typical ceramics f

z ^0.4.

For a 10 MHz wave incident on a 100 ym crack located in the Fraunhbfer zone (z = 1.5 cm),

n D for a beam width of 2 mm (typical for long wavelength transducers) is MD.44. Thus, for a transducer efficiency nj ^0.1 eo.- (9) yields a scattering coefficient Sn ^9.18 x 10 _t

*. This

corresponds to a 60 dB reduction from the input signal. This level should be discernable.

However, the equivalent reduction for a 50 pm crack is 78 dB: a level which would be difficult to detect.

III. The Transducers

III.l Calibration Procedures

The transducers used in this study were calibrated by firstly measuring the total loss

L 0 between two closely spaced transducers, Li + L 2 (where Li = 10 log 10 n-ri, L 2 = 10 logi 0 n T

2'

the subscript T referring to transmission). The losses L~i and Lq2 associated with signals

reflected from a closely spaced corner are then measured. These losses can be combined to

yield values of Li or L 2 in accord with the following relations,

L

l + L2

" L 0

2L

1 + LC

= L

C1

(Ha)

(lib)

2L

2 + L c

= L c2 (lie)

381
It follows that

2L

0

" L C2 + L C1

(12)

III. 2 Transducers Designs

Schematics of three types of transducer configuration suitable for launching acoustic

surface waves on ceramics and glass are shown in figures 1, 2, and 3. The first is a highly

efficient type of wedge transducer for use at frequencies below 10 MHz. A pulsed rf signal

is excited by means of a bulk wave transducer bonded to the wedge material; this bulk wave

is incident at an angle e to the normal at the surface such that

sin = vw

/v

s (13)

TRANSDUCER

WEDGE

iCRACK

I

GtASS f/,

Fig. 1. Schematic diagram of surface crack and test set up with a wedge transducer.

LITHIUM NIOBATE

INTERDIGITAL TRANSDUCER

Fig. 2. A schematic of the surface to surface (STS) transducer,

where vw

is the wave velocity in the wedge material and v

s the surface wave velocity along

the substrate. A new design procedure [12] for wedge transducers predicts that, for good

conversion efficiency, it is necessary to have Pwvw/ps v

s % ( wnere Pw> ps are the mass densities of wedge materials and substrate, respectively). Also, an optimum length of contact

between the wedge and the substrate (^3x1 0~ m) is required to obtain a maximum theoretical

conversion efficiency from bulk to surface waves of 0.81. The measured conversion efficiency

of the current transducers approaches the theoretical value, being as high as 0.5 in the

382
EDGE BONDED

TRANSDUCER

CERAMIC

=POLYETHYLENE

SAW

CERAMIC

amy ^POLYETHYLENE

SAW

Fig. 3. A schematic of a surface to surface (STS) transducer using a ceramic substrate

and excitation with an interdigital transducer on ZnO.

best cases,

9 dB.

For example, the measured loss for an individual wedge transducer is typically

Normally, for the excitation of waves on glass or metals, the wedge material consists

of a RTV silicone gel; this material has an extremely low wave impedance Zw

(=pwvw)- F°r excitation of surface waves on ceramics, which have a higher surface wave velocity, lucite

is a superior wedge material. For reproducibility, it is usually convenient to use a liquid

(such as water or ethylene glycol) between the wedge and the ceramic or glass substrate.

However, at higher frequencies it is difficult to construct suitable bulk wave transducers

and to design a wedge that is sufficiently thin to ensure that acoustic losses in the wedge

medium do not dominate the performance of the system. Thus, a new configuration, shown in

figure 2, has been adopted for exciting very high frequency surface acoustic waves of ^100 MHz.

A surface wave is first excited on a piezoelectric (LiNb0 3 ) delay line by an interdigital

transducer. A fluid couplant is placed between the delay line (outside the transducer region)

and the substrate on which an acoustic surface wave is to be excited. The surface wave

becomes a leaky wave in the presence of the fluid and excites a bulk wave at an angle [13]

(14)

where Vf is the longitudinal wave velocity in the fluid and v< on the Li Nb0 3

. This longitudinal wave is, in turn, arranged i the critical angle e 2 that permits surface wave excitation:

-1

'2

is the surface wave velocity

be incident on the sample at

sin (15)

where v s2 is the surface wave velocity on the substrate of interest. The angle e between

the two solids is, therefore

1 = sin

-1

fe) - •<•"

ft)

(16)

The distance between the LiNb0 3 and the delay line is made as small as possible in order to

minimize the propagation loss in the fluid couplant. For the case of a surface wave on the

Y-Z cut Li Nb0 3 delay line, coupled to Si^ ceramic, the calculated maximum efficiency (0.54) occurs for an overlay distance i (fig. 2) of .5 mm, at a center frequency of 100 MHz. This

constitutes a loss of about 2.7 dB for conversion of the surface wave from one material to

the other. The current transducers yield a loss comparable to this predicted value. Ethylene glycol was used as the fluid couplant because it does not evaporate quickly and because

it forms a good meniscus. The inclination e for this couplant is 13°.

383
Although this type of transducer operates well at frequencies in the 30 to 100 MHz range,

it is still not ideal for the present nondestructive testing application, because of the liquic

medium. An alternate transducer, illustrated in figure 3, is thus under development. An inter

digital transducer is deposited on a substrate material identical to that being tested. Piezoelectric ZnO is sputter deposited onto the transducer to give strong piezoelectric coupling.

A polyethylene layer, a few micrometers thick, is then deposited onto the ZnO to form a com- pliant contact region. With this design, the component and wave guide, having identical wave ! velocities, can be parallel; thereby, permitting the polyethylene layer to be relatively thin

and hence, allowing little acoustic loss.

IV. Low Frequency Measurements

The reflection coefficient for a surface crack obtained at low frequencies can be used to yield a direct prediction of the fracture strength of a brittle material. The requisite result can be derived from the scattering coefficient given in eq. (9). Firstly, the reflection coefficient is expressed in terms of a geometric stress intensity factor ki (=2/aAf for a penny

crack) as follows [14]:

2 \

1/6

'6(1 - v)X

Z |S„|w \

k

< = 2

\ V J

(i7> I

indicating that ki can be obtained immediately from a measurement of the scattering coefficient. It will be noted that the geometrical stress intensity factor ki is proportional to

ISnl 1 / 6

, so that errors in the measurement have a very small influence on the predicted ki

values. Then, since the fracture strength a of a material is related to ki and the fracture

toughness K JC by

K

1C = ak (18)

a determination of Sn and knowledge of yields a direct prediction of o

c

.

Pyrex discs 3 mm thick and 7.6 cm in diameter in the annealed condition were prepared for

fracture testing by introducing small semi -elliptical surface cracks in the center of each

disc. These pre-cracks were made in a controlled way by applying a bending moment of approximately 2 N.m to the disc and, on the side in tension, introducing a Knoop indentation with

a force of 0.5 to IN. This technique produces roughly half penny-shaped cracks with radii

between 100 and 500 pm. Such small cracks were required to limit the crack size to less than

the acoustic wavelength ^9 x 10 m. Additionally, the half penny-shaped cracks may be coaxed

to an extended semi-elliptical shape by applying a subsequent bending moment to the disc.

The strength of each disc was determined in biaxial flexure using the method of Wachtrrjan,

et al. [15]. The discs were loaded to fracture at a loading rate of approximately 60 MPas" . The loading rate was chosen to minimize possible slow crack growth effects [16]. To further

insure against environmental influences, the specimens were subjected to a vacuum for one hour prior to strength testing, and exposed to a stream of dry nitrogen gas for three minutes

prior to and during the flexure test.

Once the specimens were fractured, the geometry of the pre-cracks was studied by examining

the fracture surface using reflected light. The aspect ratio (a/c) of the semi-elliptical

surface pre-cracks could then be easily measured.

The stress intensity factor for the pre-crack was evaluated by the method of Shah and

Kobayashi, [17],

MRaR

/rra

k

where E(k)

i -fro- m

is an elliptic integral of the second kind, M

B is a magnification factor which

takes into account the aspect ratio and the proximity of the crack depth to the neutral

axis, a

B is the maximum bending stress at the surface of the specimen, and a is the length

of the semi -mi nor axis (depth) of the pre-crack.

384
The ten samples tested gave a normally distributed set of fracture toughness measurements,

and the mean value of fracture toughness, Kj

c

, at a confidence level of 90 percent was:

0.76 ±0.04 MPa m

1 ' 2

, comparable to the values obtained by Wiederhorn [18].

The experimental arrangement for measuring the acoustic surface wave reflection coefficient |S n | is shown schematically in figure 1. A wide band, high efficiency wedge transducer was used to excite and detect the surface acoustic waves. The transducer has a center

frequency of 3.4 MHz, a bandwidth of 50 percent, and a one way insertion loss of 9.2 dB.

Calibration of the transducer was conducted as described in Section III. All measurements

were taken in the far field of the transducer (z > x

2 /x).

A set of acoustic measurements were taken with the front of the wedge, distance z = 2.3

cm from the crack and with the acoustic surface wave normal to the crack surface. The mea- surements were taken in the bending fixture, using a bending moment of approximately 2 N.m

to ensure that the crack was open.

A series of reflection measurements was conducted on ten different cracks, and kr calculated from eq. (17). In figure 4 the acoustic prediction of the fracture stress o

c is

plotted against the measured values of o

c

. A least squares fit of this data indicates a linear relationship between theoretical and experimental measurements with a slope of 0.85.

We see that for larger, more irregular cracks (where the assumption of the theory that the

diameter of the crack is less than 1/4 wavelength is not well satisfied), the error between

theory and experiment increases. It is interesting to note that the effect of the ellipticity

of some of the small surface pre-cracks appears not to affect the accuracy of the acoustic

prediction of kj max to any appreciable degree. This confirms a prediction of Budiansky and

Rice [19] that the effect, if any, should be less than 10 percent.

30 40 50 60 70 80 90

<rCA (MPa

)

Fig. 4. Actual fracture stress versus predicted fracture stress.

V. High Frequency Defect Detection

V.l Measurements on Flat Surfaces

The material selected for this study was a commercial, hot-pressed silicon nitride

(NC 132). Plate specimens (3 x 3 x 0.5 cm) were machined from an as-pressed billet and one surface optically polished. Specific surface grinding treatments were then applied to each

sample using 30 ym and 70 urn diamond abrasive pastes and a standard diamond wheel. Then,

385
three surface cracks were introduced using the Knoop indentation technique [13]. The sizes

of the cracks were characterized using optical or scanning electron microscopy (fig. 5); the

results are presented in figure 6. Subsequently, the plastic zones were removed by grinding 1

(the grinding condition corresponding to that applied prior to indentation). Thereafter, the

cracks were not detectable, using either optical or electron microscopy, because the thin

layer of surface plasticity that accompanies grinding tends to obscure the residual crack

opening. It is assumed that all cracks are semicircular in shape, in accordance with observations of Petrovic and co-workers [20].

i 4 <V

Fig. 5. Micrographs of surface cracks introduced by indentation:

(a) scanning electron micrography, P = 20 N,

(b) polarized light reflected micrography, P = 50 N,

(c) reflected light optical micrograph, P = 110 N.

The surface roughness of each sample was also measured using a Datatrak instrument.

The results are summarized in figure 7. The expected increase in the amplitude of the sur- face irregularities with the grinding severity is apparent. It is also evident that

uncharacteristically deep troughs exist in the sample subjected to a 30 ym grind (the conse- quence of a poorly sequenced grinding procedure) -- the significance of this feature will

emerge later.

x The plastic zone in Si 3 Nif is relatively shallow, and its removal does not effect a significant change in the dimensions of the surface cracks.

386
50 100 150 200 250

CRACK RADIUS (urn)

Fig. 6. The relation between indentation load and surface crack radius.

Cracks were introduced into each sample using indentation loads of 10, 20, and 40 N cor- responding (fig. 6) to crack radii, a, of 60, 80, and 120 ym, respectively. The crack planes

in each case were normal to the long axis of the sample; this coincides, in turn, with the

incident direction of the acoustic surface wave. For the polished sample, the background

scattering was minimal, and all cracks could be clearly distinguished, as illustrated in

figure 8a, for the 60 ym crack.

For the sample with the 30 ym finish, the background level increased markedly (fig. 8b),

and although distinct signals above the background were obtained for the 80 ym (fig. 8b) and

120 ym cracks, the 60 ym crack could not be discerned from the background scattering. Additional signals that could not be related to the deliberately introduced cracks were also

apparent (fig. 8b). It is recalled that this sample contained unusually deep surface troughs

(fig. 7b), which may signify the presence of relatively large underlying cracks introduced

during a prior stage of the surface preparation sequence. The additional signals are thus

almost certainly related to the presence of large cracks other than those introduced by

indentation. This observation thus offers encouragement that the large size extreme of

naturally occuring cracks will be distinguished by the present acoustic method.

For the sample with the 70 ym finish the even larger background level prevented the

detection of both the 60 and 80 ym cracks, although a clear signal was obtained from the

120 ym crack (fig. 8c). On the rough ground sample, none of the deliberately introduced

cracks could be isolated. However, distinct signals were apparent in other areas of the

sample.

To ascertain the origin of the background (i.e., surface roughness and/or surface

microcracks) each sample was polished to remove the surface irregularities, but to retain

the sub-surface microcrack distribution. The acoustic measurements were repeated. No significant change in the signal amplitude or distribution could be detected. The backscattering

is thus attributed, almost exclusively, to the arrays of surface microcracks introduced

during the grinding process.

387
a) POLISHED SAMPLE

' 0.2mm 1

I'

SC5064.2TR

O.Olum

b) GROUND SAMPLE - 30u n DIAMOND

0. 1pm

GROUND SAMPLE - 70um DIAMOND

0.2mm 0. 5pm

d) ROUGH GROUND SAMPLE

1 0.2mm

1 I

0. 5um

Fig. 7. Surface profiles of the various surfaces used in the present study, obtained

using a Datatrak instrument.

388
END OF SAMPLE

END OF SAMPLE

Fig. 8. Oscilloscope records of signals obtained from surface cracks.

389
V.2 Measurements on Turbine Blades

A schematic of the turbine blade configuration (NC 132 silicon nitride) and the trans- tit ducer arrangement used for the study is shown in figure 9. The reduced section was afforded li.jict particular emphasis, since the highest stress obtains in this region during operation. A

LiNb03 transducer was used as both transmitter and receiver. Experiments were conducted at |ij

64 MHz or 100 MHz.

Fig. 9. Schematic of turbine blade and experimental set up.

The initial inspection of the turbine blade provided several indications from the

reduced section. However, this section (in the blade examined) contained a series of parallel machining grooves ^20 ym in wavelength and several microns on depth. To ascertain the

relative contribution to the scattering from the groove and the underlying cracks, the surface

in this region was finely polished to remove the macroscopic grooves. This procedure eliminated several signals, but the dominant ones, signifying the presence of the cracks, remained

(fig. 10).

Fig. 10. Reflections obtained from surface cracks in the next region of a turbine blade

390
Estimation of the magnitude of a crack in the reduced section of the blade differs from

the analogous problem for flat plates by virtue of the presence of the curvilinear propagation surface and the concomittent propagation loss. The most convenient method for obtaining

accurate crack size estimates is to calibrate the measurement system by obtaining the scattered

amplitude for a well-defined crack located in the reduced section. Such a crack was generated

by a 10 N Knoop indentation, followed by removal of the plastic zone. The signal amplitude

obtained was then used as a reference for predicting the size of surface cracks in other

blades with the same configuration.

VI. Discussion and Conclusion

The detectability of a crack is determined by the power scattered by the crack, relative

to the insertion and propagation losses, and the input power, and is strongly dependent on the frequency employed. The detectability levels in the presence of microcracks are not yet

theoretically predictable, but a theory analogous to that devised for grain scattering [21]

could be developed. The present observations represent an upper bound for the smallest

cracks that can be detected for each surface condition. There is an optimum frequency for

crack detectability that depends on the size of the crack with respect to the large extreme

of the background crack array; this frequency will be in the range 0.1 z ka z 1 (k = 2tt/x is

the wave number), a range in which the scattering amplitude from the crack of interest

approaches its maximum value (ka ~ 1) but varies less rapidly with frequency than the backscattering from the crack array (ka < 0.1). Subsequent studies will be concerned with the

evaluation of optimum frequencies for detectability and crack size characterization. The

transducer designs needed for this purpose have been established, as described in the present

paper. Also, it is already apparent that cracks in the 50 - 200 ym range will be detectable:

the specific level of detectability being governed by the characteristics of the background

micro-crack arrays.

In instances where individual cracks cannot be detected, the backscattered amplitude

may be an appropriate measure of the tendency for fracture [20]. The amplitude of the backscattering reflects the large extreme of the cracks that exist in that area element of surface,

in a manner comparable to the attenuation caused by grain scattering [20]. This possibility

is presently being explored.

Finally, it is emphasized that the method can be readily applied to components of complex

shape, as demonstrated by the tests conducted on turbine blades.

Acknowledgment

We would like to thank J. Shyne and M. T. Resch who carried out the mechanical tests on the glass samples and some of the comparison calculations.

This work was supported by Office of Naval Research Contract No. N0014-78-C-0283, NSF

Contract No. DMK-76-00726 through the Center for Materials Research at Stanford University,

and NSF Contract No. ENG-75-18681

References

[1] Evans, A. G. and Jones, R. L., J. Amer. Ceram. Soc . 61 , 156 (1978).

[2] Mecholsky, J. J., Freiman, S. W. and Rice, R. W. , J. Mater. Sci . H, 1310 (1976).

[3] Rice, R. W., Ceramics for High Performance Applications , J. J. Burke, ed. p. 287,

(A. E. Gorum, and R. N. Katz, Brook Hill, 1974).

[4] Mecholsky, J. J., Freiman, S. W. and R. W. Rice, J. Amer. Ceram. Soc . 60 , 116 (1977).

[5] Evans, A. G. , Kino, G. S., Khuri-Yakub, B. T. and Tittmann, B. R., Mater. Eva! ., April

1977, p. 85.

[6] Kino, G. S. , J. Appl . Phys . 49, 6 (1978).

[7] Auld, B. A., Wave Motion 1 , 1979.

391
[8] Born, M. and Wolf, E., Principles of Optics , (Pergamon Press, New York, 1959).

[9] Cuozzo, F., Cambiaggio, E. L., Damiano, J. and Rivier, E., IEEE Trans. Sonics Ultanson . SU-24 (1977).

[10] Domarkas, V., Khuri-Yakub, B. T. and Kino, G. S., Appl . Phys. Lett . 33, 7 (1978).

[11] Auld, B. A., Acoustic Fields and Waves in Solids, vols. I and II, (Wiley-Interscience,

New York, 1973T

[12] Fraser, J. D., Khuri-Yakub, B. T. and Kino, G. S., Appl. Phys. Lett . 32, 11 (1978).

[13] Khuri-Yakub, B. T. and Kino, G. S., Appl. Phys. Lett . 32, 9 (1978).

[14] M. T. Resch, B. T. Khuri-Yakub, G. S. Kino and J. C. Shune, "Stress Intensity Factor

Measurement of Surface Cracks", presented at 1st International Symposium for Materials

Characterization, 1978.

[15] Wachtman, Jr., J. B., Capps, W. and Mandel , J., J. of Mater . 7, 188-194 (1972).

[16] Evans, A. G. and Johnson, H., J. Mater. Sci . 10, 214-222 (1975).

[17] Shah, R. C. and Kobayashi, A. S., Proc. of 1971 Nat. Symp. on Frac. Mech ., Part 1 (ASTM

STP 513, 1972).

[18] Wiederhorn, S. M., J. Amer. Ceram. Soc . 52, 2 (1969).

[19] Budiansky, B. and Rice, J. R., J. Appl. Mech , 45, 2 (1978).

[20] Petrovic, J. J. and Mendiratta, M. G. , J. Amer. Ceram. Soc . 59, 163 (1976).

[21] Evans, A. G., Tittmann, B. R., Ahlberg, L., Khuri-Yakub, B. T. and Kino, G. S., J. Appl

Phys . 49, 2669 (1978).

Discussion

BENNETT

Does the ultrasonic technique for crack detection make it possible to

detect cracks which reseal themselves and thus cannot be detected readily

by other techniques?

EVANS

The surface acoustic wave technique certainly detects cracks which are

not visible on the surface (even using the scanning electron microscope).

This reflects the surface obscuration of the crack (effected by the plastic

deformation at the surface), rather than a crack resealing.

I will speculate that partially resealed cracks will be detectable using

ultrasonic waves, whenever the strain amplitude associated with the acoustic wave is large enough to overcome the 'resealing' crack surface force.

These reopening requirements are not yet well understood.

BOWLES

1. How critical is the setting of the transducer on the substrate to generate the required surface waves?

2. Does the coupling medium affect the intensity of scattering from local

flaws?

392
EVANS

1, With a LiNbC>5 wave guide the setting is critical (0=13°). However, if

the wave guide is made from the same material as the material being investigated, this difficulty can be eliminated.

2. The penetration of the coupling medium into surface cracks can reduce

the scattering amplitude. The effect has not been intensively explored;

but should not cause difficulties if either the couplant is made from a polymer or if the transducer is remotely located from the region to be inspected.

WU

Can you comment on the orientation effect of the cracks on the detectibility of this technique?

EVANS

We are now exploring orientation effects. I will speculate that, because of the relation between k and Sn, that the orientation dependence of

the scattering will be similar to the dependence of the fracture stress on

flaw orientation (see J. J. Petrovic, M. Mendiratta, J. Am. Ceram. Soc.)

393
I
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

CHARACTERIZATION OF SURFACE DAMAGE IN CERAMICS USING

SELECTED AREA ELECTRON CHANNELING

J. Lankford and D.L. Davidson

Southwest Research Institute

San Antonio, Texas 78284

The technique of selected area electron channeling has been applied to

the problem of characterizing subsurface plastic damage introduced during

ceramic machining and pojishing operations. The surfaces of ceramics spanning a wide range in hardness were prepared under controlled grinding and

polishing conditions. Subsurface regions of sectioned specimens were subjected to electron channeling analysis, enabling the boundaries of the

subsurface damage (plastic) zones to be mapped out. It was determined that

under conditions of constant applied load, the depth of damage during

polishing is inversely proportional to the square root of the hardness;

this finding is in good agreement with predictions based on indentation

theory.

Key Words: Electron channeling; plasticity; subsurface damage; surface

damage.

1. Introduction

Damage produced in ceramic materials through conventional surface preparation techniques can manifest itself in at least two ways. First, a population of microflaws

capable of degrading the mechanical properties may be introduced; Becher, for example,

has shown [l] J that twins generated in alumina by surface grinding constitute failure

initiation sites upon subsequent bend testing. Second, it is known that microplasticity

attending mechanical surface preparation can be reflected in significant surface compressive stresses [2]. These likewise can exert an appreciable influence upon mechanical

properties.

A consistent problem in dealing with these damage factors is the inherent difficulty

in characterizing them. To a certain extent, topographically sensitive techniques such

as scanning electron microscopy and profilometry can be used to characterize features

such as surface cracks, pullouts, and surface roughness. However, the depth and severity

of near subsurface microplastic damage is much more difficult to ascertain. In a few

materials such as MgO, dislocation etch-pitting can be used to accurately map the subsurface plastic zone. However, the technique is not applicable in general, especially for

materials of engineering interest such as AI2O3, SiC, etc. In these latter cases the

only available data regarding the depth of machining damage has come from TEM studies

such as those of Hockey, et al . [3,4], which because of the difficulty in knowing precisely

the depth from which the foils are taken, provide rough estimates at best.

In this paper, a new technique is shown to be capable of providing hitherto unavailable information concerning near subsurface plastic damage layers in engineering ceramics.

The technique of selected area electron channeling has been used primarily in the study of

metal plasticity, but was recently applied by the authors to the study of indentation

plasticity in SiC[5] . It is possible that the technique will prove helpful in other

applications of ceramics where microplasticity is a consideration.

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

395
2. Choice of Materials

The principal criteria for the selection of materials were a wide range in hardness,

and that each material be available in single crystal form. The ceramics chosen, and

their measured hardnesses, are indicated in table 1

,

2 It was felt that statistical

scatter could be reduced by using single crystals, since varying polycrystal 1 ine grain

orientations would be expected to produce a non-uniform subsurface damage profile,

reflecting crystal! ographic anisotropy in hardness. All crystals were approximately the same size.

Table 1.

Materials Chosen and Hardness Values

Hardness (GNm~ 2

)

34.0

20,4

11,0

7.8

3. Experimental Procedure

For the surface finishing operations, two crystals of every specimen material were

each mounted separately upon stainless steel right circular cylinders whose faces had

been ground parallel. One set of four mounted specimens was then selected for diamond

surface grinding, while the other set was used for metal lographic surface polishing.

Grinding was carried out by gripping each cylinder in the chuck of a lathe, and making

a pass across the face of the exposed specimen with an 800 grit, diamond bonded cup

wheel. The three inch diameter wheel was turned at 3500 rpm, using a feed rate of

25pm/sec and a 25ym depth of cut; a steady stream of water soluble oil coolant was applied throughout the grinding pass.

For polishing, the other set of specimens was mounted in an automatic metal lographic

polishing apparatus. Wet polishing was performed using 600 grip SiC polishing compound

and a pressure of about five pounds. Polishing was continued until an equilibrium situation was achieved with respect to apparent (optical) surface finish, i,e,, with further

polishing, no change in surface texture was evident.

Following characterization of these surfaces in the scanning electron microscope

(using palladium to coat the non-conducting A1 2 0 3 ), it was necessary to section the

specimens to allow electron channeling characterization of subsurface damage. Since it was anticipated that the damage depth probably was not very deep, it appeared desirable

to section the specimens on a taper. This "magnifies" depth inversely as the sine of

the taper angle, thus increasing the damage zone dimension relative to the diameter of the

probing scanning electron beam. The taper sectioning was performed as shown in figure 1.

First, a cylindrical epoxy mount is machined so that one face (A-A 1 ) lies 15° from normal

to the axis. Next, a ceramic specimen is cemented to this inclined surface with the

ground/ pol i shed surface facing the mount. These are then placed in a mold, and epoxy is

added to reform the cylindrical mount, sealing the specimen within. The mount is

sectioned down to B-B' by sequential polishing, taking care to intersect the machined/

Material

SiC

A1 2 0 3

Si

GaAs

z M

g 0, with a hardness of -3.5 GNm" 2

, was also selected, but chipping of the taper edge

during preparation unfortunately precluded damage depth measurements.

396
polished surface of the ceramic, and finally the damage introduced into the ceramic along

B-B' during sectioning is removed by ion milling. This process was carried out at the

National Bureau of Standards by B. Hockey. Each specimen was argon-ion milled for a minimum of ten hours, so that flaw-free channeling patterns were produced in unground/

unpolished areas of the specimens.

Selected area electron channeling has been applied extensively by the authors in the

study of localized plasticity in metals [6] and ceramics [5], The technique, generally

described elsewhere [7], is as follows: The electron beam of an SEM is made to rock

about a small area on the specimen surface, and the electrons backscattered from that

surface are collected. Since electron backscattering is crystallographically specific, a pattern related to the orientation of the crystal is obtained. The acuity of lines within

the pattern depends on instrument parameters and on the defect state of the crystal. As

the defect density increases, the acuity of the lines on the channeling pattern is

decreased. The sensitivity of the line acuity to defect density is material specific, but

generally it has been found that very small amounts of crystallite damage may be detected.

Suitable theoretical and experimental techniques are not now available to make quantification of this effect generally possible. Thus, while it is not difficult to define the

extent of plasticity, a suitable technique has not been developed to quantify the magnitude

of the plasticity determined, except in general terms.

The depth within the crystal from which information is being derived by the channeling

technique depends on a combination of materials properties, accelerating voltage of the

electron beam and detector characteristics. For the present work, an accelerating

voltage of 20 keV was used for all the specimens and instrument parameters yielding a pattern of four degrees were obtained from an approximately lOpm diameter area of the surface. The same instrument parameters were used for all materials.

Since A1 2 0 3 is an excellent insulator, it was necessary to vacuum deposit a very thin

layer of amorphous graphite on the surfaces of the specimens studied. By carefully con- trolling the coating thickness, it is possible to obtain channeling patterns from A1 2 0 3 which have the same line acuity as from conducting specimens.

The electron channeling data in the present circumstances were obtained from the

ion-milled taper-section specimens as sketched in figure 2, Starting near the machined

surface, a local region on the order of 10pm in diameter was interrogated by the electron

beam, yielding a channeling pattern. This process was repeated as the beam was moved to

other locations further from the surface, until a pattern of identical quality to that

found at an "infinite" (large) distance was obtained, thus establishing the depth of

plastic damage, dp.

397
Plastic Zone

Boundary

Plastic

Zone Boundary

Plastic Damage Zone

lOym

Ground/Polished Surface

Ground/Polished Surface

Plastic Damage Zone

Fig. 2. Scheme for determining depth of subsurface plastic zone; arrows and spots indicate locations at which channeling

patterns are taken; spot size ~10ym

4. Results

SEM views of the ground and polished surfaces are shown in figure 3, where hardness

decreases moving from top to bottom. It is clear that despite the fact that both grinding

and polishing were carried out under identical conditions for all four materials, there is no correlation between surface features and hardness. For example, ground Si C and Si

exhibit chipping along with considerable plastic scoring, while ground A1 2 0 and GaAs seem to be chipped and burnished. Polished alumina (fig. 3d) has an extremely smooth sur- face interspersed with microspalling, while harder Si C (fig, 3c) and softer GaAs (fig, 3h) are heavily scored by plastic gouges. In this case, surface appearance is obviously a poor parameter by which to rank ceramics with regard to abrasion resistance.

GROUND POLISHED

SiC

(a) (c)

Fig. 3. High magnification views of ground and polished surfaces (700X)

(Continued on next page.)

398

An example of subsurface damage characterization using electron channeling is

shown for GaAs in figure 4. Five microns beneath the surface, damage is sufficiently

extensive that the corresponding channeling pattern is nearly destroyed (fig. 4a).

Somewhat deeper (d = 12pm, fig, 4b), the pattern begins to emerge, and is quite

recognizable just inside the plastic zone (d = 32pm, fig. 4c). The difference between

patterns obtained from just inside, versus anywhere outside, the plastic zone can be

clearly discerned by comparing figures 5c and 5d,

(b) d = 12pm (d) d = »

Fig. 4. Typical channeling pattern results for polished GaAS.

Similar results were obtained for harder materials such as SiC, for which the extent

of plasticity is confined to a region nearer the surface. (Table 2 summarizes the damage

zone results for all four materials.) Figure 5a shows a distorted pattern 10pm below

the surface, and figure 5b a much clearer version of the same pattern obtained right at

the apparent plastic zone boundary. Close comparison between the latter photomicrograph

and that corresponding to d = 35pm (fig. 5c) show that a subtle but resolvable loss in

clarity and fineness of detail characterizes the d = 20pm pattern.

400
Table 2. Summary of damage zone results.

Material Plastic Zone Depth (ym)

Ground Pol ished

Sir 19.5

A1 2 0 3 32.0 21 .0

Si 32.0 30.4

GaAs 46.5 38.5

(a) d = 10ym (b) d = 20ym (plastic zone boundary)

(c) d = 35pm

Fig. 5. Typical channeling pattern results for polished Si

401
The effect of hardness upon the depth of equivalent damage is suggested in figure 6.

At d = 3ym, polished silicon exhibits the pattern of figure 6a, and at d = %m, that

of figure 6b. At approximately the latter depth, polished GaAs produces a pattern

(fig. 6c) lower in quality than that of Si at d = 3ym. The lower hardness of the GaAs

clearly exerts a major influence upon the damage gradient.

(c) GaAs, d = 8ym

Fig. 6. Comparison of effect of polishing damage upon channeling patterns

for Si and GaAs, showing the influence of hardness.

This fact is emphasized in figure 7, a plot of the apparent plastic zone boundary d versus hardness H. Chipping of the specimen disallowed the determination of dp for

ground SiC, so that the slope of the ground specimen results is questionable. However,

for the polished state, it seems reasonable to consider the potential significance of

the observed relationship between d

D (polished) and H. As it turns out,

d

n (polished) « -V (1)

which has interesting theoretical implications.

402
100

o Ground

• Polished

(urn)

dp _

10 j i i

i i i i I J I I I L

1 10 100

H(GPa)

Fig. 7. Depth of plastic zone versus hardness.

5. Discussion

The preceding results indicate, as expected, that polishing introduces less subsurface

damage, for approximately equivalent grit size, than does grinding. Moreover, the

functional dependence of the polishing damage depth is controlled principally by hardness

(under constant load conditions). This observation can be rationalized by appealing to

the results of recent studies of the physics and mechanics of the elastic-plastic indentation process.

It is by now well known [8] that a sharp indenter under a load P produces a plastic

impression of characteristic dimension a, which is related to hardness through

where a is a geometric constant. If the load is constant, then

Recently, Lawn and Evans suggested [9] on the basis of geometrical similitude considerations that the depth of the plastic zone dj beneath such an indent should be proportional

to a. This prediction was borne out for Si C by recent work by the authors [5]. From

equation (2), one would then expect

for indentations in materials of varying H under equal loads. This result also is

predicted independently by the recent elastic-plastic analysis of Perrott [10]. The

similarity between equations (1) and (4) is obvious.

From the preceding, it appears that the depth of the subsurface damage zone caused

by polishing can be predicted on the basis of indentation theory, whereby a dynamic

(sliding) indenter (i.e., a grit particle) behaves like a static indenter. This would

a <* 1 (3)

(4)

403
imply that if one could determine the depth of damage in a ceramic of a given hardness

subject to prescribed polishing conditions, the damage zone in another ceramic polished

under the same conditions could be predicted, based on its hardness alone. It is clear,

in addition, that electron channeling recommends itself as a technique for obtaining information on plastic damage in materials which do not lend themselves to conventional

damage assessment techniques (such as dislocation etchpitting)

.

The authors appreciate the support of the Office of Naval Research through Contract

No. N00014-75-C-0668 during the course of this work. We are also very grateful for the

kind help of Bernard Hockey at NBS, who performed the ion milling.

[1]

[2]

[3]

[4]

[5]

[6]

[7]

[8]

[9]

References

Becher, P. F., Fracture-strength anisotropy of sapphire, Jour.

[1-2], 59-61 (Jan-Feb. 1976).

Am. Cer. Soc. , 59.

Rice, R.W., Panel discussion (Proc. Symposium Science of Ceramic Machining and

Surface Finishing, National Bureau of Standards, Gaithersburg, Maryland, (Nov. 2-4,

1970), Book, The Science of Ceramic Machining and Surface Finishing , NBS Spec. Pub.

348, S. J. Schneider, Jr., and R. W. Rice, eds., pp 407-408 (National Bureau of

Standards, Gaithersburg, Maryland, 1972).

Hockey, B. J., and Lawn, B. R. , Electron microscopy of microcracking about indentations in aluminum oxide and silicon carbide, Jour. Mat. Sci., ]0_, [8] , 1275-1284

(August 1975).

Hockey, B. J. , Wiederhorn, S. M. , and Johnson, H. Erosion of brittle materials by

solid particle impact (Proc. 2nd, Int. Conf. Frac. Mech. Ceramics, Pennsylvania

State University, State College, Pennsylvania, July 27-29, 1977), Book, Fracture

Mechanics of Ceramics , Vol. 3, R. C. Bradt, D.P.H. Hasselman, and F. F. Lange,

eds., pp 379-401 (Plenum Press, New York, NY, 1978).

Lankford, J., and Davidson, D. L. , Indentation plasticity and microfracture in

silicon carbide, Jour. Mat. Sci. (in press).

Davidson, D. L. and Lankford, J., Plastic strain distribution at the tips of propagating fatigue cracks, Jour. Eng. Mat. and Tech., 98, [1] , 24-29 (Jan 1976).

Booker, G. R. , Scanning electron microscopy-applications (electron channeling effects)

in Modern Diffraction and Imaging Techniques in Materials Science , S. Amelinckx, R. Gevers, G. Remant, and J. van Landuyt, eds., pp 597-653 (North Holland, Amsterdam,

1970).

Lawn, B. R. , Jensen, T., and Arora, A., Brittleness as an indentation size effect,

Jour. Mat. Sci., ll_s [3] , 573-575 (March 1976).

Lawn, B. R. , and Evans', A. G., A model for crack initiation in elastic/plastic

indentation fields, Jour. Mat. Sci., 12, [11] , 2195-2199 (Nov 1977).

[10] Perrott, C. M., Elastic-plastic indentation: hardness and fracture, Wear 45_, [3]

293-309 (Dec. 1977).

Discussion

VAN GROENOU

1.

d

Your results suggest that "the" force is constant, otherwise

1 ^-Vu would be difficult to understand. Our results on scratching on MgO crystals show etch pits at rather large depth, d being proportional to JF^ (F

n

= load). P

404
2. Any plans for changing removal rates? Note that more particles become

involved and that "the" force could well increase less than Z.

LANKFORD

1. We agree with the comments concerning the load dependency. Since

d. is shown to be proportional to a , it follows from Equation (2") that

one would expect to see d

i a

^

n

/H ; our results were obtained for Fn ~

constant

2. We have no current plans for changing removal rates, although this

is an obvious next step. The last statement is true. However, for

materials polished under the same "increased force" conditions, one would

still find dp w/l/H , but with all of the d 's increased by -v/F^.

NOVAK

1. How is the 15° angle sample generated?

2. Have you studies effects of heat treatments on the extent of work damage?

LANKFORD

1. The 15° angle is discussed in the paper.

2. No, not in ceramics. In metals, it has a very significant effect, reducing the extent and severity of damage.

405

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

Y-DEFLECTION MODULATED SECONDARY ELECTRON IMAGES IN THE

EVALUATION OF CERAMIC SURFACE FINISH

L. P. Cook, E.N. Farabaugh, and C. D. Olson

National Bureau of Standards

Washington, D.C. 20234

Variations in secondary electron emission intensity as a function of

surface topography have been examined for single crystal silicon surfaces

of three different roughnesses. It is possible to calibrate absorbed

current, and under carefully controlled conditions, to calibrate secondary

emission for quantitative study of surface angular variation. Such data

can be statistically processed to give information on surface slopes

analogous to that obtained from prof ilometry. The agreement between

profilometer and preliminary SEM data appears to be good for the roughest

of the silicon surfaces. The utility of this as a general method for the

examination of ceramic surfaces is suggested for instances where specimen

charging can be properly eliminated.

Key Words: Ceramic; scanning electron microscope; secondary electron

emission; silicon; surface roughness; y-deflection modulation

I . Previous Work

In an earlier symposium Sundahl and Berrin [I] 1 noted the difficulties encountered in

interpreting the secondary electron signal produced by electron bombardment of machined

surfaces of oxides such as A1 2 03 , most of which are excellent electrical insulators in their

pure state. While application of a conductive coat a few hundred nm thick is sometimes a solution, uniform coating of irregular surfaces is difficult. Possibly for these reasons,

White, et aj_ [2] did not derive quantitative data describing surface roughness from their

image analysis system. It is the intent of this paper to explore the potential of

secondary emission for use in the quantitative evaluation of surface finish by studying

carefully prepared surfaces of silicon. Pure silicon has a resistance lower than most

oxides at room temperature (e.g., 2.3xl05 ohm cm for Si vs 1016 ohm cm for Si02 ) [3]. It

does not charge sufficiently in the SEM to alter the secondary emission properties.

II. Signal Generation in the SEM

Because of the low energies involved (<50 eV) , the bulk of the secondary electrons

leaving the surface of a specimen originate in the top few nm where the probability of

absorption is low [4] and the energy is sufficient to penetrate the surface potential

barrier - hence the use of secondary emission for high resolution images. Both n [ratio

of backscattered electron current to primary beam current] and 6 [ratio of secondary

electron current to primary current; or secondary electron yield] increase with increasing

inclination. Production of secondaries is in part related to backscattering in that a high

proportion of secondaries (up to 80%) are produced by exiting backscattered electrons [5] as opposed to those produced by entering primaries. Whereas the backscattered electrons leave

the surface with a strong directional trend [6] secondaries do not. Efficient collection of

secondaries is facilitated by application of positive bias to a collector in the vicinity

of the sample. The ease with which secondaries are liberated depends upon the chemistry of

the solids involved: insulators such as MgO and A1 2 03 generally show a higher value of 6

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

407
than metals, in spite of the increased backscattering in the latter [7]. Although a com- pletely satisfactory theory of secondary emission has not been developed, a cosine

relationship between 6 and angle of incidence has been proposed [8]. One expression used

for approximation of the secondary electron yield is [6,9]:

_J

cos e (f(o)e

•z/x

s) <i

5

dx

!•) dx

where: e is the.angle of incidence, defined with respect to the normal of the specimen

surface; f(o)e~ 2' x s is the probability that an electron produced at depth z will escape to

the surface if x

s is the mean free path of the secondary electron and f(o) is a ratio of

1 . dE

secondary current to primary current at the surface; (— ' is a quantity relating to the dx

number of secondaries produced in a given thickness dx if E is the energy of the incident

electron beam, and e is the excitation energy required to produce a secondary electron.

In practice, quantities such as e, f(o) and x

s may not be known. If accurate values of

emission as a function of surface orientation are required for a specific application the

researcher is forced to rely upon empirically determined curves.

III. Experimental Approach

Total collection of electrons emitted from a sample surface in the SEM is experimentally

difficult. On the other hand current absorbed by the specimen under conditions of constant

incident current is perhaps the most unambiguous indication of surface beam/specimen interaction:

^Absorbed ~ *Beam " ^Backscattered

+

^Secondary-'

*

If a combined coefficient of emission is defined as

6 ' = 6 + n,

then 6 ' as a function of angle of incidence can be derived from the absorbed current, since

I, 6' = 1 Absorbed

^eam

Such a curve for a finely polished silicon wafer with a 20 KV incident beam is shown in

figure 1. The resulting curve is quite smooth, and the increased sensitivity of the combined

yield 6' to angular variation at values of e > 60° is apparent. Also, note that <5

' > 1 at

high angles of inclination, in accordance with other results [4].

Figure 1 thus represents the basis for quantitative SEM study of surface roughness.

In practice, compositional effects, noise and other electronic factors, electron beam

spreading and surface contamination limit the utility of this approach. In the present

study, the build-up of hydrocarbons due to pump oil polymerization in the electron beam

which frequently occurs [10] has been eliminated by performing the experiments in a drypumped SEM. High purity single crystal Si cut normal to [111] was used, ensuring freedom

from local compositional effects.

408
i

1 1 1 r

I i i i I i I 0 20 40 60 80

0, degrees

Figure 1. Variation in 6' with angle of inclination relative to incident beam for finely

polished Si at 20 kV.

The absorbed current is the most reliable indicator of surface topography - yet the

use of the secondary electron signal has certain advantages. Photomultipl ier amplification

of the secondary electron signal is highly efficient and rather straightforward when

compared with the amplification procedures required for absorbed current. Furthermore,

the secondary electron yield, 6, is more sensitive to variations in e than is the backscattered yield, n. Therefore the absorbed current, which combines these effects, is as a whole a less sensitive indicator than secondary emission current. Finally, the secondary

electron signal bears a more direct relation to the appearance of the sample surface, and

secondary electron images are familiar to most researchers. Absorbed current signals must

of course be electronically inverted to produce a similar effect. Most instruments are designed for easy manipulation of the secondary electron signal, but use of absorbed current

is in general not as convenient.

The principle difficulties in the use of the secondary electron signal for quantitative

study of surface roughness are geometrical in origin and have to do with signal collection

factors. To study the effect of variation in surface position relative to the collector

axis in our instrument, measurements of total signal were made while rotating the sample

at 5° intervals away from the collector axis and keeping the angle of inclination relative

to the electron beam constant. This experiment showed that a variation of a few tens of

degrees in surface orientation relative to the collector axis produced only small changes

in the collector signal which were well within the limits of error (1 to 5%) required for

this preliminary study. We recognize that for other studies and other instruments, this

condition may not hold. A second factor of equal importance is the differential contribution

of backscattered electrons, via 1 ine-of-sight shadowing, to the signal. In our study, this

effect has been minimized by use of high angles of incidence. Monte Carlo studies, such as those of Murata [6], have indicated a strongly peaked distribution of exiting backscattered

electrons. At high angles of incidence, the great majority of backscattered electrons

follow a trajectory close to the surface. Because of this, for a surface of high average

inclination, relatively few backscattered electrons enter the secondary collector, and

1 ine-of-sight shadowing effects are minimized (refer to fig. 2 for a schematic of the

geometrical arrangement in our system). The generation of additional secondaries as these

409
backscattered electrons impinge upon the sample chamber is an effect which can in large

part be compensated by calibration.

Incident

elect ron beam

Scintillator

at +12 kV

AC & DC

Amplification

Current Photomultiplier

Figure 2. Signal collection process in the SEM (modified after ref. [11])

With these restrictions in mind, we have studied variations in secondary emission for

silicon surfaces. Operating conditions were kept constant by utilizing constant AC and

DC gain (fig. 2), constant accelerating potential (calibrated with the continuum), and

constant beam current (measured in a 200 ym aperture Faraday cup). For comparative work

secondary collection geometry was optimized by use of e = 70°, and a 12 mm working distance

such that the scintillator (with collector at +300V) was positioned horizontally facing

the area of observation. By utilizing such operating conditions, the secondary emission

calibration micrograph in figure 3 was made. The 35° line scan differs only slightly in

position from the zero degree scan (refer to fig. 1), and gives an indication of electronic

noise, presumed nearly constant regardless of angle. The additional AC superimposed upon

the DC level at higher angles in figure 3 is related to surface irregularities on the

polished silicon wafer used for calibration. The electronic noise leads to estimated

angular detection limits as follows for various angles of inclination: ±10° at 35°, ±1.5° at

70° and ±0.5° at 80° under typical operating conditions for our instrument. Comparison of

figures 1 and 3 thus affirms, both at the signal generation and the signal collection steps

in the signal processing train, the strong dependence of 6 ' on e for e > 60-70°. Because

of play in the SEM stage, angles were always approached from the same direction (read to

within ±0.1° or better). Calibration from the image at 90° was relatively easy; however

additional calibration points were not available, and it is possible that a slight

systematic error at lower angles was present.

Figure 3. Calibration micrograph showing photomultiplier signal for scans across finely

polished Si at four different values of 6. Arrows indicate estimated electronic noise.

For quantitative or semi-quantitative comparison (fig. 4), single secondary emission

line scans are useful: these may be regarded as but one element of a y-deflection modulated

image. Contrast (A.C. amplification) may then be adjusted for maximum response (for rapid

410
survey of a smooth surface for presence of irregularities, a complete y-deflection image

is highly useful, as will be discussed in the ensuing section). Prof ilometry, for comparative purposes (fig. 4) was done using a diamond-tipped stylus of ^3 ym radius. The

instrumental set-up and data processing techniques used for profilometry have been described

by Teague [12].

500

Figure 4. Comparison of secondary electron emission (upper) and profilometer scans (lower)

for A) rough cut Si, B) chemically polished Si, C) finely polished Si.

In the present study, three specimens of single crystal Si cut normal to [111] with

the following surface preparations were used: rough-cut [R- = 0.57 ym) , chemically

polished (Ra

= 0.17 yra) finely polished (Ra

= 0.04 ym). Windowless energy dispersive

x-ray spectrometry verified the relative cleanliness of the surfaces in that carbon and

oxygen peaks were not detected (normally any very thin adsorbed layers would be vaporized

411
by local electron beam heating).

IV. Results

A. General Correlations

Results of comparative SEM and profilometer scans ^1.7 mm in length are shown in

figure 4. The SEM profiles indicate surface inclination vs position on the surface. A

detailed comparison of the two types of profiles is not possible because of the impossibility of positioning a stylus tip a few ym in diameter and an electron beam a few hundred

nm in diameter in precisely the same fashion to traverse a surface of fine detail. The

fluctuations in amplitude of the signals show a rather similar decrease in intensity from

rough cut to finely polished. Coarse periodicities appear to correlate somewhat for both

secondary emission and profilometer. The grooves of -v-1.5 pm depth, visible on the profilometer scan of the rough cut Si (fig. 4a) are associated with cutting striations; the fine

structure produced by the cutting action also varies in a periodic fashion, apparent in

the coarse undulations of the secondary emission signal. Undulations are visible as well

in both profilometer and secondary emission profiles of the chemically etched surface

(fig. 4b). The finely polished surface (fig. 4c) is essentially "flat" with respect to

both profilometer and secondary emission scans with the exception of a few undulations in

the latter which may indicate local variations of a fraction of a degree in the surface

uniformity.

B. Quantitative Comparison of SEM and Profilometer Data for Rough Cut Si

It is possible to gather quantitative angular information from secondary emission line

scans analogous to that obtainable from profilometry . This has been done for the rough cut

Si with the aid of a template constructed using a calibration scan (see fig. 3). Results

are given in figure 5, a histogram which is a probability density distribution function in

e. It was obtained by integrating manually the portion of the sampled surface lying within

each selected angular interval. The angle of surface inclination, e, is different from the

analogous parameter that would be measured by profilometry. In profilometry only the

component of slope in the direction of travel is measured. In the SEM, e is measured with

respect to the direction of travel of the electron beam. An additional degree of freedom

is thus present along a topographic traverse measured by the SEM, and the SEM-derived values

should be thought of in terms of the relevant spherical coordinates. Since the average

inclination of the surface with respect to the beam is 70°, the set of measured angles

should be centered approximately around 70°. The SEM-derived average angle of inclination,

e, for the rough Si surface is indeed 70.6°, close to the specimen stage goniometer reading

of 70.0°. The SEM-derived average slope, e^Q» is 8.4°, in reasonable agreement with the

profilometer average slope value of 9.7° (refer to Table 1 for description of the calculations) .

C. Y-Deflection Modulated Images of Finely Polished Si

As figure 4c indicates, the finely polished surface shows few definable irregularities.

It is in such cases that complete y-deflection modulated images can be employed to good

advantage for examination of large areas of surface. The increased detail with which

imperfections are visible on the surface in the image of figure 6b relative to the

unmodulated image of figure 6a can be maximized by appropriate instrumental settings. Such

increased clarity results from use of y-deflection as opposed to grey level as the contrast

device. It is possible that some contribution from subsurface defects may be involved in

the linear irregularities of figure 6b; the extent to which this is likely should perhaps

be investigated further [13], However, the potential of the method is emphasized when

considering the sensitivity to very local variations of a fraction of a degree in surface

orientation.

L

1 f

Ra is defined as -i— J [ y | dx where y is the height of the surface with respect to the

L o mean line; L is the traverse length of the surface profile.

412
/o

6, deg

Figure 5. SEM-derived data for rough-cut Si (from 650 ym sample) showing the distribution

of surface angular irregularities.

Table 1. Surface angular roughness parameters.

Angular arithmetic mean, e

e = zf.Q./zf.

Angular mean deviation, e MD

6 MD

= Ef

i' 9

i " § l /zf

i

f. is a weighting factor calculated from the

histogram in figure 5.

V. Conclusions

Under appropriate operating conditions, quantitative information concerning surface

angular roughness can be obtained for silicon surfaces from secondary emission. Suitable

calibration can be obtained with minimal difficulty on most instruments. The sensitivity

of secondary emission to variations in surface orientation is better than 0.5° for Si under

typical operating conditions. Undoubtedly this could be improved by careful attention to

the electronics involved.

Full y-deflection modulated images are useful primarily in examining very smooth

surfaces for irregularities; this can be done quickly for large areas. For quantitative

comparisons, line scans are needed. These can easily be correlated with normal SEM images

and if desirable, in a general way with profilometer data. The use of absorbed current

413
Figure 6. A - Secondary electron image of

smooth Si surface with coarse scratch.

B - Y-deflection modulated image of area

in A. Note increased clarity with which

minor surface flaws at left and right are visible.

B

2 |im

rather than the secondary emission signal is advisable for improved accuracy- this is likely

to result in decreased sensitivity, however.

Before the approach described here can be applied extensively to ceramic surfaces, it

will be necessary to know if conductive coatings can be applied. If not, curves of total

yield must be determined for each material; with these data operating conditions can be set

such that 6' = 1, hopefully supressing the build-up of surface charge sufficiently that

quantitative studies of surface roughness can be made. For efficient data collection,

data processing equipment should be employed. Further work concerning accuracy and precision of this approach is desirable, especially studies elucidating the effect of

electron beam spreading, and the possibility of reducing noise.

VI. Acknowledgements

We are grateful to E. Muth, who prepared the polished Si wafers used for calibration,

and to F. Scire, who completed the profilometer scans. The manuscript has been improved

through the suggestions of B. Hockey, A. Feldman, and D. Newbury, and particularly those

of T. Vorburger.

414
VII. References

[1] Sundahl , R. C. and Berrin, L., Analysis and characterization of ceramic surfaces

for electronic applications, NBS Spec. Publ . 348, 271-292 (1972).

[2] White, E. W., McKinstry, H. A. and Diness, A., Quantitative surface finish

characterization by CESEMI, NBS Spec. Publ. 348, 309-316 (1972).

[3] Wolf, H. F., Silicon Semiconductor Data , pp. 1 and 602 (Pergamon, New York, 1969).

[4] Bruining, H., The depth at which secondary electrons are liberated, Physica 3_, 1046-1052 (1936).

[5] Robinson, V. N. E., The origins of the secondary electron signal in scanning

electron microscopy, J. Phys. D 7_, 2169-2173 (1974).

[6] Murata, K. , Spatial distribution of backscattered electrons in the scanning

electron microscope and electron microprobe, J. Appl . Phys. 45, 4110-4117 (1974).

[7] Dekker, A. J., in Solid State Physics , pp. 422-426 (Prentice-Hall, Englewood

Cliffs, New Jersey, 1957).

[8] Oatley, C. W., The Scanning Electron Microscope, Part 1. The Instrument , pp. 156-165 (The University Press, Cambridge, 1972)

[9] Ichinokawa, T., Iiyama, M., Onoguchi, A., and Kobayashi, T., Charging effect of

specimen in scanning electron microscopy, Jap. J. Appl. Phys. T3, 1272-1277 (1974).

[10] Conru, H. N., and Laberge, P. C, Oil contamination with the SEM operated in the

spot scanning mode, J. Phys. E 8, 136-138 (1975).

[11] Wells, 0. C, Scanning Electron Microscopy , p. 25 (McGraw-Hill, New York, 1974).

[12] Teague, E. C, Evaluation, revision and application of the NBS stylus/computer

system for the measurement of surface roughness, NBS Tech. Note 902 , 142 pp. (1976).

[13] Baines, M., Clarke, D. R., and Howie, A., Backscattered images of dislocations and

crystal defects, in Scanning Electron Microscopy: Systems and Applications 1973 , pp. 443-335 (London, Institute of Physics, Conf. Series No. 18, 1973).

415

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

LATENT SCRATCHES ON OPTICAL GLASSES

THAT APPEAR AFTER CLEANING

T. Izumitani, E. Miyade, S. Adachi

Hoya Corporation

572, Miyazawa-cho, Akishima-shi, Tokyo, Japan

S . Harada

Hoya Electronics Co., Ltd.

3280, Nakamaru, Nagasaka-cho, Kitakoma-gun

Yamanashi-ken, Japan

The latent scratch which appears on polished optical glass surfaces was

studied. The latent scratch, as the word implies, can be hardly detected in polishing process, but becomes evident after cleaning or after an tiref lection coating.

The appearance of the latent scratch is affected by (1) the polisher, (2) the polishing grains, (3) the mechanical condition of the polishing, and (4) the glass property. And it is also affected by the corrosion power of the cleaning solution. It

is concluded that the latent scratches are the scratches which are formed as un- detectable fine scratches in the polishing process and grow to the detectable size

in cleaning process. On the basis of this fact, a method of preventing latent

scratches was suggested.

Key Word: Cleaning, chemical durability; glass surface; latent scratch;

optical glass; polishing.

1. Introduction

Recently the latent scratch has become a serious problem to the lens makers who coat

multilayer films on the lens surface, because antireflection multilayer coating makes it

easier to recognize the presence of the abnormality (stain, contamination and scratch) on

the glass surface. In addition the severe cleaning condition, which is necessary to get a completely clean surface, would accelerate the growth of the scratch. The latent scratch

was experienced by the authors when the staining of optical glass in the polishing process

was investigated. The latent scratch on the polished glass surface of BaCD 16 (Si02-B

2 03RO) was recognized during electron-microscopic observation.

Photo 1(a) is the electron-micrograph of the polished glass surface of BaCD 16 before

the cleaning, which shows the several spot-like Shiroyake (clouding) . When at first BaCD 16 was immersed in trichloroethylene , the surface hardly changed as shown in Photo 1(b). Then

this sample was immersed in water with ultrasonic agitation; the Shiroyake disappeared from

the surface, as shown in Photo 1(c). Photo 1(d) shows the surface immersed in detergent

solution with ultrasonic agitation. In this photograph, several scratches can be observed.

These scratches were not detectable before the sample was immersed in detergent solution.

It was necessary to make clear whether these scratches are formed in the cleaning process or

in the polishing process.

The purpose of this paper is to make clear the nature of latent scratches and how they

appear by studying the relationships: (1) between the polishing and the latent scratches,

(2) between the cleaning and the latent scratches, and (3) between the glass properties and

the latent scratches. We will also analyze the cleaning process and discuss detergents and

the relationships between the detergent and the latent scratches. Then we will suggest how

to prevent the formation of the latent scratches.

417
(a) (b)

(c) (d)

Photo 1. Appearance of latent scratch (BaCD 16)

(a) After polishing

(b) Cleaned in trichloroethylene with ultrasonic agitation

(c) Cleaned in water with ultrasonic agitation

(d) Cleaned in detergent with ultrasonic agitation

2. Latent scratch and polishing process

In order to study whether the latent scratch formation is correlated to the polishing

process, the relationships between (1) the latent scratch and the polisher, (2) the latent

scratch and the polishing grains (cerium oxide) , and (3) the latent scratch and mechanical

polishing condition were examined.

In this experiment, BaCED 4(Si02 -B2 03 -R0) was used. The cleaning process is shown in

Fig. (1). As the cleaning solution, (a) water, (b) detergent solution, and (3) 0.01N

nitric acid were used, respectively.

Trichloroethylene

^Trichloro^

ethylene

Trichloroethylene -5» Isopropyl

Alcohol —=» Water Cleaning —Solution^

Water Water Water —

Isopropyl —"isopropyl"1 -h Isopropyl —b — Freon

-5> ^Alcohol^ Alcohol ^ ^» Alcohol Vapor

with ultrasonic agitation

Fig. 1. Cleaning Process

418
2.1, Latent scratches and polisher

Photo 2(A) and Photo 2(B) show glass surfaces polished by Politex and Polyurethane,

respectively. On the surface polished by Politex, no scratches are seen after cleaning in

water or in the detergent solution. However, on the surface polished by Polyurethane,

many scratches are seen after cleaning in water or in detergent solution.

(a) Cleaned in water for 120 sec. with ultrasonic agitation

(b) Cleaned in detergent for 40 sec. with ultrasonic agitation

(c) Cleaned in acid for 40 sec. with ultrasonic agitation

(A) Polytex (B) Polyurethane

Photo 2. Effect of the polisher on latent scratch formation (BaCED 4)

Especially when the surface was cleaned in acid, the difference between the surface

polished by Politex and that polished by Polyurethane is remarkable. These results suggest

that the formation of the latent scratches is affected by the polisher.

Politex is the elastic polisher and Polyurethane is the viscoelastic polisher. This

may cause the considerable difference in the latent scratch formation even under the same polishing condition.

419
2.2. Latent scratch and polishing grains (cerium oxide)

The hardness of cerium oxide increases with the sintering temperature. Two kinds of

cerium oxide were obtained from cerium hydroxide by sintering at 1200°C and 900°C. The

former is harder than the latter. In this experiment, BSC 7 (I^C^-SiC^-l^O) was polished

with the two kinds of cerium oxide. Photo 3(a) and Photo 3(b) show the surface of BSC 7,

polished with softer cerium oxide and harder cerium oxide, respectively.

(a) Polished with soft grains (b) Polished with hard grains

Photo 3. Effect of the hardness of the polishing grains on latent scratch formation

(BSC 7 cleaned in detergent for 40 sec.)

(a) High polishing rate (b) Low polishing rate

Photo 4. Effect of the polishing condition on latent scratch formation

(BSC cleaned in detergent for 40 sec.)

The severe latent scratches are seen on the surface of the BSC 7 polished with harder

cerium oxide. This indicates that harder polishing grains make the more latent scratches

than softer grains.

2.3. Latent scratches and mechanical polishing conditions

The mechanical condition of the polishing can be changed by (a) pressure, (b) revolution speed and (c) the type of the polishing machine. Photo 4(a) shows the surface of

BaCED 4 which was polished at higher pressure and higher revolution speed than that shown

in Photo 4(b) . The polishing rate was approximately four times faster. The result shows

that the harder mechanical condition of polishing makes more latent scratches on the sur- face of the glass.

These results suggest that the latent scratches appearing after cleaning are affected

by the polishing process. The latent scratches appear under the severe polishing conditions, i.e. the polishing the glass surface by the hard polisher with hard polishing grains

under the hard mechanical condition.

From the fact that the number and size of

ing condition, it is concluded that the origin

very fine scratch, which cannot be detected in

latent scratches are affected by the polishof the latent scratch is already made as a

the polishing process and that the fine

420
scratches will grow up to detectable size in a cleaning solution, such as water or es- pecially detergent solution.

3. Latent scratch and cleaning process

Figure 1 shows a common automatic cleaning system for camera lenses. Pitch and the

protective film on the polished glass surface is dissolved into trichloroethylene. Detergent is used for removing the polishing grains (cerium oxide, etc.) from the glass surface.

Freon vapor is used to dry the lens. The latent scratch does not appear in the organic

solvent. The latent scratch becomes visible in detergent solution.

3.1. Corrosion behaviour and cleaning power of cleaning solution

A synthetic detergent usually consists of surfactant and inorganic builder (table 1)

The corrosive behaviour of the surfactant and the inorganic builder were examined. The

cleaning power in removing polishing grains and finger marks adhering to the glass surface

were also tested. The results are listed in tables 2 and 3.

Table 1. Main components of synthetic detergent

Surface active agent

Inorganic

builder

Anion

ROS0

3

Na

R-^ ^-SQ3

Na

NaHC0

3

Na

2

C°3

Na

3

P0

4

Na

4

P

2

0

7

Na

5

P

3°10

Na

2

S0

4

Powier

Nonion

RO (Ch„CH„0) H

2 2 n

R^ 0 (CH

2

CH

2

0) n

H

Liquid

Corrosive power is measured as the percentage of weight loss after immersing the glass

in the surfactant and inorganic builder solution for one hour at 50°C with stirring. All

these data in Table 2 were normalized to the weight loss for water. Table 2 shows the sur- factant itself cannot attack the glass but that water, which dissolves the surfactant, attacks the glass. Table 2 also shows the detergent exhibits almost the same corrosive power

as the inorganic builder. This means that the corrosive power of the detergent is caused by

the inorganic builder. Therefore, if the inorganic builder were eliminated, the corrosive

power would be decreased and the latent scratch would hardly appear.

The cleaning powers of the surfactant, the inorganic builder and the detergent were

compared qualitatively by judging whether the polishing grains or the finger mark adhered on

the glass surface are removed by immersing in those solutions. As shown in Table 3, the

surfactant and the detergent exhibit cleaning power against the finger mark, but does not

have the cleaning capability against polishing grains. Only Na5P30^Q which has the highest

corrosion power among the inorganic builders can slightly remove the polishing grains. This

suggests that in order to remove the polishing grains from the glass surface, the corrosion

of the glass surface is inevitable. In other words, only high corrosive agents can remove

the polishing grains from the glass surface.

3.2. Latent scratch and cleaning

The cleaning is considered to remove the contaminant from the glass surface. This is

achieved by two methods, as illustrated in Fig. 2, (1) to dissolve the foreign matter or

(2) to remove the foreign particles.

421
Table 2 Corrosion of glass by solutions

Glass

Solution ^vDH ^"^^^^

BaCED

4

BaCD

16

LaCL

6

LaC

12

LaFL

3

Detergent A 9.6 4.6 4.5 4.1 6.6 3.3 anion

Detergent B 11.9 3.8 1.9 2.2 2.2 0.7 nonion

Surfactant A 6.5 1.9 1.7 1.7 2.1 1.2 nonion

Surfactant B 6.5 2.7 2.0 1.1 1.1 0.9 nonion

Surfactant C 6.5 1.6 2.9 anion

H

2

0

6 1 1 1 1 1

Na

5

P

3°10 8.7 4.1 7.7 16.1 36.8 7.6

NaHC0

3

8 3.6 0.8 1.9 4.3 1.4

Na

4

P

2

0

?

11 5.8 2.7

Na

2

C0

3

12 2.7 1.1 0.9 0.3 0.3

Table 3 Effectiveness of contaminant removal

\^\Gontan

Solution >sP*

__glass

ii-" BaCED 4 BaCD 16

nant Cerium

oxide

Finger

mark

Cerium

oxide

Finger

mark Time

Detergent A 9.6 + + 3-5 mir i 1.0 wt%

Detergent B 11.9 +• + 1.0 vol%

Surfactant A 6>.5 + +

H

Surfactant B 6.5 + +

ii

Surfactant C 6.5 + +

ii

H

2

0 6

Na

5

P

3°10 8.7 + + 1.0 wt%

NaHC0

3 9

ii

Na

4

P

2

0

7

11 ii

Na

2

C0

3

12 ii

422
(1) Dissolution of the foreign matter

(I) Organic contaminant

Organic

Solvent

(2) Removal of the foreign particles

(I) Organic contaminant

7

Surfactant

o-

(II; Inorganic contaminant

(a) Dissolution of the interface

Acid /-\

Alkali W

(j Detergent

(b) Physical rubbing

Q O

Fig. 2. Mechanism of Cleaning

It is an example of the former that the pitch is dissolved into trichloroethylene

while the mechanism of removing the adhering polishing grains from the glass surface is

rather the latter phenomenon. It is considered that the interface between the polishing

grains and the glass surface is attacked by the corrosive cleaning solution, and as a re- sult, the polishing grains are removed.

Because the detergent contains the inorganic builder which attacks the glass, it can slightly remove the adhering polishing grains. But at the same time, the inorganic

builder attacks the glass surface and makes the latent scratches on the glass surface.

After all, these results indicate that in order to remove the adhering polishing

grains from the glass surface, the formation of the latent scratch on the glass surface

is inevitable. This means the polishing grains on the glass surface cannot be removed

without the latent scratch formation. We cannot use the detergent to remove the polishing

grains and avoid the appearance of the latent scratches.

The surfactant can not remove the polishing grains. As shown in Fig. (1), the sur- factant can react with the organic contamination with its oleophilic group. But it can not react with the inorganic material such as the polishing grains, because it has no affinity for inorganic substance.

The conclusion is that the detergent can hardly remove the adhering polishing grains

of cerium oxide without the appearance of the latent scratches. The adhering polishing

grains on the glass surface should be removed by other methods, for example, by scrub off.

4. Latent scratch and the glass property

The mechanism of the appearance of the latent scratch was considered as follows.

The glass surface is scratched by the polishing grains in the polishing process. The

invisible, submicron scratches might be caused by some coarse polishing grains. These

scratches will grow up to the extent of the detectable size in the cleaning process.

423
(a) Cleaned in water for 120 sec. with ultrasonic agitation

(b) Cleaned in detergent for 40 sec. with ultrasonic agaitation

(c) Cleaned in acid for 40 sec. with ultrasonic agitation

(A) F2 (B) BaCD 16

Photo 5. Effect of glass property on latent scratch formation

Due to the mechanism, the relationship between the latent scratch and the glass

property will be predicted. The tendency of generating the latent, fine scratches will

be correlated to the glass hardness, and the tendency of the fine scratch to grow to a detectable size scratch and the glass properties were examined.

4.1. Latent scratch and glass hardness and chemical durability of glass

In this experiment, BaCED 4, F2 and BaCD 16 were polished. BaCED 4 and F2 have almost

the same chemical durability, but F2 is softer than BaCED 4. Photo 5(A) shows the glass sur face of F2. From the comparison of the surface of F2 (Photo 5A) with BaCED 4 (Photo 2B) , it

is evident that the more latent scratches can be seen on the surface of F2.

BaCED 4 and BaCD 16 have the same glass hardness, but BaCD 16 is less durable than

BaCED 4. Photo 5(B) shows the glass surface of BaCD 16. From the comparison of the sur- face of BaCD 16 (Photo 5B) with BaCED 4 (Photo 2B) , it is evident, too, that the more

424
latent scratches appear on the surface of less durable BaCD 16 and softer F2. This indicates that the latent scratch is correlated to the hardness and chemical durability of

glass. The hard and high durable glass is desirable from the viewpoint of preventing the

latent scratches.

5. Conclusion

The cause of the latent scratch is the scratches made in the polishing process. Relatively large, submicron scratches caused by the larger grains mixed in the polishing grains

will grow up to the detectable size in the cleaning process. The size of the latent

scratches detectable with unassisted eyes is considered about 3-4 and Q the latent

scratch is essentially different from the substantial scratches (about 80A) related to the

polishing mechanism which can be observed with an electron-microscope after etching. The

following three counterplans would be suggested for preventing the latent scratches.

1) Polishing condition

It is desirable that the glass should be polished with the soft polisher

with the soft polishing grains under a mild mechanical condition.

2) Cleaning condition

The adhering polishing grains cannot be removed with the usual detergent without the formation of the latent scratches. It is desirable that

the adhering cerium oxide be removed by physically rubbing the glass surface.

3) Improvement of the glass property

It is desirable that the hardness and the chemical durability of the

glass are improved.

Discussion

BURKE

1. Is the groove that is revealed by the corrosive action of the detergent

(the latent scratch) revealed prior to cleaning by observing a change in index of refraction of the glass at the site where the latent scratch would

appear? This would be expected to be the case if the hard particle causing the latent scratch were to have caused compaction of the glass and thus

decrease its corrosion resistance.

ADACHI

I think that the latent scratches are not recognized prior to cleaning by observing refractive index changes for two reasons:

1. Refractive index changes in the surface are complex. While polishing

grains will cause compaction of the glass surface to increase the refractive

index, leaching of soluble components of glass lowers the surface refractive

index. Generally, refractive index of polished layer of optical glasses

are lower than that of bulk glasses because of low chemical durability.

2. Because the size of the latent scratches is very fine (perhaps less than

lju) , we can hardly find them or measure the refractive index change at the

latent scratch.

VAN GROENQU

Is the corrosion due to the grooves or possible cracks?

ADACHI

Grooves

425

Session IV

Effects of Machining and Surface Finishing

on Ceramic Properties

Session Chairman

R. C. Bradt, Pennsylvania State University

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

MACHINING FLAWS AND THE STRENGTH GRAIN

SIZE BEHAVIOR OF CERAMICS

R. W. Rice

Naval Research Laboratory

Washington, D. C. 20375

A model for the strength-grain size dependence of ceramics

failing from machining flaws is presented based on the observation that the size of machining flaws (C) shows little or no

dependence on grain size. Two regimes of behavior are seen.

At finer grain sizes, little or no strength dependence on grain

size (G) exists because flaws are > G. For large G, the significant decrease of strength with increasing G is attributed to

a transition from polycrystalline to either lower single crystal

or grain boundary (i.e. bicrystal) fracture energies as the ~

constant (C) becomes significantly less than the G as G increases.

The transition between the two regimes is when C and G are similar. Extensive analysis of strength-grain size data, and more

limited directly correlated f ractographic data support the proposed model. Spontaneous cracking in non-cubic materials is

shown to support the model, but not be an alternate explanation

for the transitions seen in machined AI2O3 or BeO. This analysis

also further supports the observation that machining flaws do not

vary greatly with typical variations in machining parameters.

Key Words: Strength; grain size dependence of strength; flaw

sizes; brittle failure.

1. Introduction

Single phase as well as some multiphase ceramic polycrystalline bodies,

especially those of limited to zero porosity, show a definitive and characteristic dependence of their brittle tensile failure on grain size (G) Ll-4],

For most materials, such behavior occurs to temperatures of at least a few

hundred °C and for many more refractory bodies to temperatures of 1000°C or

more. Typically, two regimes of such G dependence of brittle fracture are

observed. At larger G, there is typically a significant increase in tensile

(or flexure) strength(s) as G decreases. However, as G decreases below some

intermediate level, e.g. below 10 to 50 /jm, much less or possibly no, increase

in S is observed with further decreasing G (fig. 1). The latter regime showing more limited or no increase in S with decreasing G typjcally gives a nonzero intercept on a Petch plot, i.e. a plot of S versus G~2 . In the past,

such non-zero intercepts have often been interpreted as indicating strength

control due to microplastic nucleation or growth of cracks [l-3]. However,

recent fractographic information shows that this is commonly not the case,

that flaws, commonly from machining, are typically the source of failure [2-

4]. In the past, it was often assumed that the G dependence of S was due to

the flaw size (C) being related to G, i.e. most commonly it was assumed that

C = G. However, recent f ractographic information shows that the size of

flaws introduced by machining do not vary much, if at all, with G [5]. Thus,

flaws are typically much smaller than large grains and typically much larger

than fine grains. This extensive f ractographic information thus invalidates

Figures in brackets indicate the literature references at the end of this

paper

429
GRAIN SIZE (G)

t

(J

tr

U\Tg

(J

z

LU

DC

I—

(/)

Figure 1. Schematic representation of proposed strength-grain size model.

Note that this is divided into two regions. Region I is for larger grain

sizes where strength significantly increases with decreasing grain size.

Region II is for finer grain sizes where strength shows significantly less,

and possibly no, increase with decreasing grain size. The transition between these two regions occurs over a grain size range where the grain and

flaw sizes are similar as indicated by the vertical wavy line. The two

sets of trends indicated by 1) and 2) are for samples of the same material

but having different surface finish, sizes, or testing conditions. There

may be an overlap of one or both regions for two such sample sets over

the full grain size extent, or partial overlap (as shown), or no overlap.

However, it is commonly expected that due to scatter and other factors

that there will commonly be at least some degree of overlap in the larger

grain size regime. Note also on the right hand side of the figure that

430
the strength ranges for single crystals of the same material in their

weaker orientations with the same set of finishing or testing parameters

are indicated for the two sets of polycrystalline samples by 1) and 2) are

shown respectively by X^l and X£2.

the previous explanation for the grain size dependence of brittle failure of

ceramics under tensile loading. The purpose of this paper is to review the

pertinent f ractographic observations on machining flaws and S-G data in order

to relate the two together and to discuss the mechanisms of failure that they

demonstrate or suggest.

2. Mechanisms of Failure

Before reviewing the f ractographic and S-G data and their interrelation,

it is appropriate to first briefly review the two basic mechanisms of failure. One mechanism is failure due to microplastic crack nucleation, due to

blocking of slip bands or twins by other slip bands or other twins, or more

commonly, by grain boundaries. Microplastic failure is typically represented by the equation:

S = O

q + KG~^ (1)

where cr is the stress to activate the slip or twinning phenomena that leads

to failure, and K is a constant. The second mechanism, that of flaw failure,

typically follows the relationship:

S = A (2)

where A is a geometrical factor associated with the shape of the flaw, and

may also be dependent on the size of the flaw relative to that of the specimen, as well as the location of the flaw in the specimen, E is Youngs modulus,

and y is the fracture energy. As noted earlier, it has commonly been assumed in the past that C was related to G, commonly C = G, implying that

one should be able to tell whether the mechanisms implied by eqs. (1) or (2)

were operative by plotting S versus G~2. Equation (1) would result in a

straight line with a non-zero intercept at o while eq. (2) would result in

a straight line with a zero intercept. However, as noted earlier and discussed in detail later, C is generally not related to G so this is not a

valid method of separating the two mechanisms. Further, there are other complications such as the fact that the two mechanisms may interact, e.g. slip

may aid the growth of pre-existing flaws to grow to a critical size. Also,

there can be competition between the two processes such that one may dominate in one grain size region and the other in another grain size regime,

with a transition between the two regimes shifting with such parameters as

impurities, porosity, and surface finish, e.g. machining.

3. Proposed Mechanism of Failure

Any mechanism of brittle failure of polycrystalline ceramics must address the following factors:

1. Foremost is the fact that failure typically appears to initiate

from machining flaws that have little or no dependence on G as noted earlier [5].

2. The observed dependence of strength on surface finish parameters

such as grit size [6-9] and machining direction [5,10].

3. The possible dependence of strength on specimen size [5],

4. That strengths of polycrystalline specimens can often extend below

431
the strengths of single crystals of the same material given the same surface i finishing which are tested in those orientations that give lower strength j e due to favorable orientation of lower fracture energy planes [2,4].

\ 5. The expected interaction or competition between microplastic and

brittle mechanisms of failure in materials such as MgO where the former mech- t anism can be operative Cl,2],

The proposed mechanism of tensile (or flexural) failure is as follows. j |j First consider the smaller G regime indicated by II in figure 1, i.e. where i strength shows either limited or no dependence of G. Fractography has clearly shown that in this region flaws are typically many grains in size [2-5]. | Since E does not depend on G, y often does not depend significantly on G [2, c 4,ll], and the flaws are multi-grained in size and show no relationship to

the grain size, no significant dependence of S on G would thus be expected.

This absence of any significant G dependence for failure with C > G gives a j n non-zero intercept on a Petch plot, e.g. figure 1. This is now believed to

be the predominant source of non-zero intercepts. The possible limited increase of strength with decreasing G in this region (II) can arise from

either of at least two sources. The first source is some possible limited

decrease of C with decreasing G [5], e.g. possibly following the limited

increase of hardness with decreasing G. A second possible source is the effect of internal stresses, i.e. due to incompatible strains between grains in

non-cubic materials or materials that have undergone a phase transformation.

Since there are statistical variations of these stresses, as the number of

grains encompassed by C decreases, internal stresses can increasingly aid

failure. Such effects can begin to occur even when flaws encompass many

grains Cl,2,4,l2]. With C nearly constant, the number of grains encompassed by the flaw increases with decreasing G, which would thus reduce the

internal stress effect, and hence increase S.

This leaves the significant dependence of S on G in region I and the

reason for the transition between regions I and II of figure 1 to be explained. A key observation is that the transition between these two different regimes of S-G behavior typically appears to occur when C ~ G [2-4, 13],

Two explanations for the significant decrease in strength as G increases so

the ~ constant C becomes progressively < G have been considered. The first

is that the machining flaws start propagating below the failure stress but

are arrested when they encounter the adjacent grains. Failure would thus

result when the stress increases to the level to propagate the flaw now the

size of the grain in which it has grown, or the size of one or more grain

boundary facets for intergranular flaws, into or around the surrounding

grains. This mechanism provides a ready explanation for the G dependence

of S in region I (fig. 1) since it gives the final flaw size ~ G. However,

this mechanism leaves unexplained how cracks can grow several fold in size

so their increase in stress intensity is greater than the change in fracture

toughness on going from a single- to multi-grain crack can be arrested.

Further, this mechanism does not explain how polycrystalline bodies with

initial flaws similar to those in single crystals of the same material can

have strengths similar to, or below, those of the weakest orientation of

such crystals when polycrystalline fracture energies are several fold times

those for common single- or bi-crystal fracture.

The second explanation for the significant G dependence of S in region I rests on the recognization that for C sufficiently smaller than G flaws would

effectively see themselves in a single crystal environment for flaws within

grains, or in a grain boundary environment for grain boundary flaws. In the

extreme of such cases, failure would then be controlled by the single crystal

fracture energy (yc ) for failure on the easier cleaving or fracture surfaces of the single crystal or the fracture energy (Yg) for crack propagation along a single grain boundary, e.g. bicrystal fracture energies. yc values vary from ~ 1/3 to ~ 1/10 of typical polycrystalline fracture energy

(ypc ) values. can be anywhere from a fraction, e.g. 1/2, to 100% of yc

. Thus, as previously suggested [2,4,14], it is proposed that the significant

432
increase in strength with decreasing grain size in region I arises from an

effective transition from yQ or Yq to ypc as G of different specimens

decrease relative to an ~ constant C.

Recent calculations indicate that unless the flaw is a sufficient fraction of G, e.g. 1/3 to 1/2, that propagation will become catastrophic before

the crack reaches the grain boundary where a major, if not the total, increase of fracture energy towards ypc would occur Cl4]. This transition has

been sketched out in some cubic materials based on calculations using the

initial flaw size [4,14], However, some crack growth may also occur El5].

There are also questions of whether the change from yc or Yq to ypc values

is really gradual or possibly a step function. The above noted transition

could, in fact, be a result of different more abrupt transitions of individual flaws in different specimens due to varying grain-flaw parameters

such as sizes and location [14], In any event, it appears that the second

mechanism, that of a y transition, applies when C is sufficiently small relative to G, e.g. C < G/3-, while the first is operative for C larger relative

to G, e.g. C > G/3.

The above transition in y is completed at different C/G ratios in different materials, e.g. ~ 1 to 3. This range depends upon a variety of

factors such as texture and the orientation dependence of y within a single

grain [14]. The C/G ratio over which this transition occurs can also vary

depending on the size, shape, and location of the flaw. Thus, for example,

a flaw may be rather shallow but fairly elongated so it crosses into two or

three grains, so it probably reached ypc sooner than a penny shaped flaw

of the same depth located entirely within one grain or entirely along one

grain boundary. Similarly, a flaw located entirely within a grain, but

located adjacent to one grain boundary may have a different transition than

a flaw whose ends are well away from the grain boundaries. Internal stresses

from thermal expansion anisotropy (TEA) or phase transformation and redistribution of applied stresses due to elastic anisotropy (EA) may also effect the C/G ratios for completing this transition.

Note that the polycrystalline strength in region I can extend below

the strengths for the weaker orientation of single crystals having the same

surface finish as the polycrystalline materials for several reasons (fig. 1).

There are extrinsic causes of this, such as impurities, and porosity. However, there are also intrinsic causes such as failure from grain boundary

flaws in materials in which values are sufficiently below those of yc values. Also, internal stresses from phase transformation or TEA can contribute to failure either by adding to the applied stress or by allowing

larger machining flaws to form, e.g. by machining flaws linking up with preexisting cracks from the internal stresses or cracks forming from such

stresses during machining operations. Stress concentrations due to EA,

which also occurs in cubic materials, may also contribute to lower polycrystalline versus single crystal strengths.

The above model is clearly consistent with all of the observed effects

of machining. Thus, for example, grinding perpendicular to the tensile axis

of bars generally introduces more elongated and hence more severe flaws

whether they are much larger or smaller than G [5]. They also lower the

transition because, as noted above, elongated flaws should reach ypc values

sooner because of their interaction with a greater number of grains. Similarly, effects of changing grit size would have the same effect. Again, in

a similar fashion, if larger flaws occur in larger bodies, they will lower

strengths in both large and fine G bodies and thus give transitions at

larger G. Finally, the competition and trade-off between microplastic and

flaw failure can also be anticipated. Since slip or twin nucleation of

cracks is typically associated with slip bands and/or twins propagating across a single grain, a pre-existing flaw larger than that grain will dominate failure. Thus, entirely brittle failure of materials that can exhibit microplastic failure, such as MgO and CaO should occur for C > G.

433
However, as the C approaches G, the microplastic mechanisms can become competitive so that as flaws become smaller than the grain size, microplastic

failure can begin to occur „ When flaws are still a sufficient fraction of

the G, slip assisted crack growth may be a dominant mechanism of failure.

Subsequent to the authors original discussion of this model, Bradt and

colleagues also discussed a similar model in which the fine grain size region was also attributed to C > G [6,]. They, however, attributed the transition and the subsequent larger G region where S depends more extensively

on G to a change from machining to "intrinsic" flaws of unspecified character which they apparently felt depended upon G for an explanation of the G

dependence of S. Several aspects of the model proposed by the present

author were also independently proposed by Rhodes et al [16] who similarly

attributed the transition in S-G behavior to C ~ G. Also, some aspects of

the model were suggested in an earlier review of devitrified glasses by

Emrich Cl7], i.e. a two branch S-G relation depending upon flaw size. Emrich attributes this model to Stookley Cl7], but the latter reference does

not show a definite base for the model, so it appears that it is really

more due to Emrich himself.

4. Comparison of the Strength Grain Size

Model with Available Data

Considerable S-G data exists in the literature. Much of it lacks definitive characterization of either the surface finishing conditions or the

strength test methods, and little fractography has been done. Much of the

earlier data does not necessarily show a clear transition for several reasons. First, often only limited G ranges were investigated. Second, much

of the data was not plotted in a fashion to clearly show such transitions,

e.g. log-log plots which tend to obscure such transitions. Finally, there

is also an important experimental difficulty in many materials; namely,

variable grain sizes. Thus, for example, bodies such as AI2O3 which often

readily exhibit exaggerated grain growth may have one large grain at or

near the surface in which one or more machining flaws may form. If a flaw

forms in or along such a large grain, that flaw is likely to be the source

of failure which will then reflect a large, if not total, influence of the

size of that large grain. In such cases, plotting S as a function of the

average G can be quite misleading. Some have, therefore, plotted strength

as a function of the largest G. While this author was one of the first to

recognize such possible effects of large grains just inside the machined

tensile surface, he has also clearly observed cases where large grains are

present, but machining flaws did not form in them (e.g. fig. 2) and hence

plotting their strength as a function of the large G is incorrect. This uncertainty of which G is controlling thus leads to substantial scatter in the

data and must be considered. Despite these problems, there is still a substantial amount of data which either individually or collectively supports

the proposed theory.

Figure 3 shows one of the more complete studies which, in fact, was

one of the factors precipitating the development of the model of section 3.

Note in figure 3 that the inserted fractographs of fracture initiating flaws

clearly show an increase of the strength as G decreases as discussed in

the model. Also, note since this is a cubic material in which internal

stresses from TEA, or other complications such as those due to phase transformation are not present, one can calculate fracture energies from the indicated flaw sizes. Such calculated fracture energies do, in fact, show a transition from approximately 2 J/m^ for C << G in good agreement with DCB

measurements of yc to ~ 7 J/m^ for C >> G in good agreement with DCB measurements of ypc .

The only other set of S-G data that has associated fractography with it

are the more limited studies of B4C (fig. 4). Evaluation of earlier B4C

data suggests a transition as shown in figure 4, and the more recent data of

434
Figure 2, Example of fracture not initiating from larger grains. This lower

magnification photo of B4C (the same specimen shown at higher magnification in the right hand portion of figure 4) shows large grains to the

left of the machining flaw origin (arrow). The largest grain (G) despite

almost reaching the tensile surface is clearly not the source of failure

showing that larger grains, while often sources of failure, are not always

such sources.

the author is reasonably consistent with this, possibly suggesting a transition at somewhat larger G„ Also, it shows a larger transition and lower

strengths for specimens ground perpendicular in comparison to those ground

parallel to the tensile axis. Because of possible differing compositions,

and more generally because internal stresses may alter the apparent fracture

energies calculated from the flaw sizes, such analysis is not attempted here.

Consider now AI2O3, one of the more extensively studied materials; i.e.

See figures 5-7. Also note table 1 where much of the machining and mechanical testing parameters for these and other studies are given. Figure 5

shows earlier AI2O3 data of various investigators each indicating to various

degrees, a S-G transition as expected. More recent flexural data of AI2O3

(fig. 5) again indicates various transitions, e.g. at larger G for specimens

ground perpendicular than for those ground parallel to the tensile axis.

Figure 7 represents various data from diametral compression tests of AI2O3

with indicated specimen compositions and sizes. While clearly significant

less certain because of the limited data, these results are again consistent

with a S-G transition. Note that all of the AI2O3 transitions are reasonably consistent, especially when one considers the uncertainty of the G values due to exaggerated grain growth. More specifically, they are also consistent with flaws observed in single and polycrystalline AI2O3, e.g. figure 9

Consider next, BeO, one of the more extensively studied ceramics after

AI2O3. Note that the extensive studies of carefully characterized material

by Fryxell and Chandler [27] clearly suggest a transition at G ~ 25 /Jin (fig. 7)

While this data is corrected to zero porosity (P), the individual sets of

samples of ~ constant porosity each show similar trends. Also, this corrected data agrees well with the data of Bentle and Kniefel [28] for hot pressed

BeO with P ~ O. Note that Greenspan's incomplete data [30] would suggest a possible transition at finer G, which is consistent with the fact that his

specimens were probably ground approximately parallel with the tensile axis

while those of Fryxell and Chandler, and Bentle and Kniefel, were circumfer435
TABLE 1

MACHINING AND FLEXURE TEST DATA

Flexure Test

Investigator

A . A1203

: Spriggs &

Vasilos

Passmore

et al

Rice

Evans &

Tappin

Tressler

et al

B„ BeO:

Fryxell &

Chandler

Veevers

Bentle &

Kniefel

Machining

G (30 ju in)

Cross

Sect

.

0.38x

0.63

Specimen (GM)

Length Type^

G (ER ~ 0.4mm) 0.38x

then an. 24hr 0.63

900°C

G(320 grit) ER 0.25x

G, ER 0.5

0.4 x 0.5

CG rods

G rods

Core drilled

rods

Greenspan G(25-35 rms)

O'Neill &

Livey

Hill et al

C . MgO

:

Evans &

Davidge

Bradt et

et al

Rice

D. SiC:

Gulden

Cappola &

Bradt

Prochazka

& Charles

G?

G?

Sawn

0.58(dia)

0.41(dia)

0.25(dia)

~0.7 x 0.7

~0.43 x 0.43

~0.43 x 0.43

0.35x

0.45

1.27x

1.27

D(C) G, J- t.A.

100,220,400,

or 600 grit

G(320 grit), then n „_ DS (600 grit SiC)

II T.A., ER U ' OJ4.45

4.45

1.5

2.5

G various grits,

(ER ~ 1/16", 0.63 x 0.63 5.1 ~ 1 . 6 mm)

8.88

2.54

2.54

> 5

2.54

2.54

2.6

3.8

3ju lapped

(E.F.?)

G?

G(220 grit,

0.025mm per

pass, 0.31 m/min

EC)

.Olx

0.1

0.64x

0.64

0.254x

0.254

< 0.63

5.1

2.54

4

4

3

3

3

4

4

3

3

3

3

4

4

Span(s) LR3

(cm) (mm/min)

3.80

1.27

3.80

1.27

1.27

2.0

3.05

7.61

2.54

1.90

0.79

1.90

5.08

2.0

2.0

2.4

1.2

3.05

1.27

or

1.9

0.43

.016

3.81

1.27

1.59

~ 0.15?

1.27

0.051

0.5

1.52

0.76

0.051

1-2 min

to failure

unspecified

0.051

0.51

1.27

0.051

0.051

0.051

(continued)

436
TABLE 1

Flexure Test

Specimen (GM)

I

Cross Span(s) LR3

Investigator Machining Sect

Length Type2 (cm) ( mm/min)

Cranmer D(C) G x T.A.

et al 220,320,400 or 0.3x „ „ 1Q n nt- n 600 grit 0.3 3 '° 3 19 °" 051

(0 0 005 mm pass)

All machining with diamond abrasive unless otherwise specified. G = ground,

D(C)G = down (climb) ground, J- or h T.A. = perpendicular or parallel to tensile

axis (i.e. bar length), ER = edges radius or rounding, EC = edges champfered, EF = edge fininishing; DC = dry sanded, an = annealed.

i'3 point or 4 point flexure.

t'LR = loading rate.

437
GRAIN SIZE, G (/im)

100 25 10 4 2

I//G (cm"l/2

)

Figure 3. Strength-grain size behavior for MgAloO^. The crossed hatch region shows the extensive strength-grain size data for near theoretical

density MgA^CXj. of good purity, as reported by Rice and McDonough Cl3],

Note the inserted fractographs showing flaws found at fracture origins

for the indicated levels of strength. Also note the fracture energies

calculated from the observed flaws show the transition from single to

polycrystalline values as discussed in the text. Note that the flaw for

the insert of intermediate strength consists of both a transgranular machining flaw portion, marked F, as well as a portion along the adjacent grain boundary, marked GB. Also note that the weaker orientation of

single crystal specimens of the same size and surface finish as for the

polycrystalline specimens give typical strengths of ~ 210 MPa (30,000 psi

The strength of polycrystalline specimens extends below this level.

438
GRAIN SIZE <pm)

400 100 25 10 6 4 3 2

5 10 15 20253035404550556065 70 75 G' % (cmr*

Figure 4. Strength-grain size data for hot pressed B

4C Note the crossed

hatch area representing analysis of an earlier survey of B4C data [l],

which gives a fair indication of S-G transition. More recent results of

the author are shown by individual data points as well as the two fractograph inserts. The right insert shows a distinct fracture origin from

a flaw much larger than the grain size in a fine grain body ground parallel

with whe tensile axis (see also fig. 2). The left insert shows a flaw

whose depth is comparable to the intermediate grain size, but the flaw

is fairly elongated perpendicular to the tensile axis. Internal stresses

due to the non-cubic structure of B4C as well as possible variations of

stoichiometry in the different bodies makes fracture energy calculations

from these flaw sizes and their interpretation in terms of the proposed

fracture energy transition uncertain.

439
GRAIN SIZE G, ( M m)

400 100 40 20 10 7 4 2

l/\ZG" (cm"l)

Figure 5. Earlier strength-grain size data for dense, machined AI2O3. Note

that while there are some differences i*n corrections for limited amounts

of porosity, e.g. between Carniglia's and Rice's plot of Spriggs et al

earlier data Cll, all show strength-grain size transitions in the range

of ~ 20 to < 100 jjm. Note also that although specimens of Passmore et

al Cl9] were annealed for 24 hours at 900°C after machining that both

the fabrication temperature 1400°C) as well as the temperature (~

1700°C) for crack healing in larger grain AI2O3 [20] indicate that this

annealing did not significantly effect the machining flaws.

440
GRAIN SIZE, G (jum)

400 100 50 25 10 7 4 2

0 10 20 30 40 _50, 60 70 80 90

1/\AG (cm 1/2

) Figure 6. More recent strength-grain size data for AI2O3. Limited data

for parallel grinding, though scattered, shows a clear transition as

does the limited data for perpendicular grinding, although this is more

ill-defined [21]. Note that these transitions are reasonably consistent

with the data of Evans and Tappin [22] for sintered alumina, especially

when strengths are corrected for the ~ 5% porosity. Also note that the

data for surface finishing of Tressler et al [7] with various grit sizes

is consistent with the other data and the range of transitions that they

indicate

441
GRAIN SIZE, G ipm)

400100 30 10

Q.

CO

DCI-<

Q

ou

45 —

I I I J I

o /

1 1 i 1

HU

z. —

/ m X

ob / + —

30 — gJ *

25 A HOT PRESSED Al 203 23°C

/ SYMBOL CdMPOSITION INVESTIGATOR ^PFPIMFNI Of LUllVlLl>J (rm) \Ul 11/ DIA. THICK

Al2 03 + No ADD., GAZZA ET AL. 1.91 0.51

15 * MgO, OR CoO

• Al 2 03 + 5 V/O Mo \ RANKIN ET AL. 0.64 0.51

10 O Al 2 03 ONLY

J

+ Al 2 03 STEELE ET AL. 0.91 0.76

5

I 1— I i i i

i i l I i 0 10 20 30 40 50 60 70

1/v^G (cm- ,/2

)

80 90

- 30

25

20

- 15

10

100 110 12C

Figure 7. Strength-grain size data for diametral compression testing of

AI2O3. While the data [23-25] is limited, it is clearly consistent

with a transition which would also for example be suggested by the

data of Gazza et al [23] alone. Note that the use of Mo additions

appears to have no significant affect on the strength other than

controlling grain size, e.g. as shown by similar flexural studies of

AI2O3 specimens with a range of Mo additions including those beyond

the 5% level, by McHugh et al [26], Note also that Rankin et al are

the only ones to report a head travel rate (0„27 mm/min) ; however,

Gazza (private communications) notes that they used the same head

travel speed as Rankin et al . Steele et al's tests were conducted

in a hand-operated press.

442
Figure 8. Example of machining flaws in AI2O3. A and B) Machining flaw

at the fracture origin of a sapphire laboratory test bar failing at

138 MPa (20,000 psi) . C and D) Flaw at the fracture origin of a circumferentially machined large hot pressed AI2O3 tensile specimen having a

gauge diameter of 14.4 mm, failure stress 275 MPa (40,000 psi). Note

elongated flaws (arrows) due to machining perpendicular to the tensile

axis in both samples.

entially machined. Also, Veavers' data [29], although low, in part due to

its greater P, would also suggest a transition in the range observed for data

of Fryxell and Chandler, and Bentle and Kniefel, were it corrected to P = 0.

Figure 10 presents more recent BeO data that suggests a transition at finer

grain sizes in comparison with those of Fryxell and Chandler, and Bentle and

Kniefel. This is consistent with the data of figure 10 having higher strength

expected from their probable grinding ~ parallel with the tensile axis.

An important question is whether the indicated C ~ G transitions for

AI2O3 and BeO are due to machining or other flaws. In view of the densities

and qualities of the bodies, dominance by processing flaws whose size is

similar to machining flaws is unlikely. However, intrinsic flaws resulting

from the TEA associated with their non-cubic crystal structures must also be

considered. Data of Charles [33] on Lucalox rods with as-fired surfaces

indicates a S-G transition at G ^ 70 /Km. This is generally beyond the range

of observed transitions for machined AI9O3, but is in the G range for initiation of spontaneous cracks from TEA [341. Similarly, Virkar and Gordon's

S-G data on P" - AI2O3 with as-annealed surfaces [35] shows a S-G transition

at G ^ 120 jLtm. Though this value of G is probably somewhat high because of

the tabular shape of the grains [4], the apparently similar thermal expansions and anisotropics of QL and P aluminas again indicates that AI2O3 S-G

transitions from intrinsic cracking are beyond those seen for the machined

samples. This is also indicated by S-G data of as-annealed TiC«2 exhibiting

spontaneous cracking and an associated S-G transition at G s 50 \m [36],

443
400 100 25 10

GRAIN SIZE, G (urn)

6 4 3 1.25

50

45 BeO

23 °C

T

12%

20%

7%

5%

30

40

& 35

25

Xlu TM

20

15

10

SINTERED AOX,

UNORIENTED BeO

SINTERED UOX

BeO + 0.5% MgO

ORIENTED

HOT PRESSED

BeO + 1 % MgO

P: 0.2 - 0.7%

SINTERED UOX

BeO + 0.2% MgO

P ~ 5%

25 „~

O

20 jE

FRYXELL AND CHANDLER

CIRCUMFERENTIAL GROUND

CORRECTED TO P = 0

BENTLE AND KNIEFEL CORE DRILLED

VEEVERS

CIRCUMFERENTIAL

GROUND

15

LU

ccXLU

10

HOT PRESSED BeO + % Al 2 03 AS SHOWN P: 0.3 - 0.7%

OTHER BeO, MOSTLY SINTERED

P: TYPICALLY 1 - 6%

1

GREENSPAN

.10 -20 .30 .40 . 50 .60

1/V^G (cm-'/!

)

.70 .80 .90 .100

Figure 9. Strength-grain size data for earlier studies of BeO [27-30]. Note

that the data of Fryxell and Chandler [27] is a compilation of their three

sets of specimens having approximate constant porosities of 3.6, 8.6, and

13.6% all of which have been corrected to zero porosity using the porosity

dependence determined by their study. Each of the individual sets of the

specimens uncorrected for porosity would show approximately the same grain

size transition. Note also that these corrected data agrees extremely

well with hot pressed BeO of near zero porosity by Bentle and Kneifel [28].

444
GRAIN SIZE, G iyim)

10 20 30 40 50

1A/G (cm" 172

)

Figure 10. More recent strength-grain size data for BeO [31,32]. Note the

distinct transition shown by the data of O'Neil and Livey L31] and that

this transition is at a finer grain size than that indicated by data of

Fryxell and Chandler, and Bentle and Kneifel. This is consistent with

the proposed model since O'Neil and Livey' s samples apparently were both

smaller and were probably ground parallel with the tensile axis while

those of all authors in figure 9, except probably Greenspan, were ground

perpendicular to the tensile axis.

445
since TiC>2 exhibits such cracking at smaller G than AI2O3 [44], BeO exhibits

spontaneous cracking at similar grain sizes as AI2O3 [34 J, and hence the S-G

transitions for machined BeO of this paper are below those expected for intrinsic cracks.

Thus, two factors should be noted. First the above S-G transitions

associated with cracks from TEA significantly reinforce the model and the

concept of S-G transitions at C ~ G. Second, the machined AI2O3 and BeO data

is consistent with S-G transitions at G ~ the size of machined flaws, not intrinsic cracks, e.g. from TEA. This latter fact is also consistent with observations on cubic materials, such as SiC and MgO below where such TEA and

associated cracking does not occur.

Next consider the compilation of SiC data shown in figure 11. Note that

both the small polished samples of CVD SiC suggests a S-G transition at small

er G than does the data of Cranmer et al [9], the latter indicating progressively larger G transitions for progressively coarser finishing. Correction

for Cranmer 's et als use of the largest G would lower the G transition values

but not below that for the CVD SiC. Both of these sets of SiC data are consistent with the proposed model.

Finally, consider the data for MgO (fig. 12), which can exhibit microplastically induced or assisted failure. Note that the limited data for

parallel ground specimens, at best, only faintly suggests transitions at

somewhat small, but not necessarily unreasonable, grain sizes. However, perpendicular grinding results show a more definitive transition at quite reasonable grain sizes. Also, note that this data is quite consistent with

earlier extensive studies [40] as well as more recent studies [42,8].

A summary of the tentative, as well as the fairly well established

grain sizes for transitions in S-G behavior, are summarized in table 2.

Table 2 also shows transitions suggested, or fairly well indicated, for

other specimens based on re-evaluation of an earlier survey of strength

grain size data [l]. Note that all of these are consistent with the typical

flaw sizes observed for machining laboratory size specimens; the type of

specimens tested in these studies.

5. Summary and Conclusions

A model has been presented for the grain size dependence of ceramic

strengths controlled by machining flaws whose size has typically little or

no relation to grain size. The model consists of two regimes. In the finer

grain regime, flaws are substantially larger than the grain size so there is

little or no decrease in strength with increasing grain size. Thus, strength

grain size curves are typically straight lines extrapolating to non-zero

intercepts on a Petch (S-G~2) plot.

In the larger grain regime, flaws are smaller than the grains and

strengths decrease substantially with increasing grain size. When flaws are

not too much smaller than G, e.g. C > G/3 subcritical growth to the grain

size probably occurs. At small C/G ratios, e.g. C/G < 1/3 the fracture energies undergo a transition from grain boundary (for intergranular flaws) or

single crystal (for transgranular flaws) to polycrystalline fracture energies

as the flaw size to grain size ratio increases (mainly due to changing grain

size). In any event, the transition between the finer and larger grain size

regimes occurs when the flaw size (C) and grain size (G) are similar. Besides the smearing out of the transition in strength behavior due to C growing to G, the transition may not necessarily occur at C = G for reasons such

as flaw shape and location, as well as preferred grain orientation and the

crystal orientation of fracture energy. Thus, an increase in flaw size

shifts both the fine and large grain portions of the curves downward and increases the grain size of the transition between the two curves.

446
140

120

a

CD

100

80

hiw 60

ccXLU 40

20

100 25 10

GRAIN SIZE, G ( Mm)

4

SiC T

/

/

/

/

/

"RANGE CVD DATA OF GULDEN AND PRICE

(SEE RI CE) SAMPLES POLISH ED

LARGE GRAIN DATA OF

PROCHAZKA AND CHARLES:

-AS REPORTED

CORRECTED

/

/

/

/

Am

/a

/

/

CRANMER,

ET AL.

a 600 GRIT

400 GRIT

* 320 GRIT

• 220 GRIT

RANGE OF CAPPOLA

AND BRADTS DATA

BOUNDS OF PREVIOUS SiC

DATA SURVEY (RICE)

J I I L

.10 .20 .30 .40 .50 .60

1/\/"G (cm- %

)

.70 .80 .90

80

60 §

x

s-z

LU

cc H 40 to

LUCCX

20

Figure 11. Strength-grain size data for SiC [1,4,6,9,37-39]. Note the consistency between the range of the previous data surveyed by Rice LU and

the more recent data of Coppola and Bradt [30]. Also note that these

two ranges are more consistent with the correction of Prochazka and

Charles [38] data as suggested by Rice et al [4], These would suggest a

transition for the small CVD SiC specimens of Gulden and Price from a few

to ~ 20 microns. Note also the studies of hot pressed SiC ground or

lapped perpendicular to the tensile axis with the various grit sizes L9J

would indicate transitions at larger grain sizes. However, it should also

be noted that Crammer et al plotted their data as a function of the largest grain size in the body. While larger grain sizes within a given body

often mean the data should be represented by a somewhat larger grain size

than the average, complete use of the largest grain size is probably an

overcompensation (e.g. see fig. 2) so the G values for the transitions

indicated by Crammer et al's data are probably somewhat high.

447
40010050 20 10

GRAIN SIZE, G (jim)

7 4 2 ~

i—i—i 1 rrr

HOT PRESSED /

\- AND ANNEALED /

/

300

250

200

150

100

50

0

AS HOT PRESSED

MgO

23 °C

BRADT ET AL, MgO GROUND

OR LAPPED 1 TO TENSILE AXIS

WITH GRITS SHOWN

X: EVANS AND

DAVIDGE, HOT

PRESSED AND

ANNEALED MgO,

AS-SAWN

MgO

HOT PRESSED AS

AND HOT

ANNEALED PRESSED

x

x SPECIMENS GROUND PARALLEL TO THE TENSILE AXIS o

SPECIMENS GROUND PERPENDICULAR TO THE TENSILE AXIS •

_L JL JL J_

45

40

35

30

25

20

15

10

5

0

0 10 20 30 40 50 60 70 80 90 100

1/^G (cm 1/2

)

Figure 12. Strength-grain size data for MgO [8,21,40-42]. The dashed lines

represent earlier extensive data of Rice [40]. Note that the more limited

data for grinding parallel with the tensile axis [21,41] suggests that

the transition for the as-hot pressed and hot -pressed, and annealed bodies

at grain sizes from ~ 3 to 15 /im. Note the earlier specimens of Rice LU

were tested with a 600 grit dry sanded surface for as-hot pressed samples

while most of the hot pressed and annealed samples were tested with asannealed surfaces. This more definitive indication of a transition in

strength-grain size behavior for both types of samples ground perpendicular to the tensile axis and both of these transitions are clearly at

larger grain sizes than those indicated for parallel ground samples. Note

also the good agreement of the data of Bradt et al for samples ground or

lapped perpendicular to the tensile axis with the various grit shown.

Also, the data of Evans and Davidge [42] for hot pressed and annealed

samples of MgO tested with as diamond-sawn surfaces, is also reasonably

consistent with the other data, i.e. falling below the curve for hot

pressed and annealed MgO tested with annealed or ground surfaces as expected.

448
TABLE 2

TETCH PLOT" BRANCHING

Material

AI2O3

AI2O3

AI2O3

AI2O3

AI2O3

A1203

BeO

BeO

BeO

BeO

MgO

MgAl204

TiB2

B4C

B4C

B4C

SiC

SiC

SiC

SiC

Processing

Hot Pressing

Hot Pressing

Hot Pressing

Sintering P ~ 5%

Hot Pressed

Press Forged

Sintered

Hot Pressed

Sintered (P:l-5%)

Hot Pressed

Hot Pressed

I Hot Pressed

and annealed

Sintered andj

Hot Pressed ( Hot Pressing

Hot Pressing

Hot Pressing

Hot Pressing

CVD

Hot Pressed

Hot Pressed

Hot Pressed

P:0-2

Machining

Various grinding

and sanding

II grinding

J- grinding

Grinding

Grinding (500 grit?)

Grain Size (^m)

of Branching

30-90 ~ 20

30-50

20-25 ~ 8

5

Diamond Ground 10

Circumferential 25

(i.e. -l) grinding

Grinding ( h?) 20

Grinding ( 11?) 10-25

Grinding ( n?) 15?

Grinding \\ 4??

Grinding -l 40

Grinding \\ 15?

Grinding -l 25

Grinding \\ 25 ± 10

Unspecified machining 5 20

Unspecified machining 15 20

Grinding -l **** 20

Grinding -l Sr. 25?

Polishing 5 20

Grinding 1 (400 or 600 grit) 35

Grinding 4- (320 grit) > 35?

Grinding x (220 grit) 50?

All data from flexure tests at room temperature except for two noted cases

of AI2O3.

1 Approximate % porosity P shown for some bodies. Typically hot pressed

bodies had ~ 1% or less porosity.

1

II and J- refer to direction of grinding relative to the tensile axis. ? indicates uncertainty in direction. Unspecified machining means samples were

machined from larger pieces by unspecified processes which would most commonly be grinding, probably at or near n to the bar axis.

Question marks indicate values indicated by limited data.

'From data for diametral compression at 23°C.

'From tests at - 196°C Cl].

FData corrected to zero porosity. Consideration of values at the different ~ constant levels of a few percent porosity would give similar results.

449
A survey of strength-grain size data shows good support for the model

and the transition grain sizes most commonly ~ 20-50 /urn, are typically consistant with commonly observed machining flaw sizes.

References

[ l] Rice, R. W., Strength/Grain-size Effects in Ceramics, Proc. British

Ceramic Soc. [20], 205-257 (1972).

[ 2] Rice, R„ W. , Microstructure Dependence of Mechanical Behavior of Ceramics, in Treatise on Materials Science and Technology, Vol. II, Prop- erties and Microstructure , R. K. MacCrone, ed. , 199-381 (Academic Press,

Inc., New York, N.Y., 1977).

C 3] Rice, R. W 0 , Fractographic Identification of Strength-Controlling Flaws

and Microstructure, Fracture Mechanics of Ceramics, Vol. 1 , R. C. Bradt,

D.P.Ho Hasselman, and F. F. Lange, eds., 323-345 (Plenum Pub. Corp.,

New York, N.Y. , 1974)

.

[ 4] Rice, R. W., Freiman, S. W., Pohanka, R. C, Mecholsky, Jr., J. J. and

Wu, C. Cm., Microstructural Dependence of Fracture Mechanics Parameters

in Ceramics, Fracture Mechanics of Ceramics, Vol. 4 , R. C. Bradt,

D.P.H. Hasselman, and F, F. Lange, eds, 849-876 (Plenum Pub. Corp.,

New York, N.Y., 1978).

[ 5] Rice, R. W. and Mecholsky, Jr., J. J., The Nature of Strength Controlling Machining Flaws in Ceramics, this conference.

C 6] Bradt, R. C. and Tressler, R. E., Microstructure/Surface Finish Control

of the Strength of Polycrystalline Ceramics, Proc. of 6th Int. Mat.

Symp. Ceramic Microstructures-76 , R. M. Fulrath and J. A. Pask, eds.,

785-795 (Westview Press, Boulder, Colorado, 1976).

C 7] Tressler, R. E., Langensiepen, R. A., and Bradt, R. C, Surface-Finish

Effects on Strength-vs-Grain Size Relations in Polycrystalline AI2O3,

J. Am. Ceram. Soc. [57] 5, 226-227 (1974).

[ 8] Bradt, R. C, Dulberg, J. L. , and Tressler, R. E., Surface Finish

Effects and the Strength-Grain Size Relationship in MgO, Acta Metallurgica [24], 529-534 (Pergamon Press, Great Britain, 1976).

[ 9] Cranmer, D. C, Tressler, R. E., and Bradt, R. C. , Surface Finish Effects and the Strength-Grain Size Relation in Sic, J. Am. Ceram. Soc.

[60] 5, 230-237.

[10] Rice, R. W., The Effect of Grinding Direction on the Strength of Ceramics, The. Science of Ceramic Machining and Surface Finishing , NBS

Special Publication 348, S. J. Schneider and R. W. Rice, eds., 365-

376 (Govt. Printing Office, Wash., D. C. 20402, 1972).

[ll] Rice, R. W. and Freiman, S. W., The Grain Size Dependence of Fracture

Energy in Ceramics, submitted for publication.

[12] Rice, R. W., Pohanka, R. C. and McDonough, W„ J., The Effect of Stresses

from Thermal Expansion Anisotropy, Phase Transformations and Second

Phases on the Strength of Ceramics, submitted for publication.

[13] Rice, R. W. and McDonough, W. J., Ambient Strength and Fracture Behavior of MgAl2C<4, in Mechanical Behavior of Materials , Proc. of 1971

Int. Conf. on Mechanical Behavior of Materials, Vol. 14, 422-431

(The Society of Materials Science, Japan, 1972).

450
[14] Rice, R. W. , Freiman, S. W. and Mecholsky, Jr., J„ J., The Dependence

of Strength Controlling Fracture Energy on the Flaw Size to Grain Size

Ratio, submitted for publication.

[15] Singh, J. P., Virkar, A. V., Shetty, D. K. and Gordon, R. W., On

Strength-Grain Size Relationships in Polycrystalline Ceramics, submitted to J. Am. Ceram. Soc

[16] Rhodes, W. H. , Berneburg, P„ L. , Cannon, R. M. and Steele, W. C.,

Microstructure Studies of Polycrystalline Refractory Oxides, Summary

Report, 29 March 1972 to 28 April 1973, prepared for U.S. Naval Air

Systems, Contract No. N00019-72-C-0298

[17] Emrich, B. R. , Technology of New Devitrified Ceramics - A Literature

Review, Technical Documentary Report No. ML-TDR-64-203 , Air Force Materials Laboratory Research and Technology Division, Air Force Systems

Command, Wright-Patterson Air Force Base, Ohio (1964).

[18] Stookey, S„ D., Ceramics Made by Nucleation of Glass-Comparison of

Microstructure and Properties with Sintered Ceramics, in Proc. Am . Ceram. Soc. Symp. on Nucleation and Crystallization in Glasses and Melts , 1-4 (1962).

[19] Passmore, E. M. , Spriggs, R. M. and Vasilos, T„ , Strength-Grain SizePorosity Relations in Alumina, J. Am. Ceram. Soc. [48] 1, 1-7 (1965).

[20] Lange, F. F. and Radford, K. C, Healing of Surface Cracks in Polycrystalline A1203

, J. Am. Ceram. Soc. [53] 7, 420-421 (1970).

[21] Rice, R. W., Machining of Ceramics, in Ceramics for High-Perf ormance

Applications , Proc. of Second Army Materials Technology Conference,

J. J. Burke, A. E. Gorum, and R. N. Katz, eds, 287-343 (Brook Hill

Pub. Co., Chestnut Hill, Mass., 1974).

[22] Evans, A. G. and Tappin, G., Effects of Microstructure on the Stress

to Propagate Inherent Flaws, Proc. British Ceramic Soc, No. 20, 275-

297 (1972).

[23] Gazza, G. E., Barfield, J. R. and Preas, D. L. , Reactive Hot Pressing

of Alumina with Additives, Am. Ceram. Soc. Bull. [48] 6, 605-610

(1969)

[24] Rankin, D. T. , Stiglich, J. J., Petrak, D. R. and Ruh, R., Hot-Pressing

and Mechanical Properties of AI2O3 with an Mo-Dispersed Phase, J. Am.

Ceram. Soc. [54] 6, 277-281 (1971).

[25] Steele, B. R. , Rigby, F. and Hesketh, M. C. , Investigations on the

Modulus of Rupture of Sintered Alumina Bodies, Proc. British Ceramic

Soc, No. 6, 83-94 (1966).

[26] McHugh, C. O. , Whalen, T. J. and Humenik, Jr., M. , DispersionStrengthened Aluminum Oxide, J. Am. Ceram. Soc. [49] 9, 486-491 (1966).

[27] Fryxell, R. E. and Chandler, B. A., Creep, Strength, Expansion, and

Elastic Moduli of Sintered BeO As a Function of Grain Size, Porosity,

and Grain Orientation, J. Am. Ceram. Soc. [47] 6, 283-291 (1964).

[28] Bentle, G. G. and Kniefel, R. N., Brittle and Plastic Behavior of HotPressed BeO, J. Am. Ceram. Soc. [48] 11, 570-577 (1965).

[29] Veevers, K., Recrystallization of Machined Surfaces of Beryllium

Oxide, J. Australian Ceram. Soc. [5] 1, 16-20 (1969).

451
[30] Greenspan, J., Development of a High-Strength Beryllia Material, Army

Materials and Mechanics Research Center Report No„ AMMRC-TR-72-16

(1972).

[31] O'Neill, J. S. and Livey, D„ T., Fabrication and Property Data for Two

Samples of Beryllium Oxide, United Kingdom Atomic Energy Authority

Research Group Report AERE-R4912 (1965).

[32] Hill, N. A., O'Neill, J. S. and Livey, D. T., The Properties of BeO

Containing up to 15 w/o SiC and up to 2 w/o MgO, Proc. British Ceramic

Soc, No. 7, 221-232 (1967).

[33] Charles, R. J., Static Fatigue: Delayed Fracture, in Studies of the

Brittle Behavior of Ceramic Materials , Technical Report No. ASD-TR-628,

370-404, Aeronautical Systems Division, Wright Patterson AFB, Ohio (1962

[34] Rice, R. W. and Pohanka, R. C, The Grain Size Dependence of Spontaneous Cracking in Ceramics, accepted for publication in J. Am. Ceram.

Sloe (1978).

[35] Virkar, A. V. and Gordon, R. S., Fracture Properties of Polycrystalline

Lithia-Stabilized £"-Alumina, J. Am. Ceram. Soc. [60] 1-2 , 58-61 (1977)

[36] Kirchner, H. P. and Gruver, R. M. , Strength-Anisotropy-Grain Size Relations in Ceramic Oxides, J. Am. Ceram. Soc. [53] 6, 233-236 (1970).

[37] Gulden, T. D. , Mechanical Properties of Polycrystalline 3-SiC, J. Am.

Ceram. Soc. [52] 11 , 585-590 (1969).

[38] Prochazka, S. and Charles, R. J., Strength of Boron-Doped Hot-Pressed

Silicon Carbide, Am. Ceram. Soc. Bull [52] 12, 885-891 (1973).

[39] Coppola, J. A. and Bradt, R.C., Measurement of Fracture Surface Energy

of SiC, J. Am. Ceram. Soc. [55] 9, 455-460 (1972).

[40] Rice, R. W. , Strength and Fracture of Hot-Pressed MgO, Proc. British

Ceramic Soc, No. 20, 329-363 (1972).

[41] Rice, R. W., The Effect of Grinding Direction on the Strength of Ceramics, in Proc. on the Science of Ceramics Machining and Surface Finishing, S. J. Schnieder, Jr. and R. W. Rice, eds., 365-376 (National

Bureau of Standards Special Publication 348 (1972).

[42] Evans, A. G. and Davidge, R. W., The Strength and Fracture of Fully

Dense Polycrystalline Magnesium Oxide, Phil. Mag. [20] 164 , 373-388

(1969).

Discussion

RHODES

How general are these results? If the crack size is not related to

the grain size, how do you explain the dependence of strength on grain

size?

RICE

The results should be quite general based both on the breadth of materials already considered as well as on the principles involved. The concepts appear to have no restriction to any particular material class. The

dependence of strength on grain size due to the initial flaw size not gen452
erally being related to the grain size is the subject of the paper "Machining Flaws and the strength Grain Size Behavior in Ceramics" by R. W. Rice

elsewhere in the conference.

TRESSLER

I think it is important to emphasize that in the systematic studies of

surface finish / grain size effects on strength of polycrystalline alumina

and silicon carbide conducted at Penn State that in the fine gram high

strength region the strength systematically increased with decreasing

grit size with very slight dependence on grain size. This result is very

important for machined high performance ceramics since one would like to

machine fine enough to have the grain size control the strength. For

fine grain ceramics one must use very fine polish to achieve this.

RICE

Although the effects of different grit sizes on the strengths of samples generally do not show up as being very pronounced on the scale of

strengths and grain sizes considered in this survey, it is none-the-less

agreed that grit size effects are quite important both in practice and

from the standpoint of scientific understanding of the strength behavior

of ceramics. Certainly the finishing studies at Penn State have been an

important contribution in demonstrating and documenting the extent and details of such surface finish effects. However, I question whether or not

one will commonly attempt to machine to a fine enough grit size to have

grain size control strength for two reasons. First, in many fine grain

bodies I expect that it will be fairly common that with finer finishing

that internal processing defects will take over control of strength rather

than the grain size itself. This is shown, for example, in the work of

Wu and McKinney on machining of Si3N4 elsewhere in this conference. It

is aiso indicated in your work on AI2O3 where the strengths of specimens

ground to 600 grit finish and those to a quarter micron diamond polish cannot be distinguished and in your work on SiC where the strengths of your

400 and 600 grit ground bodies cannot be distinguished. There will also

often be an economic constraint in limiting one to the degree of surface

finish that is conducted. Since the improvement in strength with finer

surface finishing commonly approaches a limit as indicated by the above

three cited cases, and this limit is often approached asymptotically because of increasing competition between internal processing defects and

machining flaws, one will have to consider the trade-off between decreasing improvement in strength vs continued increase in cost. However, I agree with what I believe is the general thrust of your comment, i.e.

that one needs to know what the rate of increase is as a function of surface finishing procedures in order to make this judgment, and that this can

have important practical results.

MALKIN

Can your results be explained alternatively in terms of the effect of

grain size on the relative amounts of transgranular and inter granular

fracture?

RICE

The results of different regimes of strength grain size behavior cannot

be explained in terms of grain size dependence on the amount of transgranu - lar -intergranular fracture. However, I am glad that you asked this question

since it is a fair one and I failed to address it in the text. The reason that I state that the mode of fracture is not a cause of the strength change is based on three observations and also is supported by some simple theo453
retical considerations. First, if one considers the overall trends in fracture mode of the whole fracture surface, one observes the trend for changing from one mode of failure to another to typically occur over a greater

grain size range than the transition in strength-grain size behavior shown

in this paper, see for example data on Mg0 aand AI2O3. Second, AI2O3

and MgO show opposite trends of the mode of fracture with grain size r i.e. MgO shows increasing transgranular failure with increasing grain size

while AI2P3 shows decreasing transgranular failure with increasing grain

size,(lj2y yet they both show similar strength grain size trends. Third,

MgAl204 and the B^C materials typically tend to show very nearly if not

total transgranular failure over the entire grain size range.

The question of the role of the mode of failure must, however, be

considered further since the overall mode of crack propagation may not have

a relationship to the mode of initial crack propagation. Thus, for example,

the author has observed cases of intergranular fracture immediately at the

fracture origin surrounded by 100% transgranular failure and more recently

my colleagues have shown that cracks can propagate in an intergranular mode subcritically due to environmental, e.g. moisture, effects and then change to transgranular propagation when failure becomes catastrophic in materials such as ZnSe and MgF2. However,

focusing in on the question of the mode of failure at fracture origin

in the materials of this study, the answer to your question still remains no.

MgAlgO^ is the only material of those studied here observed to indicate a change in the mode of fracture at the origin. MgAlgO^ fracture

origins of fine grain bodies are all transgranular while the fracture

origins of larger grain bodies were either inter-or transgranular or

mixed with greater occurrence of intergranular failure origin in contrast to

general transgranular fracture at all grain sizes. However, there is no

strength difference observed between specimens with a transgranular initiating flaws and those with intergranular initiating flaws.

One can also ask what effect the mode of failure might have on the different properties controlling strength. Fracture energy is the primary

candidate for this. However, in studies of MgO, fracture energy was constant over the broad range of grain sizes considered although there was a

change from essentially complete intergranular failure at fine grain sizes to

to essentially completely to transgranular failure at large grain sizes, (2,3)

Finally, some simple theoretical considerations support this lack of control

by whether the failure mode is intergranular or transgranular. Both estimates based on surface energy measurements of fracture energy of grain

boundaries show that typically they will range from ~ 50% to ~ 100% of

the lower single crystal fracture energies. However, simple geometrical

considerations show that the surface area created by intergranular failure

vs transgranular failure varies from a factor of ~ 2 down to only slightly

greater than 1. Thus, the differences in surface area between intergranular and transgranular failure approximately compensate for the potentially

lower fracture energy of intergranular failure suggesting that one should

not observe larger effects as previously discussed.

1) R. W. Rice, "Fracture Topography of Ceramics," Surfaces and Interfaces

of Glass and Ceramics , Plenum Publishing Corp., New York, New York, 1974

2) R. W. Rice, "Microstructure Dependence of Mechanical Behavior of Ceramics," Treatise on Materials Science and Technology , Vol. II,

Academic Press, New York, N.Y., 1977, pp. 199-381.

3) R. W. Rice, S. W. Freiman, & P. F. Becher , "Grain Size Dependence

of Fracture Energy in Ceramics" submitted to Am. Ceram. Soc. for publication.

454
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SURFACE MACHINING EFFECTS ON THE STRENGTH DISTRIBUTIONS OF ALUMINAS

T. E. Easier, T. A. Countermine

1

, R. E. Tressler and R. C. Bradt

Department of Materials Science and Engineering

The Pennsylvania State University

University Park, PA 16802

Three microstructural ly different, dense polycrystall ine aluminas were machined with 400-grit diamond at two different rates of surface removal.

Flexural strengths were measured in four-point bending under toluene and

the strength distributions analyzed by two-parameter Wei bull statistics.

Effects of edge-bevelling on the distributions were also examined. Machined

surfaces and the bevelled and non-bevelled sample edges were examined by

SEM. For all of the materials, the lesser rate of surface removal was less severe in edge chipping, and yielded higher average strengths and increased

values of the Wei bull parameters, m and a .

Key Words: Alumina; edge-bevelling; strength distributions; surface removal

rate; Weibull function.

1. Introduction

Machining of the surface of ceramic materials dictates their strength in cases where

machining induced surface flaws act as the failure origins. Surface flaws (features)

resulting after a given machining process are related to several machining variables, such

as the type of abrasive, the grit size of the abrasive, the direction of grinding, and

the rate of surface removal [1-4] 2

- Variations in the magnitudes of these effects from

one material to another result from differences in the physical properties and/or microstructures of the machined materials. This paper reports the results of a study of the

effects of the surface removal rate on the strength distributions of three polycrystal 1 ine

aluminas. Although more general strength distributions can be applied, Weibull statistics

were utilized because of their success in describing statistical distributions of the

strength of ceramics [5-7].

The three-parameter form of the Weibull distribution can be expressed as:

where F is the probability of failure, a

f

is the observed fracture stress, a is the lowest possible failure stress, o

Q

is a scaling parameter, m is the Weibull modiJlus, and S is the surface area of the specimen. If the lowest possible failure stress, a , is

assumed equal to zero, then the three-parameter form is reduced to the two-parameter form:

where S is the surface area in pure bending. In the four-point bend test when all of the

specimens fail between the inner knife edges, S is the surface area subjected to uniform

tensile stress.

Present mailing address: 307 Gilbreath Hall, East Tennessee State Univ., Johnson Ci

TN 37601

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

455
2. Experimental Procedure

2.1 . Test materials

Three dense sintered polycrystall ine aluminas were chosen for this study because they

represent a range of interesting microstructures. Their physical properties are listed

in table 1. The AL300 material has a rather large average grain size of 25ym, while the

AL600 has a much finer average grain size, about 10pm. The AL995 has a bimodal grain size

distribution consisting of large grains on the order of 25ym distributed in a finer grain

size matrix of about 12ym. It appears that the fine grain matrix of AL995 is comparable

to the microstructure of AL600. A semiquantitative spectrochemical analysis of these

materials, reported elsewhere, has shown the major impurities to be Si, Mg, Ca, and Fe [4k

Table 1. Physical Properties of Aluminas

Materials* P (gms/cc) K

I(

. (MN/m3/2) ^

E(MN/m2

)

AL300

AL600

AL995

3.74

3.71

3.85

2.53 ± 0.27

2.76 ± 0.10

2.48 ± 0.12

3.3 x 10'

3.2 x 10

5

3.6 x 10

5

Average

G.S.(ym)

25

10

12 .(25)

*Wesgo, Belmont, CA 94002.

2.2. Machining procedure

Specimens of the three alumina materials were finish machined at two different re- moval rates of down (climb) grinding [9], 2.5 x 10~4 and 2.5 x 10~5 inches/pass, to yield

six combinations of surface finish and microstructure.

Between one hundred and one hundred and ten specimens were preoared and tested at

each condition. Specimens were initially diamond sawed from 3/8 x 1-1/8 x 4-1/4" plates

to approximate dimensions and then mounted on a steel plate using glycol phthallate.

The surfaces were ground parallel, to size, with a 400-grit diamond wheel with the machining direction parallel to the subsequent direction of maximum tensile stress. The tensile

surface was then finish machined, removing 2 x 10~3 inches with 400-grit diamond, perpendicular to the testing direction of maximum tensile stress. Thus, all specimens were initially sized by machining in the direction to create the least severe flaws, but

finish machined in the direction yielding the most severe flaws. Machining was performed

on a Brown and Sharpe surface grinder at a wheel speed of 2860 rpm and a table speed of

20 ft/min. using a water/water soluble oil coolant. The eight-inch diameter metal bonded

wheel was a D400-L50M grading. Final specimen size was 1/8 x 1/8 x 1-1/8". After final

finishing, specimens were cleaned in acetone and dried in air at 120°C. Edge bevelling

was done manually on 600-grit SiC paper in the longitudinal orientation, the orientation

for introduction of the least severe flaws.

2.3. Testing procedure

After drying, the specimens were immersed in dry toluene prior to and during the

testing, the objective being to actually measure failure from surface machining flaws

and not extended flaws created by substantial subcritical crack growth. Dry toluene was

used because it has been reported to act as an inert environment with regard to slow

crack growth [10]. Actual testing was with 1/3-point loading in 4-point bending on an Instron Model TM commercial testing machine at a crosshead speed of 5 x 10"^ inches per

minute, a stressing rate of 1.7 x 104 MN/m2 min., or 2.5 x 106 psi/min.

456
3. Results and Discussion

3.1. Examination of machined surfaces

Tensile surfaces were examined by scanning electron microscopy to compare the surface

features of specimens finished at the different rates of surface removal, and to examine

the bevelled edges. It was found that for the greater rate of removal, the surface generally appeared "cleaner". The microstructure was more readily discerned, in that grains,

pores, grinding scratches, and some pullouts were more evident. In the case of the lesser

rate of material removal, the surface appeared much smoother, with some evidence of plastic

flow. There also appeared to be a large amount of small (<lym) particle debris scattered

across the tensile surface and within the pores open to the surface. Tensile surfaces of

AL600 are compared in figure 1 for each of the surface removal rates. These general observations applied for each of the three materials.

Figure 1. Resultant surface features of AL600 after

two rates of surface removal: a) 2.5 x 1Q-4 in/pass,

b) 2.5 x 10-5 in/pass.

3.2. Data Analysis

A statistical analysis of the results was employed to evaluate the relative changes in

the strength distributions. The two-parameter Wei bull distribution function, eq. (2), was

fitted to each of the six data sets so that the Weibull parameters could be directly compared. This method relates the range of strengths of a particular set of data to a probability of failure scale, which can be linearized by transforming eq. (2) to the following

logarithmic form:

457
I

In In (—c) z m In a f + [InS - m In a

Q

], 1-F J

y

'f

ax + b.

(3)

(3-a)

The Wei bull modulus, m, is then obtained from the slope, and the scaling parameter, a , is calculated from the intercept of the regression line.

3.3. Strength Data

Over one hundred data points for each set of conditions provided

population for a computer calculation of the least squares regression

Wei bull parameters [11]. The data sets were also computer plotted so shapes and locations of the distributions could be visually compared,

tributions for the three aluminas are shown in figure 2, with the two

a specific material on the same plot. The pertinent descriptive statistical parameters are listed in table 2. From the R^ values, it is evident that the two-parameter Wei bull distribution is a quite satisfactory description of the data. However, the vertical dropoffs of the data at the lower strengths for AL300 and AL600 does indicate a tendency toward

a good statistical

equations and

that the relative

The strength disremoval rates for

4.0*

I •1.0-

-6. 0.

(a)

4.0

High

Rate 1.0

-6,0

0.0 10,6

LN a.

11.2 4.0

(b)

0.0

> -1 .0

6.0

(c)

High

Rate ' Low

Rate

10.6

LN

11.2

10,0 10.6

LN a*

11.2

Figure 2. Comparisons of Wei bull distribution for the three aluminas

after two rates of surface removal, a) AL300, b) AL600, c) AL995.

a finite value of a in eq. (1), so that a three-parameter

fit for these data. u The nonlinearity in the high strength

central region of the AL600 data also suggests the possibi

than one flaw population as has been discussed by Johnson

fracture surfaces, flaw origins were sought but could not

is often the case with polycrystall ine alumina. Thus, mul

be conclusively confirmed. From figure 2 and table 2, it

of surface removal has a marked tendency to yield higher a Weibull modulus, m. The trend is consistent throughout wi

458

Weibull may yield an improved

region of the AL300 and the

lity of the existence of more

[12]. During SEM study of the

be conclusively identified as tiple flaw populations could not

is evident that the lesser rate

verage strengths and a higher

th the exception of the latter
for the AL300. From the differences that exist between the two different removal rates for

the three materials, it is evident that some type of material /mi crostructure dependence

exists, although the specific relationship is not obvious. The coarse grain size material,

AL300, exhibits nearly parallel distributions over the entire range of strengths, including

the low strength and the high strength tails. This indicates that the change in flaw

severity due to surface removal rate is the same for all strength levels. On the other

hand, the fine grain size AL600 material reveals a decreasing effect of the surface re- moval rate as the high strength region is approached. The high strength tails for both re- moval rates are nearly identical. Other than the low strength tail of the distribution

for the higher removal rate, the shapes for AL600 are very similar. In the case of AL995,

which contains some large 25ym grains, the strengths are comparable to those of AL300

(grain size ;25ym). This material is the least affected of the three by the surface re- moval rate during machining.

Table 2. Strength Distribution Parameters

Removal Ratps (\ in/DasO ill/ f/U J J /

Material Parameters 2.5 x 10" 4 -5

2.5 x 10

R

2 (%) 95.6 93.4

AL300 m 13.5 13.3

°o 25,850 29,650

~a

f (psi) 32,200 36,900

R

2

(%) 97.6 98.3

AL600 m 11.0 14.3

39,300 45,150

°f

(psi) 51 ,300 55,300

R

2 (%) 97.8 98.6

AL995 m 13.4 15.5

a

o

31 ,800 33,550

a

f (psi) 39,500 40,400

R

2 (%) 97.7 94.4

AL600 m 9.1 10.3

(Bevelled)

!<> 42,650 44,250

a

f (psi) 58,600 58,700

3.4. Effects of Edge-Bevelling

The AL600, the fine grain size, high strength material, was employed to determine the

effects of edge-bevelling on the resultant strength distributions. The statistical parameters for this comparison are listed in table 2, the distributions are plotted in figure 3,

and the bevelled vs. non-bevelled edges are compared in the micrographs of figure 4. The

distributions in figure 3 and the parameters in table 2 clearly reveal that edge-bevelling

yields an increase in average strength, an effect that is more pronounced for the higher

rate of surface removal. Edge-bevelling also caused a slight decrease in the Weibull

modulus, m, for both rates of surface removal. This effect probably results from not

459
Figure 3. Wei bull distributions comparing the strengths of

bevelled and non-bevelled specimens of AL600 after a) the

lesser surface removal rate, and b) the greater surface

removal rate.

removing the most severe edge chipping flaws. As a result the low strength tail is little

affected by bevelling, resulting in lower m- values but higher average strengths. Similarl

the bevelling process might be expected to tend toward a higher value of o

u in a threeparameter fit.

Micrographs comparing bevelled and non-bevelled edges are shown in figure 4. It is

evident that the bevelling process eliminated some of the edge chipping. It is likely

that since the removal of edge flaws by the bevelling procedure results in increased

strength for both removal rates, that the magnitude of the increase is probably related

to the edge damage present before bevelling. The distributions for each removal rate are shown in figure 3. The wider separation in distributions seen for the high removal rate

specimens indicates that more severe edge chipping occurred in this case than when the

lesser removal rate was employed.

4. Conclusions

A less severe rate of material removal during surface grinding results in higher

average strengths and an increased Weibull modulus. The differing shapes of the strength

distributions for the three aluminas indicates that microstructure has a decided effect

on the magnitudes of these increases. Specific details of this effect are not evident

at this time.

The Weibull two-parameter equation fits the strength data rather well, although from

the shapes of some of the distributions, such as for AL600, multiple flaw populations can- not be ruled out. Tendencies toward a finite value of a , as evident with AL300, indicate

that the three-parameter form may yield an improved fit 8ver that of the two-parameter

form.

Edge-bevelling can increase average strengths. Bevelling has a more pronounced effect

on the strength of materials subjected to high rates of surface removal, since this surface

finishing operation results in more severe edge chipping.

460
Figure 4. Comparison of a) non-bevelled and b)

bevelled sample edges of AL600.

The authors express their thanks to the Alcoa Foundation for their support of this

research, and also to Wesgo for supplying the alumina materials. .

5. References

[1] Busch, D. M., and Prins, J. F., A basic study of the diamond grinding of alumina

(Proc. Sci. of Cer. Mach. and Surf. Fin., NBS, Gaithersburg, Maryland, Nov. 2-4,

1970), NBS Special Publication 348, p. 73-87.

[2] Rice, R. W., The effect of grinding direction on the strength of ceramics (Proc. Sci.

of Cer. Mach. and Surf. Fin., NBS, Gaithersburg, Maryland, Nov. 2-4, 1970), NBS

Special Publication 348, p. 365-376.

[3] Mueller, J. A., Grinding - a means of shaping and sizing ceramics (Proc. Sci. of Cer.

Mach. and Surf. Fin., NBS, Gaithersburg, Maryland, Nov. 2-4, 1970), NBS Special

Publication 348, p. 45-52.

[4] Sedlacek, R., Halden, F. A., and Jorgensen, P. J., On the strength of ceramics as a function of microstructure, grinding parameters, surface finish, and environmental

conditions (Proc. Sci. of Cer. Mach. and Surf. Fin., NBS, Gaithersburg, Maryland, Nov.

2-4, 1970), NBS Special Publication 348, p. 89-97.

461
[5] McClintock, F. A., Statistics of brittle fracture, Fracture Mechanics of Ceramics 1 , Eds. R. C. Bradt, D. P. H. Hasselman, and F. F. Lange, p. 93-114 (Plenum Press, New

York, 1974).

[6] Batdorf, S. B., Fundamentals of the statistical theory of fracture, Fracture Mechanics

of Ceramics 3 , Eds. R. C. Bradt, D. P. H. Hasselman, and F. F. Lange, p. 1-30 (Plenum

Press, New York, 1978).

[7] Evans, A. G., and Jones, R. L., Evaluation of a fundamental approach for the statistical analysis of fracture, J. Am. Cer. Soc, 6_L [3-4] 156-160 (1978).

[8] Easier, T. E., previously unpublished results.

[9] Tressler, R. E., Langensiepen, R. A., and Bradt, R. C, Surface-finish effects on strength-vs-grain-size relations in polycrystal 1 ine Al

? 0o, J. Am. Cer. Soc, 57 [5]

226-227 (1974).

[10] Simpson, L. A., unpublished data, Atomic Energy of Canada, Whiteshell, Manitoba.

[11] Jeryan, R. A., Use of statistics in ceramic design and evaluation, Technical Report

X4-8515, Ford Motor Company, Dearborn, Michigan (1977).

[12] Johnson, C. A., Fracture statistics of multiple flaw distributions (Presented at the

Fall Meeting of the Basic Science Division of the American Ceramic Society, NBS,

Gaithersburg, Maryland, Nov. 12-15, 1978).

[13] Davies, D. G. S., The statistical approach to engineering design in ceramics (Proc.

of Br. Cer. Soc, [22], June 1973, p. 429-452).

[14] Weil, N. A., Bortz, S. A., and Firestone, R. F., Factors affecting the statistical

strength of alumina, Materials Science Research 1, Eds., H. H. Stadelmeier, and

W. W. Austin, p. 291-313 (Plenum Press, New York, 1963).

Discussion

FIRESTONE

Did you examine the specimens for fracture origins, specifically,

the beveled specimens which broke at lower stresses than the non-beveled

specimens from which it appears the specimens were damaged in bevelling?

BRADT

The fractures of specimens were examined. However, fracture origins

were not sufficiently well defined to give definitive results.

462
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

EFFECT OF SURFACE FINISH ON THE STRENGTH

OF HOT PRESSED SILICON NITRIDE

C. A. Andersson and R. J. Bratton

Westinghouse Research and Development Center

Pittsburgh, PA 15235

Statistical flexural strengths of hot pressed silicon nitride, Norton NC

132, were determined as functions of the stress axis to hot pressing direction,

the coarseness of grinding and the stress axis to grinding orientation. The

higher strengths of the material tested perpendicular to the hot pressing

direction rather than parallel to it can be attributed to a proportional

difference in the critical stress intensities between the directions. Coarser

grinding media introduces larger flaws and reduces strength. Finally, trans- verse grinding produces stress concentrating grooves, effectively deeper

cracks and possibly wider ones which reduce the applied stress necessary to cause propagation.

Keywords: Finishes; flexure; fracture; fracture mechanics; grinding, silicon

nitride; strength; surface.

1 . Introduction

The strength data reported in this paper were collected in support of a program 1 to

develop ceramic blading for industrial size combustion turbines [I] 2

. A major purpose of

that program was to design, test and evaluate blade attachment concepts. In order to enable

statisitcal failure predictions to be made of those blade forms, it was necessary to statistically determine the material strengths as functions of the orientation within the billets

from which they were machined, and the diamond grinding procedures used. Two hot pressing

orientations, three grinding meshes and two grinding directions, for a total of twelve com- binations, were evaluated. These data, as well as their assessment, are presented herein.

2. Experimental

Test specimens measuring .125 x .250 x 1.125 in. (.32 x .64 x 2.86 cm) were cut from

five billets of NC 132 hot pressed silicon nitride 3

, both parallel and perpendicular to the

pressing direction and were diamond ground either parallel or perpendicular to the stress

direction. Three diamond grit sizes were used: 325, 240, 80 mesh. Thirty 325 mesh ground

specimens were prepared for each of the pressing directions and grinding direction orientations, and fifteen specimens were prepared for each of the other grinding mesh/orientation

conditions

The specimens were flexure tested at room temperature in an Instron Universal Test

Machine using a steel fixture having inner and outer spans of 0.25 and 0.75 in. (0.64 and

1.91 cm), respectively. A crosshead speed of 0.05 cm/min (0.02 in/min) was used. Breaking

loads were autographi cal ly recorded.

Electric Power Research Institute, Research Project No. 421-1.

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

3 The Norton Company, Worcester, MA.

463
3. Results

2

The mean fracture strengths, variances, S , standard deviations, S, and the coefficient of variance, C vm , are presented in table 1. In agreement with previous investigators,

the strength is dependent on orientation within the billet, i.e., for all conditions tested,

the strength perpendicular to the hot pressing direction is 13 to 19 percent greater than

that parallel to the pressing direction. The scatter is also greater in the strong direction

as indicated by the coefficients of variation, C vm . With one exception, C vm perpendicular

to the pressing direction ranges between 10 and 100 percent greater than the parallel orientation; in that exception, it is 15 percent lower.

Comparing the effects of grinding coarseness parallel to the stress direction: there

is less than 5 percent difference between the 325 and 240 mesh surfaces, but a strength de- crease of 30 percent occurs when 80 mesh grit is used. The grinding direction transverse to

the stress axis orientation has a more profound effect: a strength decrease of 35 to 50 percent occurs for this grinding condition. For the worst cases tested (transverse grinding

with 80 mesh grit) there is in excess of a 60 percent strength decrease when compared to the

best cases tested (longitudinal grinding with 325 mesh grit).

In addition to the Gaussian distribution analyses, the data were also analyzed by two

parameter Weibull statistical methods. All the data were assessed assuming that the failures were caused by flaws existing on the surface. In addition, the data for specimens whose

surfaces were ground parallel to the stress axes were analyzed assuming that the defects were contained within their volumes.

For surface flaws the probability of failure, P^, is given by:

P (1) A

= 1- exp K

A A(a/a0

) m

where a is. the stress, m is the Weibull modulus, a0 is the Weibull characteristic strength

for surface flaws, is a constant dependent on specimen and loading geometries and A is

the specimen surface area. The corresponding probability of failure for volume flaws, Py, is

P v

= 1- exp -K

v V(a/a

0v

) mJ

(2)

where K

v is the geometric constant, V the specimen volume and a0 the characteristic strength

for volume flaws. Figure 1 contains plots of the data based on these equations. The Weibull

parameters were determined statistically and the results are contained in tables 2 and 3.

The lines in the figure represent the results of the analyses.

Comparing the data on figures 1(b) and 1(d) for the grinding directions transverse to

the stress axes, the data with one exception seems to be reasonably well behaved. The

exception is the 325 mesh ground data in which the stress axis is parallel to the pressing

direction. These data (fig. 1(b)) have a lower Weibull modulus than the other two mesh

sizes for this orientation. Thus, the probability of failure at stresses lower than 33 ksi

(230 MPa) is greater for the 325 mesh than for the 80 mesh material. The physical cause for

this is uncertain, but the fine grinding procedure seems to have put a wider dispersion of

defect sizes on the surface.

4. Discussion

Brittle materials are known to fail from preexistng flaws within the body being loaded,

and a fracture mechanics approach is generally accepted as the best criterion for failure.

This criterion simply states that the applied stress, a, at which a body fails is dependent

not only on a property which is characteristic of the material (Kj c

, the critical stress

intensity parameter) but also on the flaw size (2c). Thus,

a = Y K Ic

//c~ (3)

where Y is a constant which accounts for the flaw shape and location, the geometry of the

loaded body, the type of loading, etc.

464
Table 1. Room temperature flexural strengths and Gaussian distribution parameters of Norton NC 132 silicon nitride as functions of: 1) the orientation of the stress axis with respect to the hot pressing direction, 2) the

surface grinding direction, and 3) the grinding mesh size.

Stress Axis Grinding Direction Longitudinal with Transverse to

to Pressing

Direction

Orientation

*~-^^^ Grinding

^~~~~---^Me s h

Parameter:

Stress Axis Stress Axis

325 240 80 325 240 80

Parallel Mean Fracture

Stress, o f

, ksi

(MPa)

104.0

(717)

101.3

(698)

74.3

(512)

67.4

(465)

51.4

(354)

39.5

(272)

2 2 Variance, S , ksi

(MPa2

)

100.2

(4763)

57.1

(2715)

190.1

(90 38)

70.2

(3336)

9.1

(433)

2.8

(133)

Standard Deviation,

b , KSI

(MPa)

10.0

(69)

7.6

(52)

13.8

(95)

8.4

(58)

3.0

(21)

1.7

(12)

Coefficient of

Variation, Cvm.100 .075 .186 .124 .059 .043

Number of

Specimens, n 30 15 15 3.0 15 15

Perpendicular Mean Fracture

Stress, Of, ksi

(MPa)

123.2

(849)

116. 7 (805)85.3

(588)

72.7

(501)

59. 2 (408)

47.2

(326)

r.2 , -2 Variance, S , ksi

(MPa 2

)

zuy . _> (9960)

y 3 . 0

(4516)

TnQ "7 juy . / (14720)

CO c jO.O

(2785)

i *3 a j j . o (1597)

ib • y

(803)

Standard Deviation,

S, ksi

(MPa)

14.5

(100)

9.7

(67)

17.6

(121)

7.7

(53)

5.8

(40)

4.1

(28)

Coefficient of

Variation, C,_„ vm.117 .083 .206 .105 .098 .087

Number of

Specimens, n 30 15 15 30 15 14

Ratio: o

f (J_) :5

f

(||) 1.18 1.15 1.15 1.13 1. 15 1.19

Ratio: C ( J) :C (II) 1.17 1.11 1.11 .85 1.61 2.02

- . 2

(a-a

f

)

f (a

f

)

0- =

(2,)

1/2

S

N

^ In

exp

2S

2

S =

1jrr .v

a

£i

-G

f

) 1=1 x

P- = 1 -

(2n) 1/2

s a exp

465
Stress, MPa

100 200 300 400 500 600 800 1200

Stress, ksi

(a) Tensile axis oriented parallel to the

pressing direction; grinding direction

parallel to the tensile axis.

Stress, MPa

100 200 300 400 500 600 800 luuu

1200

Stress, ksi

(c) Tensile axis oriented perpendicular to

the pressing direction; grinding direction parallel to the tensile axis.

Stress, MPa

20 30 40 50 100 150 200

Stress, ksi (b) Tensile axis oriented parallel to the

pressing direction; grinding direction

perpendicular to the tensile axis.

Stress, MPa

300 4i

20 30 40 50 100 150 200

Stress, ksi

(d) Tensile axis oriented perpendicular to

the pressing direction; grinding direction perpendicular to the tensile axis.

Figure 1. Probability of failure of Norton NC 132 as functions of flexural stress, hot press

ing direction to stress axis orientation, grinding mesh size and grinding direction to

stress axis orientation.

466
Table 2. Surface flaw Weibull distribution parameters for Norton NC 132 silicon

nitride at room temperature as a function of: 1) the orientation of the tensile

axis with respect to the hot pressing direction, 2) the surface grinding direction, and 3) the grinding mesh size.

Stress Axis Grinding Direction

Longitudinal with Transverse to

to Pressing

n i T(=>r"t~ i on Orientation

Grinding

^^"""""••w Mesh * Parameter:

Stress Axis Stress Axis

iZD z4U OU "3 *> C 80

Parallel Weibull Modulus, mA

12.9 17.1 6.7 9.8 22.0 30.5

Characteristic

Strength, o q

, ksi

A (MPa)

88. 8

(612)

91.2

(629)

57.

(394)

54.6

(376)

47.0

(324)

37.2

(256)

Perpendicular Weibull Modulus, m 10.4 15.3 5.9 12.1 12.9 14.7

Characteristic

Strength, o Q

, ksi

A(MP a)

99.8

(688)

104.1

(718)

62.8

(433)

63. 3 (436)

50.6

(349)

41.7

(288)

P, = 1 - exp[-/ (a/a ) dA]

A A °A

Table 3. Volume flaw Weibull distribution parameters

for Norton NC 132 silicon nitride at room temperature

as a function of: 1) the orientation of the tensile

axis with respect to the hot pressing direction, and

2) the grinding mesh size.

Stress Axis Grinding Direction: Longitudinal with

to Pressing

Direction

Orientation

^^^^ Grinding

^"^--^Mesh:

Parameter:

Stress Axis

325 240

Parallel Weibull Modulus, m^ 12.9 17.1

Characteristic

Strength, o q

, ksi

V(MPa)

58.3

(402)

65.4

(451)

Perpendicular Weibull Modulus, 10.4 15.3

Characteristic

Strength, a , ksi

°V(MPa)

60. 3 (416)

72.2

(498)

P, = 1 - exp[-/ (a/a

f

J A V

m

) v

dv]

°v

467
The effect on strength of stress axis to hot pressing direction for the silicon nitride

tested can be attributed simply to the difference in the critical stress intensity factors

between the orientations. Since the strengths perpendicular to the hot pressing direction

range between 13-19 percent greater than those parallel to it for identical grinding mesh

and grinding direction conditions, it can be assumed that the critical stress intensities

differ proportionally for these pressing orientations. On the average, the Kj perpendicular

to the hot pressing direction can be taken to be 16 percent greater than that parallel.

The effects of the coarseness of the grinding media and grinding direction with respect

to the stress axis are more complicated. For example, it is expected that the size of the

microcracks introduced during grinding, which will act as the sources of failure, will be

dependent on the grinding mesh size. In addition, flaw shapes would be expected to be a function of the grinding direction; i.e., the cracks should be longer parallel to the grinding direction than perpendicular to it. Alternatively, parallel to the grinding direction,

several closely spaced, uniformly shaped cracks could effectively act as a single long one. Finally, the grooves formed by the grinding process can act as additional stress concentrators when oriented perpendicular to the stress direction, thus changing the effective stress

intensity factors.

The effect of grinding on the magnitude of flaw size is a complicated one beyond the

scope of the present study. Other than noting that flaw size is expected to be somewhat

proportional to grinding media size, it will not be discussed. However, grinding direction

to stressing direction will be considered by assessing its effect on the proportionality

constant in the fracture mechanics equation.

In the following discussion, it will be assumed that failure occurs from flaws at

tensile surfaces of the bend specimens tested. Scanning electron micrographs of fracture

surfaces such as those shown in figure 2 confirm this assumption 4

. Occasionally, failures

can originate from defects contained within the volume of finely ground specimens, but

these are exceptions to the rule in the present study. The procedure used will thus be to

establish the stress intensity relationships for surface flaws in bending and then to determine the stress concentration effects of the grinding grooves on these relationships.

(a) Stress axis parallel to grinding direction (b) Stress axis perpendicular to grinding

direction

Figure 2. Typical flexural fracture surfaces of hot pressed silicon showing origins of

failure. The bottom portion of each photo is the tensile surface of the flexure specimen.

Note that the origin of failure for the stress axis parallel to the grinding direction is

nearly semicircular, whereas it is spread parallel to the crack edge for the stress axis

perpendicular to the grinding direction.

468
4.1. Stress intensity relationships for surface flaws in bending

The analysis of Kobayashi [2] for the stress intensity relationships in bending for

semielliptic surface flaws oriented normal to the tensile axis was used. He graphically

presents these as functions of the ratio of crack depth, b, to crack width, a, and of the

ratio of crack depth to specimen thickness, h. It is assumed throughout the following, that

the crack depth is small compared to the specimen thickness. Therefore, the values were extrapolated to b/h equal to zero and replotted in figure 3.

Figure 3. Stress intensity coefficients as function of the crack width to crack depth

ratio for shallow semielliptical surface

cracks in bending [2].

In bending, the stress intensity relationship generally differs along the periphery of

the crack. For example, for a semicircular crack (a/b = 1), the stress intensity coefficient

in the direction of the a-axis is 1.41, whereas it is 1.14 at the maximum crack depth along

the b-axis. Therefore, on loading, it would be expected that the critical stress intensity

at the surface would be arrived at first and that the crack would first start growing in the

direction parallel to the surface. As the crack grows, the a/b ratio changes, causing the

stress intensity at the surface to decrease and that at the maximum depth to increase until

a/b reaches a value of 1.23. At this ratio, the value of the coefficient is 1.35 for both

orthogonal directions and crack propagation proceeds in both directions.

4.2, Stress intensity relationships for grinding

direction parallel to the stress direction

Since the grinding grooves parallel to the stress direction do not contribute to the

stress concentration, the stress intensity relationships given in figure 3 are applicable.

Also, since the elliptical axes are usually nearly equal, the previous reasoning regarding

crack growth can be applied and the relationship,

Kj = 1.35 a

b

/b, (4)

can be used.

469
4.3. Stress intensity relationships for grinding direction

perpendicular to the stress direction

Surface grooves such as would occur from grinding are known to cause stress concentrations which should affect the stress intensity relationships of surface cracks associated

with them. A simple model based on some work that was done on two dimensional cracks [3] was devised to determine the stress intensity relationships as functions of groove and crack

dimensions. In this model, a semiel 1 i pti cal crack is assumed to lie at the bottom of a semicylindrical groove. Using the previous notation, the a-axis of the crack is parallel to the

axis of the semicylinder which, in turn, is perpendicular to the stress axis in bending. Both

the groove radius, r, and the crack depth, b, are considered to be small compared with the

specimen thickness, h, and the stress intensity relationships given in figure 3 can be used.

Parallel to the surface, at the bottom of the groove, the stress is multiplied by the

stress concentration factor, k^., of the groove (kj- % 3), and the relationship in the a-axis

direction is:

Kj = f(a/b, b/h = 0)k

t a

b

/b . (5)

An example of the function f(a/b, b/h = 0)k for a/b = 2 has been plotted versus b/r in

figure 4 as the long dashed/short dashed curve.

The stress concentration effect of the groove at the maximum crack depth is dependent

on the relative size of the crack with respect to the groove. For very short crack depths,

b, the crack will be influenced by the full stress concentration, k^, and

Kj = g(a/b, b/h = 0)k

t

o

b

/b . (6)

At Surface

Figure 4. Stress intensity coefficients in bending as a function of the crack depth to groove radius for a shallow

semiel 1 i ptical crack at the bottom of a semi cyl indrical

groove, and a crack width to depth ratio of 2.

470
For very deep cracks, with respect to groove radius, the stress concentration will be negligible and the only influence of the groove is to lengthen the crack by an amount equal to

the groove radius, and:

Kj = g(a/b, b/h = 0)a

b

/b + r . (7)

Equations (6) and (7) for a/b = 2 have been plotted as the solid line segments in figure 4.

A smooth short dashed line representing the equation,

Kj = h(a/b, b/h = 0)o

b

/b , (8)

has been drawn in this figure to represent intermediate crack depths.

Similar treatments for other a/b values were made to construct the stress intensity

relationships given in figure 5. The long dashed/short dashed curve represents the stress

intensity parallel to the bottom surface of the groove in the a-axis direction. The solid

lines are the stress intensities, for various b/r values, at the maximum crack depth in the

b-axis direction.

0 I 1 1 1 1 L

0 2 4 6 8 10 a/b

Figure 5. Stress intensity coefficients in bending

as functions of the crack width to crack depth

ratio and the crack depth to groove radius ratio

for shallow semielliptical surface cracks at the

bottom of semi cylindrical grooves.

The consequences of figure 5 will now be examined. The stress concentration effect of

transverse grooves is most pronounced at small b/r ratios and becomes negligible at large b/r

471
Whether the Kj for a given applied stress is the greatest parallel to the surface or in

the transverse direction at the maximum crack depth is dependent on both ratios, b/r and a/b.

For example, for b/r = 1 and a/b = 1, the stress intensity at the surface is 2.5 times that

at the maximum crack depth. On increasing the applied stress, the critical value would be

reached at the surface first and the crack would grow laterally to an a/b ratio of 4.6 before

fracturing transversely. On the other hand, for the same b/r = 1, and for elliptical crack

axis ratios greater than 4.6, the Kj at the maximum crack depth would become critical first,

and transverse fracture would proceed initially.

For reasonably large crack depth to groove radius ratios (b/r > .5), the shape of the

crack does not have a significant effect on the stress intensity relationships. For example,

consider two cracks, one having an initial a/b = 1 and the other an a/b = «>, at the bottom

of a groove with b/r = 1. The ratio of the stress intensities would be:

Kj(a/b = °°) 2>81

Kj(a/b = 1+4.6) 2.65 1.06, (9)

since the initial semicircular crack would be expected to grow laterally before transverse

failure until a/b reaches 4.6. Thus an initially semicircular crack would be expected to

sustain only six percent greater applied stress than one having an infinite a/b ratio for

the same crack depth, b. In the absence of the groove, the stress intensity ratio would be:

K

T (a/b = -) i go

' = 1.47, (10) Kj(a/b = 1+1 .23) 1 .35

and the initially semicircular crack could sustain 47 percent greater stress than the

elongated one, if both cracks had the same depth.

The insensitivity of the stress intensity relationships to the shape of the cracks in the

presence of grooves, is one factor that can contribute to the lower scatter in strength

often observed for transversely ground materials. For example, in the present test series,

the coefficients of variation for 80 mesh transversely ground specimens were .09 and .04,

whereas for the parallel ground specimens they were .21 and .19.

4.4. Assessment of test results

In concurrence with previous investigations of hot pressed silicon nitride [4], the

difference in strength parallel and perpendicular to the hot pressing direction is attributable to the difference in the stress intensity factors between the directions. The

strength ratios parallel and perpendicular to the pressing direction ranged between .82 and

.94 for six grinding directions averaging 0.86. It is reasonable to assume that this value

can be used for the ratio of the critical stress intensities in the two directions. The Kj

of this material was measured to be 4.6 MPa nv/2 (C v = .09) with the plane of the crack c normal to the pressing direction [1]. The orthogonal Ki c is estimated to be 5.4 MPa trr' .

The effect of the coarseness of the grinding mesh on the strengths of materials ground

parallel to the stress axis is on the maximum flaw depth. Using the appropriate strengths

and the Kj values above, the mean maximum crack depths, b, were calculated and the results

presented Tn table 4.

The effect of the coarseness of the grinding mesh on the strengths of materials ground

perpendicular to the stress axis is more complicated since the transverse grinding operation

causes both cracks and grinding grooves to have affect. However, calculations show that the

crack depth to groove radius, b/r, for each case tested was greater than unity. Therefore,

the maximum crack depth is out of the stress concentration field of the groove, and the major

contribution of the groove to the crack in the b-axis direction is to increase its depth.

The effective depth is thus the sum of b and r. It has also been pointed out that for b/r

greater than one, there is little difference in the stress intensity equations between the

cases where the crack grew laterally to a maximum value at fracture or where a long crack

preexisted. The effective crack depths were calculated for two extreme conditions: one for

472
Table 4. Calculated mean flaw depths of Norton NC 132 silicon nitride as functions of:

1) the orientation of the stress axis with respect to the hot pressing direction,

2) the surface grinding direction, and 3) the grinding mesh size.

Stress Axis

to Pressing Mean

Grinding

Direction

Longitudinal with

Stress Axis

Transverse to

Stress Axis

Direction

Orientation

Maximum

Crack Depth, um Grinding

325 240 80 325 240 80

Parallel 22 24 45

beff

= b+r, b/r = a/b = *± . o A O 81

beff

= b+r, b/r =

a/b =

00

I oqC 25 43 72

Perpendicular

b

a

beff

" b+r, b/r = a/b = 4.6b

22 24 46

33 50 78

beff

= b+r, b/r = a/b =oo

t 29 44 69

aKj = 1.35 a it

b i Kz = 1.87 a /beff

°KI = 1-99 a \ff

which b/r = 1 and an initially semicircular crack grew at a/b = 4,6, and the other for which

b/r = oo and the initial crack extended across the specimen width. These values are also

given in table 4.

For longitudinal grinding, coarsening the grinding mesh from 325 to 240 to 80 mesh

causes the mean maximum crack depth to increase from 22 to 24 to 46 ym, respectively. For

transverse grinding the effective crack depths are considerably larger, averaging 29, 46 and

75 \im for the same respective mesh sizes. The transverse depths are thus approximately 30,

90 and 60 percent larger than their corresponding longitudinal grinding depths. A fraction

of this difference can be attributed to the groove radius; however, both examination of the

fracture surfaces and calculations show the transverse ground crack depths to be actually

larger than the longitudinal ones after subtracting r. Thus, other factors related to the

grinding techniques used must have contributed to the differences.

5. Summary

Flexural strengths of hot pressed silicon nitride, Norton NC 132, were determined as functions of the stress axis to the hot pressing direction, the coarseness of grinding media

and the stress axis of grinding orientation. The strengths were statistically analyzed using

both Gaussian and Weibull methods. Both the strengths and distributions were shown to be

sensitive to each of the parameters. Strengths are the highest for materials finely ground

parallel to the stress axis and tested perpendicular to the hot pressing direction. The best

case tested was over 300 percent stronger than the worst case tested. The least amount of

scatter occurred for the coarsest mesh specimens ground perpendicular to the stress axis.

The stress intensity relationships for surface flaws in bending were presented for con- ditions of both grinding parallel and perpendicular to the tensile axis. In bending, the

473
stress intensity is generally different around the crack periphery. When the stress

intensity is greater at the surface than at the maximum crack depth, a crack will initially

grow laterally, simultaneously decreasing the stress intensity at the surface and increasing

the stress intensity at the maximum depth until its critical value is reached. If the stress

intensity is greater at the maximum crack depth than at the surface, as in the case for

wide cracks, transverse crack propagation will occur from the start.

Grinding perpendicular to the stress axis causes loss of strength for several possible

reasons. Crack lengths parallel to the grinding grooves can be deeper and longer than those

perpendicular. The groove depths can increase the effective crack depths. Grooves can act

as stress concentrators, increasing the effective stress intensities. Each of the above

effects reduce the applied stress necessary to reach the critical stress intensity and subsequent crack propagation.

The authors wish to thank Dr. A. Cohn and the Electric Power Research Institute for

supporting the program for which the data presented herein was obtained. In addition, they

wish to thank Mr. J. J. Nalevanko and Mr. W. J. Carmen for their experimental contributions.

References

[1 ] Andersson, C. A., Bratton, R. J. and Booher, Jr., C. R. , Ceramic turbine components

research and development, part 1 - ceramic rotor blade development, Final Report,

Electric Power Research Institute Research Project 421-1, (Dec. 1978).

[2] Kobayashi, A. S., "Crack opening displacement in a surface flawed plate subjected to

tension or plate bending, (Proc. 2nd International Conf. on Mechanical Behavior of

Materials, Boston, Mass., August 1976), pp. 1073-1077.

[3] Dowling, N. E., Fatigue at notches and the local strain and fracture mechanics approaches

Paper presented at the 11th National Symposium on Fracture Mechanics, Blacksburg, VA,

(June, 1978).

[4] Miller, D. G., Andersson, C. A., Singhal, S. C, Lange, F. F., Diaz, E. S., Kossowski , R. and Bratton, R. J., Brittle materials design, high temperature gas turbine material

technology, Final Report, Advanced Research Projects Agency, Contract No. DAAG-46-71 -C0162, AMMRC CTR 76-32, Vol. IV, (Dec. 1976).

i

f

474
Discussion

RICE:

First, I understand that you consider a semi-cylindrical gouge in the surface,

i.e., a gouge with a width twice the depth. If this is correct, this leads to two questions.

The first being what justification do you give for this in view of the extensive evidence

shown in earlier work as well as quite extensively at this conference that groove width to

depth ratios are typically _> 4 due to substantial fracturing along the path of the grit

particles. The second related question is what effect would width to depth ratios of fourto-one and six- to-one have on your resultant calculations.

Secondly, I had concerns about your proposed failure criteria. As I understand it, you

defined failure when the stress intensity for propagation of the flaw parallel with the

surface equals the stress intensity for propagation inward from the surface. However, this

appears to violate the normal failure criteria of the net stress intensity on the original

flaw reaching the critical stress intensity value for the material. Thus, for example, if one had failure when the stress intensity was equal in two perpendicular directions, this

would appear to imply that the propagating crack front would tend to preserve the shape of

the flaw at the time of failure. This, however, is not observed since typically fracture

mirror and related markings which outline the crack front characteristically approach a circular profile in a uniform stress field regardless of the shape of the original flaw.

Your comments?

ANDERSSON :

With regard to your first question, as to the use of the semi-cylindrical

gouge for my model, it was, of course, selected for the sake of simplicity in an attempt to

explain observed trends. The model also assumes that fracture will occur from a single gouge

or alternatively a gouge that is deeper than its neighbors since it ignores stress field

interactions from neighboring gouges. However, the only change that would be required to

account for different shaped gouges or gouge interactions is a change in the value of stress

concentration factor, k^, used in the equations. For example, using textbook values for

k t [1], it would change from 3.0 to 2.7 to 2.3 on increasing the groove width to depth ratio

from 2 to 4 to 6, while maintaining the radius of curvature at the groove bottom. The major

effect that this, in turn, would have would be to reduce the surface stress intensity

coefficients (the dashed line in Fig. 5) by 10 and 25 percent, respectively. This, in turn,

would reduce the intercept a/b ratio with the maximum crack depth stress intensity coefficients. The following table summarizes the effects for an exemplary case:

Crack Width to Depth Ratio

Parameter

2 4 6

k

t 3.0 2.7 2.3

b/r 1 1 1

a/b 4.6 3.7 2.4

Kj/a

b

,/b" 2.65 2.60 2.50

Thus, there is a two percent decrease in the stress intensity coefficient on using a width

to depth ratio of 4 and a six percent decrease on using a ratio of 6. It can be concluded

that the semi-cylindrical groove model is adequate, but may be adjusted if still more

accurate calculations are desired.

The second question that you raised concerns the failure criteria that I use. In this

paper, the failure stress has been defined as the stress at which the stress intensity factor

orthogonal to the tensile surface (in the b-axis direction) reaches the critical value. For

very wide cracks, the stress intensity in this direction is always the greatest and there

exists no problem between definitions. However, for narrow cracks under the conditions

described in this paper, the stress intensity is larger in the a-axis direction (parallel to

the tensile surface) than in the b-axis direction, and other directions would be intermediate

475
to these. Therefore, on loading, subcritical crack propagation parallel to the surface would

be expected to occur first with no propagation perpendicular to the surface. For example, assume the following initial parameters:

a = 50 ym

b = 25 ym

r = 25 ym

K T = 5 MPa m 1/2

From figure 5:

K, (a-axis) = 3.54 o

Kj (b-axis) = 2.37 a /b

On loading to a stress value of 282 MPa:

Kj (a-axis) = 5 (critical)

Kj (b-axis) = 3.35

The crack would propagate in the a-direction. If it arrested at a = 75 ym,at this

stress value then:

Kj (a-axis) = 3.10 a Jb = 4.37 MPa m 1/2

K

z (b-axis) = 2.56 a v¥ = 3.60 MPa m 1/2

,

both of which are below the critical value. Thus, the load must be increased to 323 MPa

before the critical value is again reached in the a-axis direction. At 323 MPa:

Kj (a-axis) = 5 (critical)

Kj (b-axis) = 4.13.

This process will continue until the a-axis dimensions reach 115 ym at which point

Kj (a-axis) = Kj (b-axis) = 2.65 a /b

and at a stress of 377 MPa, the crack will catastropically propagate in both orthogonal

directions, since an increasing b value in the equations will maintain the criticallity.

Finally, with regard to your comments on the shape of a propagating crack front, two

points must be made. First the values presented in this paper were for small crack depths,

b, with respect to the overall specimen thickness, h, and as such pertain to the shape of

the crack at the time of initial propagation. As propagation proceeds b/h changes and must

be accounted for in the stress intensity relationships (as has been done by Kobayashi in

the paper used). These relationships would then dictate the crack front shape. Second, as the crack depth, b, to groove radius, r, ratio increases, the effectiveness of the stress

concentration of the groove diminishes and the relationships presented break down. The

stress intensity relationships of a propagating crack would be expected to revert to the

usual form (fig. 3) and result in nearly semicircular fronts. Therefore, I find no contradiction between the present results and your observations.

Reference

[1] Peterson, R. E., Stress concentration factors, Fig. 38, p. 57, John Wiley & Sons,

New York, (1974).

476
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

THE EFFECT OF SURFACE FINISHING ON THE STRENGTH

OF COMMERCIAL HOT PRESSED Si3N

4

C. Cm. Wu and K. R„ McKinney

Naval Research Laboratory

Washington, D. C. 20375

Effects of diamond grit size on the as-ground and oxidized

strength of commercial hot-pressed Si3N4 were studied. Finer

grit sizes gave higher as-ground strengths. Prolonged high

temperature oxidation negates effect of machining on strength.

Key Words: Silicon-nitride j machining of ceramics; fracture

strength; oxidation and degradation; and fracture origin.

1. Introduction

Hot-pressed Si3N4 is a promising candidate for ceramic heat engine

parts [l,2]l, but machining is needed to meet the close dimensional tolerances of some parts. Strengths can often be increased by finer finishing,

but may be effected by subsequent oxidation. The purpose of this work was

to investigate the machining and oxidation effects on the strength of commercial hot-pressed Si3N4.

2. Experimental

Flexure bars were cut from hot pressed billets of Norton NC-132 Si3N4

such that the resultant fracture surface would be parallel to the hot

pressing direction. The fracture would propagate perpendicular to the hot

pressing direction, i. e. the bars were cut to reflect the possible greater

strength of the billet if it had a preferred orientation of grain parallel

with the hot press plunger surface. The bars were then diamond ground to

a final cross section of ~ 2.5 mm x 7.5 mm with metal bonded grinding wheels

of 180, 320, and 600 grit sizes. All surface grinding was parallel to the

length of the bars. The 180 grit grinding was on a Boyar Schultz Hydraulic

Surface Grinder with a 200 mm (8 inches) grinding wheel rotating at 1725 rpm

and automatically translated at 760 mm/min with a feed rate of 0.025 mm per

pass. The 320 and 600 grit grinding was on a Motor Appliance Corp's grinder

with 100 mm (4 inches) wheel running at 3450 rpm translated manually at a speed of ~ 760 mm/min with a feed speed again of 0.025 mm per pass. After

surface finishing, the edges were rounded off using a 600 grit bonded diamond grinding disk moving lengthwise along the bars to decrease the probability of failure initiating from the corners. Some ground samples were

oxidized in an open end tube (AI2O3) furnace at 1430°C for 100 hrs. After

oxidation, some of the specimens were refinished with the 180 grit grinding

wheel until the surface appeared comparable to those attained before heat

treatment. Flexure strengths for each were measured in an Instron machine

with a 3-point flexure span of 12.7 mm (0.5 inches) and a cross head speed

of 1.27 mm/min (0.05 in/min)

Figures in brackets indicate the literature references at the end of this

paper.

477
3. Results and Discussions

The results of the strength measurements are summarized in table 1 0 Samples were from more than one billet, with billet-to-billet differences

not significant in comparison to the finishing effects below. While there

was no clear distinction between strengths of bars as ground with 180 or

320 grit diamond, there was a definite increase (e.g. 15%) in the bars as

ground with the 600 grit wheel.

Table 1. Strength of NC-132 in various machining conditions.

Conditions

Grinding

Wheel

Grit Size

Strength . Standard Deviation Typical Failure Origin

MPa ksi MPa ksi

As

Machined

i fin 884

(18) a 128 70 10 Machining flaws

320 875

(5)

127 54 8 Machining flaws

600 1032

(8)

150 74 11 Internal flaws or machining flaws

Oxidized

at 1430°C

for

100 hours

after

machining

180 470

(13)

68 68 10 Bubble-like reaction pits

320 492

(7)

71 24 4 Bubble-like reaction pits

600 508

(7)

73 44 7 Bubble-like reaction pits

Ref inished

after

Oxidation

180 962

(ID

140 25 4 Machining flaws

Numbers in parentheses indicate number of specimens tested.

Fractographic studies of bars tested as ground with the 180 grit wheel

showed fracture markings such as mirror and crack branching patterns indicating surface fracture origins associated with a deeper surface grinding

groove, figure 1. Similar fracture origins were observed for 320 grit ground

bars. While it was not possible to unequivocally define the specific flaws

themselves, from all the indications, these failures clearly appeared to

be from machining flaws introduced by and approximately perpendicular to the

deeper grinding grooves, consistent with other observations [3,4], On the

other hand, 600 grit finished bars showed a mixture of fracture origins from

either internal or surface flaws. The implication is that the hot-pressed

material contains processing defects of comparable size to those introduced

by 600 grit size diamond grinding, i.e. ~ 5 jum . From the economical point

of view for machining of this material into final configurations, it is then

obvious that any machining better than 600 grit size finishing is unnecessary

as far as strength is concerned.

The flexure strength of specimens after 100 hours at 1430° heat treatment was about 485 MPa (70 ksi) for all three groups of specimens regardless

of finishing prior to heat treatment. This was due to a thick porous layer

that had formed on the specimen surface during the heat treatment. This

layer contained numerous pits of different sizes that penetrated into the

478
Figure 1. Fracture origin in a 180 grit ground specimen

as seen in SEM. The origin is beneath the surface under

a deep machining groove. Arrows show boundary where crack

branching began.

bulk of the material. Upon loading, these pits then acted as fracture origins as has been reported [5,6,7] and also shown in figure 2. Some pits

have reached about 100 [i in size. No original machining introduced flaws

were observable. One can see that the prolonged high temperature treatment

has completely nullified the effect of machining on strength of the material.

479
Figure 2. Oxidation pits become fracture origins in

NC-132 Si3N4 specimens oxidized in air at 1430°C for

100 hrs regardless of finishing prior to oxidation.

In order to measure the strength level of the bulk of the oxidized

material, the oxide layers were ground off with the 180 grit metal bonded

wheel. The flexure strength measured reached slightly higher than the

original at about 962 MPa (140 ksi) with smaller standard deviation. This

perhaps is not surprising if we consider the fact that the impurities or

second phases that were present may have dispersed more evenly or even

diffused out of the body due to prolonged heat treatment at high temperature. At the same time, internal flaws may have also redistributed or

shrunk in size and therefore higher strength could be obtained. Typically,

fracture originated from cracks extending from the base of a deep machining

groove

.

One interesting question is: How far into the specimen does the damage

extend due to the oxidation treatment? Although a definitive conclusion

still cannot be drawn from the result presented here, from the fact that the

strength has completely recovered to its original value after refinishing,

it appears that the damage may be confined to as deep as the reaction pits

have grown into the surface, i.e. on the order of 100 yaa.

4. Summary

Flexure strength of 1200 MPa (160 ksi) for commercial hot-pressed

Si3N4 is attainable through grinding with fine (600) grit metal bonded diamond wheels. Further increases will be limited since even at this stress

level, internal processing defects were becoming frequent sources of failure.

Therefore, grinding with finer grits does not appear productive. For

specimens to be used in high temperature, fine machining is not needed since

high temperature oxidation significantly degrades strength, e.g. by ~ 50%

regardless of original grinding.

480
The authors would like to gratefully acknowledge R. W. Rice and

P, F. Becher for their invaluable discussions.

References

[l] Richerson, D. W. , Materials characterization program overview, (Proc

of the 1977 DARPA/NAVSEA Ceramic Gas Turbine Demonstration Engine

Program Review, Maine Maritime Academy, Castine, Maine, Aug. 1-4, 1977)

pp. 187-91.

[2] Richerson, D. W. and Yonushonis, T. M. , Properties of silicon-nitride

rotor blade materials, ibid, pp. 193-217.

[3] Rice, R. W. , Freiman, S. W. , and Mecholsky, Jr., J. J., Fracture sources

in Si3N4 and SiC, ibid, pp. 665-688.

[4] Rice, R. W. , Freiman, S. W. , and Mecholsky, Jr., J. J., Fractography of

Si3N4 and SiC (Ceramics for High Performance Applications-II , Proc. of

the Fifth Army Mat. Tech. Conf., Newport, R.I., March 21-25, 1977),

pp. 669-687.

[5] Freiman, S. W. , Mecholsky, J. J., McDonough, W. J., and Rice, R. W„,

Effect of oxidation on the room temperature strength of hot-pressed

Si3N4 -MgO and Zr02 , ibid, pp. 1069-1076.

[6] Richerson, D. W. and Yonushonis, T. M. , Environmental effects on the

strength of silicon-nitride, (Proc. of the 1977 DARPA/NAVSEA Ceramic

Gas Turbine Demonstration Engine Program Review, Maine Maritime Academy, Castine, Maine, Aug. 1-4, 1977), pp. 247-271.

[7] Freiman, S. W 0 , Wu, C. Cm., McKinney, K. R., and McDonough, W. J.,

Effect of oxidation on the room temperature strength of silicon-nitride

hot-pressed with MgO or ZrC>2, ibid, pp. 655-663.

Discussion

GRAHAM

Is the improved strength (140 ksi) for the oxidized & machined (180

grit) as compared to the strength (128 ksi) for the machined (180 grit)

unoxidized samples significant, and if so how do you explain the improvement? (If it isn't significant, is the 150 ksi for the 600 grit, unoxidized

a significant difference from the 180 grit, unoxidized?)

WU & McKINNEY

The improved strength for the oxidized and refinished specimens

as compared to those only machined can be shown to be significant at the 1%

level according to the Student t test. Two additional observations not

mentioned in the paper should aid in supporting this. First, the appearance of the specimen surfaces after the same machining procedure is

different for the two groups, i.e. the as machined specimens looked coarser

than the re-finished specimens. Second, preliminary microhardness measurements showed higher hardness for the as machined group. This leads to the

belief that an appreciable change may have taken place in the hulk during the

oxidation treatment.

As to the difference between 180 grit and 600 grit as finished specimens,

this also is significant at the 1% level. Reasons for this are explained

in the paper.

481
!j1iijiI
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SURFACE CONDITION AND STRENGTH OF SILICON CARBIDE - SILICON MATERIALS

T. J. Whalen and L. L. Terner

Research Staff, Ford Motor Company

Dearborn, Michigan 48121

The measured strength of SiC-Si materials with as-sil icided , ground, and

ground and polished surfaces was determined. The influence on strength of Si

and nickel coatings and the geometry of test specimens was studied.

Key Words: SiC, Surfaces, Coatings, Strength

1 . Introduction

A process for molding and sintering silicon carbide materials has been developed and

details are given in the literature [1,2]. Briefly, the process involves the formation of

a SiC-filled, thermosetting plastic which can be injection or transfer molded into complex

shapes. Subsequent steps include carbonizing the plastic at high temperature in an inert

atmosphere to yield a carbon-SiC composite and siliciding the composite at high temperature

in vacuum or controlled atmosphere to form a fully dense SiC-Si material. The final body

has essentially the same dimensions as the initial molded shape within a tolerance of 1%.

Silicon carbide surfaces of high strength are difficult to form and maintain, and for

this reason a study of the influence of coatings of chemically deposited materials such as nickel and SiC was initiated.

The purpose of this paper is to discuss the measured strength of SiC-Si composites as influenced by the surface condition, surface composition and geometry of the test specimens

Sample bars 1/4" x 1/8" x 1 1/4" or 1/4" diameter x 1 1/4" were tested for flexural

strength in 4-point bending upper span 3/8 in. (0.95 cm), lower span 3/4 in. (1.9 cm) at room temperature. Strengths are reported as "characteristic strength" which is defined as that strength at which the probability of failure is 62.3% of the tested population. Cross

head speed in the Instron tester was 0.05 in/min. The bars, were injection molded in preheated test bar molds at 155°C, except those bars for which thicknesses were varied and

these were compression molded in a simple uniaxial preheated die.

2. Testing Procedure

Table I. Surface Conditions of Test Samples.

Surface Range of Surface Roughness (yin. )

As Si 1 icided 30 - 90

Ground 9 14

Pol ished 2 4

SiC Coated 28 57

483
3. Surface Conditions

Different Surfaces were generated by diamond grinding, polishing, hipping and coating

with nickel or SiC. As-silicided surfaces were also evaluated. The maximum and minimum

surface average roughness for typical samples as measured by a Gould Surfanalyzer #150, is

shown in table I.

4. Results and Discussion

The influence of surface condition on the measured strength of a-SiC test bars of rec- tangular cross section is shown in table II. The strength of as-silicided surfaces varies over a wide range, depending on the precise details of the siliciding operation. The

presence of excess silicon on the surface makes transverse rupture testing a poor choice for

evaluating the material since the maximum stress is generated at the test bar surface and

silicon is brittle at room temperature and initiates fractures at low stresses. Grinding

and polishing yield more reproducible surfaces and the values of strength tend to reflect

the quality of these surface operations and of the SiC material itself. Much of our work

has been done with hand-ground and polished samples with rectangular cross sections but

good-quality surfaces free of edge chips were difficult to form with such brittle materials.

Table II. Surface Condition and Strength of a-SiC Rectangular Samples.

Surface Characteristic Strength

Kpsi Mpa

As Silicided 20-40 138-276

Ground and Polished 24 165

Nickel Plated 36 250

Round rods as samples were then investigated in an attempt to improve and simplify the

flow of material in the mold during forming to reduce molding flaws. It was also anticipated that improved strengths could be attained by increasing the perfection of surfaces by

machine grinding and polishing. The results shown in table III, compared with those listed

in table II, indicate that the strengths increased significantly for the types of surfaces

generated. Individual specimens with strengths in excess of 100 Kpsi were measured with

both a-SiC and 6-SiC compositions. It appears that the SiC-Si system has the potential to

form high strength materials. The open issue that remains is how reproducibly and consistently can the high-strength SiC be fabricated and how well can the surfaces be protected

against the formation of damaging flaws.

Table III. Surface Condition and Strength of a-SiC Round Samples.

Surface Characteristic Strength

Kpsi Mpa

As Silicided 49 340

Ground 61 420

Ground and Polished 75 520

484
Coating of the surface for strength improvement was attempted with a nickel coating

and a Si C coating. Nickel was deposited el ectrolytical ly, and one can see from the data

given in Table II that the strength increased significantly over the baseline data generated

for comparable, ground and polished samples. For low and moderate temperature applications,

a simple nickel coating may suffice to strengthen and perhaps protect these materials.

A dense Si C coating prepared by Chemetal Corporation with their Controlled Nucleation

Thermochemical Deposition (CNTD) process was evaluated, and the strength data are listed in

table IV. Although a significant increase in strength of coated samples was not realized,

the coating appeared adherent and dense, as seen in the microstructures of figure 1.

An individual sample which underwent CNTD Si C coating had a porous surface layer beneath the coating which led to some strength degradation. (See figure 2). Heating of the

SiC-Si specimens in a vacuum or inert atmosphere at high temperature prior to coating may

have removed the Si phase near the surface which invariably leads to a strength reduction

when tested in a flexure test.

Figure 1. Microstructure of CNTD Si C Coating Figure 2. Microstructure of Pourous Surface

on SiC-Si material . Layer Beneath CNTD SiC Coating

Table IV. Influence of SiC (CNTD) Coating on Strength.

Test Bars Characteristic Strength

Kpsi Mpa

Round Cross Section

Baseline - Ground and Polished 75 520

CNTD Coating 77 530

Rectangular Cross Section

Baseline - Ground and Polished 32 220

CNTD Coating 35 240

The test bar geometry was varied to study the susceptibility of strength to area or

volume under stress. Figure 3 shows the usual relationship observed for brittle materials

as the test bar geometry changes. In figure 3 one sees that the strength of rectangular

cross section specimens decreases as the thickness of the specimens increases. Preliminary

data on stronger, round bars indicates that the effect of change in geometry has a less severe effect on strength.

485
30

RECTANGULAR CROSS SECTION 4 - POINT LOADING

0:250 INCHES WIDTH

X GROUND

• POLISHED

20

o

z

tsJ

ttl

10

- 200

Figure 3. Influence of Sample Thickness on Strength of SiC-Si material.

100

0.125 0.188 0.250

THICKNESS (INCHES)

As a first approximation, the quality of a ground and polished surface of a brittle

material is proportional to the particle size of the brittle phase. When applied to SiC-Si

composites, this approximation suggests that the strength, which is influenced by surface

flaws, should decrease with increasing Si C particle size. Some data generated to investigate this assumption are given in figure 4. Each point on the graph represents 16 to 18

samples, and one can see that the strength peaks at a particle size about 16 ym. Samples

of coarser particle size decrease in strength probably due to the increased flaw size generated on the surface by grinding and polishing of a phase (SiC) which tends to fracture by

cleavage during these operations. Samples finer in particle size are difficult to si 1 i ci de

completely and, therefore, contain a pore phase which may behave as a stress concentrating

flaw. The mi crostructures of the 8 ym and the 36 ym SiC materials are shown in figure 5.

Exploring ways to si 1 i ci de composites with finer SiC particles may be a fruitful avenue for

future research.

AVERAGE PARTICLE SIZE { flm) 36 jX

Figure 4. Strength as a function of Initial Figure 5. Microstructure of SiC-Si Samples.

a-SiC Particle Size (roi'nd samples).

486
Another process which can be brought to bear on altering the surface, and the interior

of SiC-Si composites for improved mechanical properties is hot isostatic pressing (HIP).

Examples of property improvement (strength and consistency) of brittle materials by HIP

especially for WC cutting tools are well documented in the literature [3]. This process

improves density by eliminating the pore phase. SiC-Si samples were subjected to HIP

(30,000 psi argon at 1350°C for one hour) and the results are given in figure 6. The

strength of the initial a-SiC and B-SiC were significantly increased by the HIP process,

and further polishing of the samples increased the strength only slightly. Nickel-plated

a-SiC of 22 ym particle size was not altered by HIP.

60

2 50

XI 40

ui

DC»-

atu

in£

ui

*?30-

20 -

I 0 -

X BASELINE STRENGTH

©AFTER HIP I350*C • AFTER HIP a SURFACE

POLISH

- 400

300

- 200

- 100

I

t-o

U3

60

50

40

S 30

</> E

uiH

O

1

<E<

xu

20

10

a-sic a-sic a-sic

22 flm 22/i.m lOfJLm

Ni PLATE

Figure 6. Strength of SiC-Si Samples After

HIP at 1350°C for 1 hour.

X BASELINE STRENGTH

O AFTER HIP I450*C 400

X 300

X

200

o o

100

o

a.

a-sic

22/Xm

/3-sic

\Oflm

Figure 7. Strength of SiC-Si Samples

After HIP at 1450°C for 1 hour.

Since significant improvement in strength was observed with the HIP process at 1350°C,

which is below the melting point of silicon, it was logical to explore the effects of the

process above the melting point of silicon, such as 1450°C. When this war done, the results

are those given in figure 7. In both the a-SiC and in B-SiC, the strength decreased to

nearly half of the original, baseline strength of these materials. This drastic decrease in

strength could be understood when the microstructures of these treated materials were

observed. The microstructure, shown in figure 8, clearly shows that the HIP process at

1450°C evaporated silicon from the surface of the specimens and produced a weak, porous

layer at the surface.

487
Concl usions

1. SiC-Si Composites have the potential to attain strengths greater than 100 Kpsi, when used1 at room temperature.

2. SiC-Si materials with ground and polished surfaces yield greater strength when tested in

flexure than do as-silicided materials.

3. Round test bars are easier to prepare and they give higher strengths than bars with

rectangular cross sections.

4. HIP processing shows promise in increasing the strength of SiC-Si composites providing

the migration of the Si phase near the surface can be controlled.

References

[1] Willermet, P. A., Pett, R.A., and Whalen, T.J., Development and Processing of InjectionMoldable Reaction-Sintered Si C Compositions, Ceramic Bulletin 5J7_ [8] 744-747 (August

1978).

[2] Whalen, T.J., Noakes, J.E., and Terner, L.L., Progress on Injection-Molded, ReactionBonded Si C (Proc. 5th Army Materials Technology Conf. Newport, Rhode Island, March 21-

25, 1977). Book, Ceramics for High Performance Applications - II J. L. Burke, E. N. Lenoe and R. N. Katz, eds., pp. 179-189 (Brook Hill Publishing Company, Chestnut Hill,

Mass., 1978)

[3] Van Cleave, D.A., Hipping Improves Properties in Castings and PM Parts, Iron Age,

August 23, 1976.

Discussion

RICE

The negligible effect of SiC coatings is not surprising since this is

not expected to fill in or bond cracks together. However, you Ni coating

did increase strength. The question is whether this really "heals" cracks

or alternatively alters, e.g. blunt cracks. Have you investigated this,

e.g. by removing the Ni layer by polishing or etching?

WHALEN

We have polished the interfacial area to microscopically evaluate the

continuity of the Ni coating. The adherence is good and it follows well

the contours of the substrate. We do not know if the coating process alters

the flaw geometry on a micro scale, but we are attempting to find out.

488
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

THE EFFECT OF GRINDING AND POLISHING ON THE

STRENGTH OF ANNEALED SODA-LIME-SILICA GLASS

R. Bowles

British Glass Industry Research Association

Sheffield SI 0 2UA England

As part of the continuing studies sponspored by the Procurement

Executive, Ministry of Defence (UK), recent work at British Glass

Industry Research Association has aimed at assessing the defects

generated in glass during its working under conventional grinding

and polishing procedures.

The assessment has been mainly based upon mechanical strength

property measurements of lath specimens typically 150 x 25 x 3.5 mm

either at specific stages of processing or after a controlled

sequence of operations designed to produce a state of minimum

surface damage.

The strength of laths was determined under four point flexure

test, at a nominal stress rate of 5 MPa/s under normal ambient

conditions of temperature and humidity.

The resulting fracture stress data was calculated on simple

bending theory after bonded strain gauge tests indicated its adequacy

under observed deflections to fracture. Group statistics were largely

calculated assuming normal distribution theory but tests on aggregated

groups of specimens were also examined by Weibull distribution theory

to determine the probability of fracture under relatively lower

stresses.

Results show that by application of prior or inter-stage proof

loading, that although a low fraction of weak specimens still emerges,

a significant improvement of strength can be achieved in annealed

glass specimens of reasonable size.

Keywords: Defects; etching; glass; grinding; polishing; strength.

1 . Literature Review

A review was made of some publications on polishing processes, methods of

examining polished surfaces and of the effects of polishing on the strength of

glass. A list of references is included herewith.

The mechanisms of grinding and polishing processes (1-12) in glass were comprehensively reviewed by Cornish (6) and he covered the period 1665-1961.

Surface examination techniques by electron microscope are numerous,

references (1,2,4,6,7,9) illustrate the potentialities. Holland (13) considered

optical microscopy using phase contrast and schlieren methods but concluded that

it was not possible to differentiate between optical path differences due to refractive

index gradient or topographical features. Holland (11) also reviewed the mechanisms

of grinding and polishing in which surface flow is sometimes said to occur.

489
Levengood (3) showed that narrow deep cracks could be bridged at the surface by

polishing glass with rouge on felt, and therefore the apparent absence of surface defects

does not exclude their presence. Rawstron (4) illustrated the

0

difference between

surface polishing under high pressure and a heated slurry at 60 C that under low pressure

and a cooler slurry at 15 C. From a microscope examination, it was concluded that

surface flow predominated under the former condition whereas the slower, latter process

produced general surface removal. Rawstron describes the use of a Talysurf instrument

to measure the thickness of a polished layer ~100 nm deep on glass. Holland (13)

described multiple beam i nterferometry to measure surface topography.

Ernsberger (14) used an ion-exchange process to examine the density of microcracks,

and Levengood (13) used water vapour condensation to reveal microscopic scratches.

Stoll (8) measured the strength of fused silica discs 102 mm diameter 6.2 mm thick

comparing a commercial surface preparation with a controlled grinding and polishing

schedule set to remove at each stage, those defects remaining from previous work stages.

Defects were estimated to be approximately 3 x abrasive particle diameters. The discs were tested under concentric rings and a strain gauge calibration had shown close agreement

between measured stress and that calculated by an overhang corrected formula by Roark (15).

The controlled schedule gave a 50% strength improvement over the commercial process,

yielding 88 MPa compared with 59 MPa. Polished samples with light or severe scratches were found to be of similar strength ~48 MPa regardless of grinding history.

Shand (16) reported on a strength comparison using glass plates processed by one of five participant companies but he could not draw firm conclusions as to differences

between processes because of large inherent variances within groups, typically

ranging 40-110 MPa.

McKinney (12) tested fine grain alumina discs in different surface conditions,

and measured rupture moduli by ball and ring test calibrated against strain gauge measurements. The strength was found to increase from 150 MPa to 250 MPa when the SiC grit

used for surface grinding was reduced from 210 ym to 17 ynu

Mecholsky (17) tested several glass compositions, either as laths under three

point bending or as discs to examine the effect of grinding in orthogonal directions

of sample relative to a diamond grinding wheel, and found two differing flaw populations.

He measured rupture moduli ~122 MPa after polishing and ~112 MPa for samples etched

and containing known flaws in different orientations. He concluded defects in polished

glass were ~23 ym deep. In an earlier paper, Mecholsky (18) examined the effects of

sputtering on the strength of silicate glasses, and found that although a significant

improvement obtained, sputtering introduced its own limiting constraint after removal

of the original defect layer.

Shelyubskii (19) examined artificial defects produced in sheet glass by abrasion

with electro-corundum grit sizes 1 to 1000 ym. He found a low dependence of minimum

strength on grit size, and quoted a relation between defect layer depth (F) and grit

size (D) as F = 1.1D. He differentiated between strengths resulting from grinding

on a soft as opposed to a hard grinding table. He also commented on strength - sample

size effect.

Orlova (20) reported on different glass compositions using discs, 30 mm diameter

by 1.5 mm thick, which were ground through successive emery stages and then polished.

Silicate glass in annealed condition was found to have a rupture modulus of 118 MPa.

Densienko (21) examined HF etched and chemically strengthened sheet glass using

44 mm square samples by 2 mm thick which were tested under central symmetric bending

at 50 MPa/s. He observed an optimum etch removal of ~25 ym before ion-exchange and

a subsequent light etch of ~2 ym to give a final strength ~700 MPa. Allowing for

a quoted surface compression of ~ 300 MPa induced by the ion-exchange, the annealed

glass strength could be construed to amount to the difference, i„e. 400 MPa for a commercial quality sheet glass after a surface etch of ~27 ym.

The question of interpretation of strength data as it relates to components of

490
different sizes, and under different stresses in service has been considered by several

authors (22,23,24,25). Fracture mechanics theories have been used to calculate the

probability of component lifetime from knowledge of original surface defecto Lenoe (23)

advised caution in assumption of Weibull distributions and pointed out the possible

variance of crack growth phenomena. Davidge (24) discussed strength-probability-time

relationships for ceramic materials, and evaluated effective times to failure,, Ritter

(25) calculated the necessary proof stresses required to predict the probability of

lifetime under particular stress, and in a later publication (26) he referred to the

advantage of proof test in defining a threshold flaw size, whereas calculation of design

data from typical stress distributions often yield low strength at low failure

probabi 1 i ties.

2. BGIRA experimental programme

The strength of annealed soda-lime-silica glass subjected to different surface

finishing procedures, was measured under ambient conditions by flexure tests of lath

and disc samples. Attempts were made to relate rupture moduli with corresponding

surface finishes produced by diamond tool grinding, smoothing by free Si C grit, by

ceria polishing and acid etching. Precautions were taken to avoid extraneous damage

of the test surface during processing to try to obviate non-characteristic defects.

Experiments were undertaken to assess the depth of surface flaws. The strength data

distributions were assumed to be normal for purposes of comparison, and the results

obtained from this programme were used to devise a process schedule so as to leave

a minimum of defects on an adequately finished surface.

This schedule was tested on a group of lath samples to ascertain its effectiveness,

and the consequent strength data were analysed by Weibull distribution theory.

3. Sample preparation

Two forms of sample were used, discs nominally 100 mm diameter by 5 mm thick,

and laths nominally 150 x 25 x 3 mm. They were prepared from either white plate glass

supplied as rough cast, 8 mm annealed strip, or from 4 mm float glass, both products

of Pi Iking ton Brothers Ltd.

The greater bulk of surface removal was necessary in the initial roughing stage

for the as-cast glass by 8" diameter diamond cup wheel, grade 125-150 pm grit size,

metal bonded. Both surfaces of the glass were ground substantially to remove surface

undulations and tonear-fi nish thickness. The discs were cut by diamond impregnated,

thin-walled drill operated from both surfaces so as to reduce edge chipping. The roughing

stage was not required for the 4 mm float glass stock which was basically worked on one surface only.

Both types of glass were smoothed by aqueous slurries using Si C grits of various

grades, and polished by ceria slurry on felt pads.

Different edge finishes were investigated to seek to minimise edge failures which

would otherwise introduce data non-characteristic of surface properties. Lath specimens

required preferential edge treatments to reduce edge failure incidence.

Contact with the critical area of the test surface by measuring instruments was

avoided, and progress through the various stages was monitored via known surface removal

rates established on coupons.

The actual surface removal was measured between stages of working on the non-critical

area of each sample, i.e. that for laths being outside the span of the four point test

rig. In some instances, measurements were possible after fracture but not in the case

of strong specimens which produced finely fragmented debris which also obscured the

fracture origin.

Discs were processed in batches of five, being wax mounted on a block, and using

491
a controlled addition of SiC grit slurry against a steel lap. The discs were polished

singly on an "Autosphere" polishing machine with a 150 mm felt pad and using a recirculating aqueous ceria slurry, they required usually 4 minutes but longer periods were

set on occasion.

Laths were usually prepared by scoring the stock with a wheel cutter, using the

scored side as the eventual compression test surface. The edges were worked in blocks

of six laths to reduce them to a common width and the edges were arrissed by SiC grit,

treating laths individually. After a selective edge etch applied whilst the main surfaces

were protected by wax coating, the test surface was stripped and machined again in

batches of six samples through smoothing and polishing stages. Working of laths in

groups was preferred to achieve consistent process conditions.

For discs, typical surface removal rates measured in simultaneous grinding of

five samples were as follows:-

SiC grit range Mean Average grinding rate

Grade (vim) (vim) (ym/hour)

320 18 - 46 32 342

400 14 - 34 24 284

600 8 - 25 16 135

Etching of samples was usually done by intermittent dip of groups of samples

supported at their ends, in HF solution and then in water baths, thermostatically

controlled,, The problem of avoiding contact with test surfaces was most acute during

this stage, and the process was again monitored by coupon specimens undergoing

simultaneous treatment. The use of coupons also served to allow for diminishing etch

rates with use of acid.

4. Test configurations

For disc samples, uniform pressure and concentric ring tests were applied and stresses

were calculated according to formulae quoted in the appendix. In early studies of

110 mm diameter, 2 mm strengthened glass discs loaded under uniform pressure, bonded

strain gauges showed that elementary theory (equation 3) in the appendix, did not apply.

An empirical equation was derived and was of similar form to that allowing for the

development of membrane stress component in addition to flexure stress but having different

coefficients to the theoretical free edge condition. The stress analysis also indicated

the differences between radial and tangential stress components off centre.

From this study it was concluded that in certain cases, allowance should be made not

only for fracture origin location, but also for original crack propagation direction

when evaluating an appropriate fracture stress. However in the current tests of the

weaker, annealed glass discs, much lower values of dimensi onless stress (a*) resulted

and therefore the elementary formula (equation 3) was deemed to be sufficiently accurate.

Concentric ring tests were applied to reduce the preponderance of edge failures

experienced in the uniform pressure tests, and initially, different ring diameters were

used to seek to minimise that tendency. Using concentric rings of 68 mm and 50 mm diameters, the incidence of edge failures decreased but there appeared to be a concentration of fracture origins near the inner ring radius which suggested an undesirable

stress concentration in that region. Using elementary theory, (equation 7), fracture

stresses were calculated and were compared with values obtained in uniform pressure

tests. Although some degree of consistency was obtained, variance of test data was so large as to obscure potential differences between various groups of samples.

Tests on lath specimens were accordingly started with promising results in respect

of statistical variance. Because these samples were easier to make, the subsequent

492
test programme was based upon such specimens which were tested under four point bending

by a Monsanto-Tensometer 'E' test machine. Adjustment of the cross-head speed allowed

the setting of a nominal stress rate of 5 MPa/s. This rate was calculated on assumption

of elementary theory but it was found in other tests of chemically strengthened,

2mm thick glass laths, that a more comprehensive analysis may be required. For example,

stress data as calculated from simple formula and as computed by numerical solution

as per Ritter and Vrooman (27) was compared with data measured by bonded strain gauge

up to fracture. A graph of stress versus load showed that non-linear stress-load

relationships might become significant for stronger specimens withstanding greater

deflections. Again however for annealed glass laths, elementary formulae were deemed

adequate and a comparison between rupture moduli for discs under concentric ring tests,

with laths under four point bend test gave agreement as illustrated.

It proved convenient to adopt a modified specimen preparation technique in which

laths were given a pre-etch before applying surface machining stages. The purpose

of this was to ensure the removal of significant defects before applying the final

stage whose characteristic strength was to be evaluated. After pre-etch, samples were proof-loaded at 400 MPa for one minute, and survivors, usually about 90% of the original

test batch, were then re-surfaced by different processes.

5. Assessment of surface defects

The defective region in a diamond ground glass surface was initially studied by

forming a wedge shaped specimen in which ca 300 ym had been removed from one end whereas

the other end was virtually untouched. The wedge was formed by using 600 SiC grit slurry,

and this ground surface was polished by ceria. The specimen was then ion-exchanged

in a UNO3-KNO3 salt bath at 200 C for 2 hours, and the resulting crack pattern developed

by HF dip for 2 minutes. Cracks were observed even at a depth of 300 ym and this

observation was checked by repolishing the wedge surface to a further depth of 30 ym,

and then subjecting the specimen to a second ion-exchange treatment. This had the

effect of suppressing cracks except at the thick end of the wedge and down to a point

corresponding to a total surface removal of ca 90 ym. This apparent discrepancy

was attributed to the presence of defects introduced by the 600 SiC grit and not fully

polished out before the first ion-exchange treatment. The second polishing stage

eliminated much of this damage, and it was concluded that the defect layer due to diamond

wheel grinding was approximately 90 ym deep.

Defective layer depths were also studied in etched and proof loaded laths which

had been ground with particular grits. Groups of similarly processed laths were re-etched to different degrees and the corresponding rupture moduli measured. Pavelcheck

(28) found that crack tips were relatively unaffected in early stages of HF etching

of soda-lime glass, and he quoted a relationship between fracture stress (a 0 ) for an original crack length (CQ

) and fracture stress (a) after etching a layer (5) from the

surface

This can be re-written as

(°o/a) = 1 - 6/C 0

(2)

2 and a plot of (°o/a) against 6 should asymptotically reach zero when 6 approximates

to C 0 . Appropriate values for C0 for different abrasive or polishing powders were evaluated for process conditions applying to this study. Those values were taken to

represent minimum surface removal depths necessary to clear pre-existing flaws, and

a schedule for progressive surfacing was evolved therefrom. The values ranged from

2 ym for a ceria polished surface ("Optipol 8G") to 22 ym for "Carborundum F400" SiC

grit.

493
A group of 71 samples were processed from the etched, proof-loaded stage through

400, 600 Si C grits and ceria polish to a final light surface etch. After fracture,

the fragmentation was almost always too fine to allow identification of fracture origin

and it was not possible to distinguish between surface and edge fracture in all but

2 samples which failed unexpectedly at 117 and 148 MPa from the surface. The remaining

69 samples ranged from 464 MPa to 1310 MPa, having a mean of 976 MPa and a standard

deviation of 185, i.e. a coefficient of variation of 0.19. These data are shown as

a Wei bull distribution and a comparison was made against a line of similar slope but

equivalent to a characteristic time (t) of 60s, i.e. the proof load duration. From

that it appeared as if 4% of the population had failed at a stress equivalent to or

less than proof test. It appeared therefore that some damage had been sustained in this

fraction of the population and to this extent the controlled schedule had not resulted

in the complete determination of a minimum strength level.

However whilst it is not known what process contributed defects to this weaker

fraction of samples, it is clear that a final proof test could be used in practical

components to establish a maximum defect intensity.

Acknowledgements

This work has been carried out with the support of the Procurement Executive,

Ministry of Defence in the UK. The author wishes to thank the sponsor and the

Council of the British Glass Industry Research Association for permission to publish

this paper.

6. References

(1) Bruche, E and Poppa, H. , The polishing of glass. Part 1 - Optical polish,

Glastech. Ber., 28 (6) 232-242, (1955)

ii (2) Bruche, E. and Poppa, H., The polishing of glass. Part II - Technical 1 polishing of plate glass, Glastech. Ber., 29,(5), 183-191, (1956)

(3) Levengood, W.C. and Fowler, W.E., Morphology of fractures in polished

glass surfaces, J. Amer. Ceram. Soc. 40_, (1 ), 31-34, (1957)

(4) Rawstron, G.O., The nature of polished glass surfaces. J. Soc. Glass Technol

.

42, 253T-260T, (1958)

(5) Bruche, E. et al . The polishing of glass (discussion). Silikattechnik, 10,

(4), 213, (1959) ~

(6) Cornish, D.C. The mechanism of glass polishing - a history and bibil iography.

BSIRA, Res. Rept. No. R 267,(1961 )

(7) Bound, U.M., Cornish, D.C. et al . The mechanism of glass polishing - II The

effects of slurry concentration and temperature. BSIRA, Res. Rept. No. R 284, (1962)

(8) Stoll, R., Forman, P.F., Edelman, J. The effect of different grinding procedures

on the strength of scratched and unscratched fused silica. Symp. Res. Mec. du Verre,

paper 1-5, 377-392, (1962)

(9) Cornish, D.C. S Verrall , A.F.D., Watt, I.M. The mechanism of glass polishing,

presented at Conference on the Physics of Optical Glass, 3rd-4th October 1963

Pilkington Bros. Ltd., Research Labs., Ormskirk, Lancashire.

(10) Holland, L. The properties of glass surfaces. 15-73, (Chapman and Hall, 1964)

(11) Holland, L. The properties of glass surfaces. 53-56, (Chapman and Hall, 1964)

(12) Mc'Kinney, K.R. and Herbert, CM. The effect of surface finish on structural

ceramic failure. J. Amer. Ceram. Soc. 53, (9), 513-516, (1970)

494
(13) Holland, L. The properties of glass surfaces, 107-120, (Chapman and Hall, 1964)

(14) Ernsberger, F.M. A study of the origin and frequency of occurrence of Griffith's

micro-cracks on glass surfaces. Advances in Glass Technology, Paper VI, 39,

511-524, (Plenum Press, New York 1962)

(15) Roark, R.J. Formulas for stress and strain, p. 195, (McGraw-Hill, 1954)

(16) Shand, E.B. Weakening defects in glass. Amer. Ceram. Soc. Bull. 46, (11),

1111-1115, (1967) ~

(17) Mecholsky, J.J., Freiman, S.W., Rice, R.W., The effects of surface finish on

the strength and fracture of glass. 1977, Proc. XI ICG Conf. (11), 489-497,

(Prague 1977)

(18) Mecholsky, J.J., Rice, R.W. The effect of sputtering on the strength of

silicate glass. J. Appld. Phys. 47, (7), 2972-2974, (1976)

(19) Shelyubskii, V.I., Velikovskaya, N.A., Gutner, S.Kh. Effect of artificially

applied surface defects on the strength of sheet glass. Steklo i Keram. 2£,

(11), 7-9, (1969). Glass and Ceramics 26, (11), 639-642, (1969)

(20) Orlova, D.A., Kuznetsov, A. Ya., Mechanical strength of components made of

various optical materials. Optikomekhanicheskaya promyshlennost (6), 79, (1975)

(21) Densienko, O.N., Shchlegova, O.V., Soboliev, Ye. V. The effect of etching in HF on the strengthening of glass by means of low temperature ion-exchange 1976,

Steklo i Keram. 33, (2), 13-14, (1976) Glass and Ceramics 33, (2), 92-93, (1976)

(22) Wiederhorn, S.M., Strength of glass - fracture mechanics approach.

Proc. X ICG Conf. (11), 1-13, (Kyoto 1974)

(23) Lenoe, E.M., Neal , D.M. Assessment of strength-probability-time relationships

in ceramics. AMMRC TR 75-13, pp. 22 (1975)

(24) Davidge, R.W., McLaren, J.R., Tappin, G. Strength-probability-time relationships

in ceramics. J. Mat. Sc., 8, 1699-1705, (1973)

(25) Ritter, J.E., Laporte, R.P. Effect of test environment on stress-corrosion

susceptibility of glass. J. Am. Ceram. Soc. 58, (7-8), 265-272, (1975)

(26) Ritter, J.E., Meisel, J. A. Strength of failure predictions for glass and ceramics

J. Am. Ceram. Soc. 59, (11-12), 478-481, (1976)

(27) Ritter, J.E., Vrooman, D.L., Rep. Themis-UM-69-4, Contract 0WR-N00014-A-0146-3

,

(1969)

(28) Pavelchek, E.K., Doremus, R.H. Fracture strength of soda-lime glass after etching

J. Mat. Sc. 9, 1803-1808, (1974)

495
Appendix 1. Formulae for disc samples

For discs under uniform pressure:

3PA 2 (3 + v)

°

= 8t^

3PA (1 - v) (5 + v)

T6Et 3

aA2

Et 7

1 .238

(y/ t

) + 0.294 (y/ t 2)

p* PA4 it 1* 1.016

(y/ t

) + 0.376 (y/ t

):

For discs under concentric ring loading:

3W

a

t

= 2^ (1 - v)

(

R2 -/2 )

, H + v ) l0 9e(

where a - stress; subsets r, t for radial, tangential components

P - pressure applied A - aperture radius

E - Youngs modulus \> - Poisson ratio

a - sample radius t - thickness

y

- sample deflection

R,r - radii of outer, inner concentric rings

W - force applied

496
Appendix 2. Bend test rig and formulae for lath samples

Where M = F£, I = bt 3 /12, 2F W

Maximum deflection (u + v) = M (3L 2 - 4«, 2

) 24EI

_ M (L - 2£)2

u 8EI

Mt max stress a = >

v =

_Mj, (3L - 4a)

6EI

max strain e Mt

2TI

substituting experimental parameters

L = 127 mm, £ = 44.5 mm, E = 73 GPa

thickness t = 3.5 mm, v = 3.5 mm/minute

stress rate a 3Etv

£(3L-4£)

strain rate e = 68 x 10" 6 /s

= 4.95 MPa/s

497
Appendix 3. Figures

Figure 1. Photomi crograohs by scanning electron microscope

of a glass surface after diamond grinding (a) and

320 Si C smoothing (b) at 2400 x magnification.

498
Figure 2. Crack patterns after first ion-exchange of a glass specimen,

wedge ground and polished after diamond grinding (a) at 60 pm from

original surface, (b) at 210 pm and (c) at about 300 ym depth.

499
Figure 3. Crack patterns in wedged specimen as in figure 2 but after

re-polishing and a second ion-exchange treatment; (a) at 20 ym

from original surface, (b) at 40 ym, (c) at 65 ym and (d) at 90 ym

depth

.

500
Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SURFACE FLAW EFFECTS ON CRACK PROPAGATION FROM SUBSONIC WATER DROP IMPACT 1

Y. Marvin Ito and Martin Rosenblatt

California Research & Technology, Inc.

Woodland Hills, Ca. 91367

The effects of surface flaws on ceramic target damage from subsonic water

drop impacts are investigated using numerical code simulations based on fundamental wave propagation and fracture mechanics concepts. The water drop impact

generates a tensile stress wave with a steep near-surface gradient. As a con- sequence of this steep stress gradient, crack activation is dependent on the

depth of surface flaws. In addition, a relatively large flaw can perturb the

stress field such that crack activation in its vicinity is suppressed.

Key Words: Crack activation; numerical simulation; subsonic water drop impact;

surface flaws.

1. Introduction

The damage of a ceramic target from a subsonic water drop impact is dependent on the

tensile stresses generated by the impact and the criteria for the extension of flaws ex- posed to the tensile stress field. Since the target response is essentially elastic, the

maximum tensile stresses are the radial stress components. Also, the tensile stress field

has a steep near-surface gradient. Thus, these radial stresses tend to produce circumferential cracks which initiate at the surface and propagate into the target [lj 2

.

As a consequence of the relatively steep near-surface stress gradient, it is expected

[l\ that crack activation is dependent on the depth of surface flaws (exposed to this

stress gradient). It is the intent of this paper to analyze this effect for specified impact conditions.

2. Approach

A series of numerical simulations are performed in order to examine the effects of

surface flaws on zinc setenide window damage from subsonic water drop impact. The normal

impact conditions of a 2Rq diameter water drop onto the window target at velocity VQ is

shown in figure 1. A finite-difference computer code, WAVE-L, with an analytic pressure

loading model is used to predict the target response.

WAVE-L is a two-dimensional (axial or plane symmetry), Lagrangian, explicit, finitedifference code based on the HEMP scheme \l\ which integrates the governing partial differential equations of motion for arbitrary dynamic problems in solid and fluid mechanics.

It has been applied in a number of particle impact investigations, including both subsonic [3\ and hypersonic [4] erosion.

The analytic pressure loading model has been developed for simulating subsonic water

drop impacts and is discussed elsewhere ]j>] . The impact model specifies the loading pressure on the target surface as a function of radius from the impact point and time, P(r,t).

"Research supported by Office of Naval Research under Contract No. N00014-77-C-0790.

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

501
Axis of Symmetry

Figure 1. Normal impact conditions of 2RQ water drop onto zinc selenide (ZnSe) window

at velocity V .

The development of this analytic loading model was based on detailed WAVE-L numerical calculations [5~\ of spherical water drop impact onto a rigid surface. The rigid surface

approximation is valid when the target surface deflection is negligible compared to the

drop radius. For elastic materials, this condition is satisfied if the target material

impedance (density times wave speed) is much greater than the water impedance. The zinc

selenide targets described in this paper have an impedance which is 14 times greater than

water. Hence, the decoupZ&d approach is justified in the present case.

The loading pressure function, P(r,t), for a water drop of 2-mm diameter (2RD ) impact

ing at 222 m/ s (730 fps) velocity (VD ) onto a zinc selenide (ZnSe) window target has been

determined Jjj] . Figure 2 gives the maximum impact pressure profile. Due to the spherical

geometry and water material properties, an off-axis peak3 pressure occurs (at early times)

which is about twice the Hugoniot value. This peak pressure occurs when the velocity of

the radius of contact (r in fig. 1) approximately equals the speed of sound (C Q

) in

water .

This impact pressure P(r,t) is used as the loading boundary condition for the axisymmetric WAVE-L numerical calculation of the ZnSe target response. Figure 3 shows the

initial computational grid for the target material using 80 cells per water drop diameter

(2RQ = 2mm). In the WAVE-L code, particle velocity is defined at the corners of the com- putational cells while the stress state is associated with the center of the cells.

3 Note that this pressure peak will not occur for subsonic solid particle impacts,

and observed differences in target response from solid versus liquid projectile

impacts may be partially due to the presence of the off-axis pressure peak.

502
0 .02 .04 .06 .08 .10 .12 .14 .16

Radius, r (mm)

Figure 2. Maximum impact pressure as a function of radial distance from impact center

for 222-m/s impact of 2-mm water drop onto ZnSe window.

CALIFORNIA RESEARCH AHD TECHNOLOGY, INC.

0.02

a

0.04.

0.08

0.06 0.02 0 0.02 0.04 0.06

Radius, r (cm)

Figure 3. Initial computational grid in target material using cells per 2-mm water

drop diameter.

503
Figure 4 indicates typical elastic response of a ZnSe window in terms of the particle

velocity field and principal in-plane (rs z) tensile stress field at 0.1 us. The velocity

vectors show the magnitude and direction of the velocity of the particle mass located at

the tail of the vector. A 15-m/s (50-fps) scale bar is shown in the upper right corner

of the figure. The principal tensile stresses are indicated by lines which show their

magnitude and direction in the r-z plane. (Hoop tensile stresses are not indicated on

these plots.) A 1-bar (14500-psi) scale bar is shown in the upper left corner of these

plots. Thus, at 0.1 us after the impact, the peak velocities are roughly 15 m/s (50 fps)

in the target and the peak tensile stresses are about 2 kbar (30000 psi)

.

CALIFORNIA RESEARCH AND TECHNOLOGY, INC. EROSION CASE

CYCLE 50 I—I 1 KB 50 FT/ SEC

0.02 -

4-1 P.

a) a

0.04. -

0.06 -

0.06 0.04 0.02 0 0.02 0.04 0.06

Radius, r (cm)

Figure 4. Elastic response of ZnSe window for 222-m/s impact of 2-mm water drop at

0.1 ys (80 computational cells per drop diameter).

It is seen that a region of high tensile stresses occurs near the surface, outside

the contact area. The material directly under the contact area is in pure compression.

Figure 5 shows the corresponding velocity and in-plane tensile stress fields at 0.2 ys.

The qualitative features are similar to the 0.1-ys plot; however, the peak velocities and

peak tensile stresses have already decayed below the 0.1-ys values.

The specified impact conditions used in the present numerical simulations involve

222 m/s (730 fps) impact of 2-mm diameter water drop onto zinc selenide windows. These

conditions are considered to be representative of subsonic liquid drop impact onto ceramic

targets. In the following, Section 3 gives details of the near-surface stress gradient,

Section 4 shows the effect of surface flaws and Section 5 has the concluding discussion.

504
-.02

CALIFORNIA RESEARCH AND TECHNOLOGY,

EROSION CASE

CYCLE 98 I—I 1 KB

1

1

'

INC.

50 FT/SEC

Position of water drop

at .2 ys

l\ » ' I » * »

i n i n m / III1IIU

J » r «. f •• r x

1 I M H ( 11111/

w / / / n m >

0.02 -

/S\\S

f / A / V l V \»\, V^* *

<• 4 / J J t MN N.'N.-v-- ~ *

< r f | | | v \ VN.^^^ * 'JIJIIA^N N. -N. -*s ~» ~

I | \ \ \ V V V, —

a.

11 Q

» V \ \ \

* \ \ \ \ V — '

w V V V V N. \. >. —

o.ou -

o.oe

Magnitude and Orientation of Principal

In-Plane Tensile Stresses

I ^ \ 1

i- r i ill

till

f i f I J

[ I i 11/ (S

f l l V V VW v*»» UVW>v

1 \ \

V V \ \ vvv>v>»»"

l v I vvvvvyvvvvvvv»'

WVVVVVVVVVVVVVVVV'

% % I V t- | Particle Velocity Vectc Vi.VVWVVVVVVVVVv. .

0.06 0.04 0.02 0 0.02 0.04 0.06

Radius, r (cm)

Figure 5. Elastic response of ZnSe window for 222-m/s impact of 2-mm water drop at

0.2 us (80 computational cells per drop diameter).

3. Near-Surface Stress Gradient

The formation of high tensile stresses near the surface of a ZnSe target impacted by

a water drop is illustrated in figures 4 and 5 (Section 2). In this numerical simulation

the initial computational grid (fig. 3 in Section 2) has 40 cells per water drop radius.

Since the effective prediction of crack activation requires that these high tensile

stresses be well characterized, an appropriate fine-zoned computational grid must be

selected. Some calculations, employing selected fine-zoned computational grids in the

near-surface impact region, provide a basis for selecting the grid structure for subsequent crack activation experiments.

Figure 6 shows the peak tensile stress versus depth in a ZnSe window impacted by a water drop at 222 m/s (730 fps) for three variations in computational cell size. Note

the very steep stress gradient near the impact surface and its increased resolution1

* with

decreased cell size (where stress is defined at the cell center) . For a water drop of

1-mm radius5 the finest-zoned (80 cells per drop radius) case is adequate to resolve peak

tensile stress from 356 MPa (52000 psi) at a depth of about 6 um to 80 MPa (12000 psi) at

100 um depth.

In fact, an analytic solution by Blowers [6] predicts infinite peak tensile stress

at the impact surface for purely elastic targets.

5 Note that these elastic solutions can be scaled linearly in dimensions and time for

other drop radii as long as the impact velocity and material properties are not

changed.

505
Figure 6. ViUlk t&tU-iZe. &£A<Li>i> veAAUA dtpth (for all radii) in a ZnSe window impacted

by a water drop at 222 m/s for variations in computational cell size.

Figure 7 shows the peak tensile (radial) stress near the surface versus radius. The

results are from the surface computational cells which have stress defined at the cell

center depth of z = Rq/160 or 6.25 um. The maximum in the peak tension (356 MPa) occurs

at radius - 0.2 RQ or 200 um. For r < 0.1 R0 or 100 um no tension develops in the

target surface. At radius 2 rp = 0.4 RG or 400 um the peak tension has dropped to only

200 MPa (30000 psi) . The nominal static tensile strength of ZnSe is about 45 MPa

(6500 psi).

The temporal development of the radial tensile stress at four radial locations

(3/4 rp, rp, 3/2 rp and 2 rp) near the target surface is given in figure 8. The pulse

duration (which is about 30 ns at the maximum tension location rp) increases with radial

location (~60 ns at 3/2 rp and -75 ns at 2 rp

)

.

The effect of a surface flaw of depth Ro /80 or 12.5 ym at radius rp

= 0.2 Rq or 200 ym on the peak radial tensile stress is shown on figure 9. For this flaw, the stress

gradient ahead of the crack tip is similar to that from the surface without a flaw, but

the maximum in the peak radial tension is about 10 percent less.

It is concluded that the effects of surface flaws of depths from 10 ym and greater,

which are exposed to these stress gradients, should be well characterized in subsequent

crack activation experiments using the fine-zoned computational grid with 80 cells per

drop radius.

506
CALIFORNIA RESEARCH AND TECHNOLOGY, INC. ONR HIGH STRESS GRADIENT CASE

„ RADIAL TENSION TIME HISTORIES AT Z = .000825 cm for R„ = .

1

1

1 1 °—

1 cm

x

as i i.o h

CH

OS

0.6

0.6

<u o.U.

-H

o 0.2

z = 0.00625RQ

Axisof

Vo \

I Water

Drop

\

—-^Window

i 1 > 1 1 1\/7

3 i 4

rP

/ / 1 /

> b

* t

' )*

p

2i"p

r„a.l51L

1/80

0.05 0.10 0.15 0.20 0 o 2S

Time, ys (for R0 = 1 mm)

Figure 8. Normalized (crmax=356 MPa at r=r

p

) mcLial. tan&ilt i>£neJ>i> kldtOfviu near Surface

(z=A/2) for impact (VQ=222 m/s) of ZnSe window by a water drop (R0=lmm)

507
Scaled Depth from Surface, z/RQ

Figure 9. Effect of surface hoop crack at r=r on normalized (amax=356MPa) ptak hJX&LoJL

t<LWt>iZ<l t>tAZAt, veAAuA Acal&d dzptk at r=r

p

for impact (VD=222m/s) on ZnSe

window by a water drop (R0=lmm)

.

4. Surface Flaw Effects

In the numerical code a flaw is simulated as "tensile" cracks across computational

cells near the surface. In a cell with an initial crack, no tensile stress across the

crack is permitted, and no shear stress is permitted on the crack if the crack is open

(the width of the crack is continuously monitored). These stress adjustments are important because a crack frequently alters the local stress field in such a way as to enhance

its own growth. Thus, any realistic method of predicting crack growth must consider the

altered stress field.

The crack activation experiments involve surface flaws of various depths located at

radius 3/2 r p

= 0.3 Ro or 300 urn. At this location, without a flaw, the maximum radial

tensile stress is about 285 MPa and a tensile stress gradient develops to a depth of

z = 0.075 RQ or 75 ym (see fig. 10). Also, the tensile pulse duration near the surface is

about 60 ns (see fig. 8 in Section 3).

Figure 10 shows the effect of surface hoop (circumferential) flaws on the radial

tensile stress gradient ahead of the crack tip for crack depths of 25, 50, and 100 ym

(R0=lmm) . In each case, the peak radial tensile stress ahead (one-half a cell dimension)

of the crack tip is enhanced, but decreases with increasing flaw depth.

Figure 11 shows the effect of the flaws at radius 3/2 rp

- 0.3 RQ or 300 ym on the

peak radial tensile stress gradient monitored downstream at radius 2 rp

- 0.4 Rq or 400 ym.

A factor of four reduction (from the no flaw case) is observed near the surface for the

100 ym (R0=lmm) deep flaw. (Note that the stress profile upstream at r p

= 0.2 RD or 200 ym remains essentially unaffected by the flaws.;

508
Scaled Depth from Surface, z/R0 Figure 10. Effect of surface hoop cracks at r=3/2r

p on normalized (amax=356MPa) p&ak

fuxduxl tZYi&ilo. i>tA.QJ>t> VZUUA &zatzd dzpth at r=3/2r

p

for impact (V0=222m/s)

of ZnSe window by a water drop (R=lmm)

0 0.05 0.10 0.15 0o2D

Scaled Depth from Surface, z/RD Figure 11. Effect of surface hoop cracks at r=3/2r on normalized (amax=356MPa) pzak

fUutiaZ tzmltz A&IZ46 VZHMM Aealzd dzpth at r=2r

p

for impact (VD=222m/s) of

ZnSe window by a water drop (R0 =lmm)

509
5. Discussion

The effects of surface flaws on ceramic target damage from subsonic water drop impacts

are investigated using numerical code simulations. The pressure pulse due to water drop

impact generates a tensile stress wave with a steep near-surface gradient. Figure 12

shows a schematic view of the effect of a surface flaw on the near-surface tensile stress

profile. Upstream profiles remain essentially unchanged, while downstream profiles have

reduced gradients and peak value at surface.

Using the stress intensity approach of fracture mechanics, the stress intensity

factor, Kj, .can be obtained 6

, as shown in figure 12, and crack activation is defined by

the condition where Kj _> K^. Figure 13 shows the peak stress intensity factor for various

surface flaws based on the results of the crack activation experiments from Section 4 (see

fig. 10). For a toughness value of Kq = 1 MPa^nT, surface flaws of depth greater than

75 um will not propagate.

Thus, as a consequence of the steep stress gradient, surface crack activation under

water drop impact is dependent on the depth of surface flaws. In addition, a relatively

large surface flaw can perturb the stress field such that crack activation in its vicinity

is suppressed.

Pressure Pulse

Due to Water Drop Impact

Crack

Upstream

Tensile Stress

Profile Remains

Essentia 1 ly

Unchanged

Downstream

Tensile Stress

Profile Has

Reduced

Gradient

Computational

Cell

Axis of Symmetry

k

t = cr/frS

Figure 12. Schematic view of affect of a surface flaw on near surface tensile stress

profile and associated stress intensity factor.

The stress intensity factor is calculated from

where a is the radial stress located one-half a cell dimension, A/2, from the crack tip.

510
25 50 75 100 125

Flaw Depth, c (pm)

Figure 13. Peak stress intensity factor (and associated time) for various surface flaws

located at r=300 um for 222m/s impact of a 2-mm water drop onto a ZnSe window.

The authors wish to acknowledge the useful discussions with A. G. Evans of the

University of California, Berkeley (formerly at Rockwell International/Science Center).

References

[l] Evans, A. G. , Impact damage in ceramics, Report under Office of Naval Research

Contract No. N00014-75-C-0609, Rockwell International/Science Center, Thousand Oaks,

California (1977).

[2] Wilkins, M. L., Calculation of elastic-plastic flow, Report UCRL-7322, Rev. I,

Lawrence Radiation Laboratory, University of California, Livermore, California (1969).

[3] Rosenblatt, M. , Eggum, G. E. , DeAngelo, L. A., and Kreyenhagen, K. N. , Numerical

investigation of water drop erosion mechanisms in infrared-transparent materials,

Report AFML-TR-76-193, Air Force Materials Laboratory, Wright-Patterson Air Force

Base, Ohio (1976).

[4] Kreyenhagen, K. N. , Rosenblatt, M. , and Isbelle, T. R. , Cratering, mass loss, and re- sidual damage due to hypersonic particle impacts on ablative materials, Report AFWLTR-75-184, Air Force Weapons Laboratory, Kirtland Air Force Base, New Mexico (1976).

[5] Ito, Y. M. , Rosenblatt, M. , Perry, F. W. , and Eggum, G. E. , Analysis of water drop erosion mechanisms in infrared-transparent materials using microphysical fracture models,

Report AFML-TR-77-219, Air Force Materials Laboratory, Wright-Patterson Air Force

Base, Ohio (1977).

[6] Blowers, R. W. , On the response of an elastic solid to droplet impact, J. Inst. Maths.

Applies. 5, 167-193 (1969).

511

Reprinted from The Science of Ceramic Machining and Surface Finishing II, B. J. Hockey and

R. W. Rice, editors, National Bureau of Standards Special Publication 562 (U.S. Government

Printing Office, Washington, D.C., 1979).

SURFACE CHARACTERISTICS OF DIAMOND -MACHINED CERAMIC SURFACES

EVALUATED BY HIGH FLUENCE DIAGNOSTICS 1

M. J. Soileau, J. 0. Porteus , and D. L. Decker

Michelson Laboratory, Physics Division

Naval Weapons Center, China Lake, CA 93555

Ceramic surfaces prepared by precision diamond-machining are

compared with surfaces prepared by conventional polishing. Comparison is made on the basis of high fluence diagnostics. The results indicate that the diamond-machined surfaces have as high or

higher failure resistance under high fluence loading as surfaces

prepared using conventional techniques. Failure of both diamondmachined and polished surfaces are strongly dependent upon surface

structure and surface contamination. These findings indicate that

diamond-machining is an attractive alternative for generating

optical surfaces on ceramic materials.

Key Words: Ceramic surfaces diamond single-point machining;

high fluence; laser irradiation

1. Introduction

Diamond-machining has been shown to be an attractive alternative to

conventional polishing for producing high quality optical surfaces on metal

mirrors [l] 2

. This technique has also been applied to relatively soft dielectrics such as KC1 and NaCl with promising results [2], In this work ceramic

surfaces prepared by precision diamond-machining are compared with surfaces

prepared by conventional polishing. Comparison is made on the basis of visable effects produced by laser radiation. The thresholds for these effects

indicate surface quality in terms of cleanliness and surface finish. The

laser source is a single-pulsed commercial HF/DF laser with a maximum output

of 200 mJ per pulse. The pulses are approximately 100 nsec in duration with

wavelengths grouped at either 2.7 or 3.8 urn. The laser output is tightly

focused on the sample surface by a 127-mm focal length ZnSe lens.

Diamond machined surfaces were produced on single-crystal CaF 2 (Irtran

II). The details of the machining process are given elsewhere in these proceedings [3]. The results indicate that the diamond-machined surfaces have

failure resistance under high fluence loading of the same order of magnitude

as surfaces prepared using conventional techniques. In both types of surfaces, failure is initiated by surface detects and impurities.

2. Experimental

Disc-shaped samples of single- crystal CaF 2 and SrF 2 , and polycrystalline

MgF 2 (Irtran II) were turned on-center with a single-point diamond tool to

produce an optical surface. Part of the polished surface was left intact for

comparison with the machined surfaces.

^his work was performed under ONR Project No. RR02202, Task RR0220202,

Dr. A. M. Diness, Program Manager.

Figures in brackets indicate the literature references at the end of

this paper.

513
The tightly focused output of a commercially available, single-pulsed

HF/DF laser was used to probe the surfaces. The laser parameters are summarized in table 1. High fluence levels were obtained by focusing the laser output (2 0 -mm diameter) with a 50-mm diameter, 127-mm focal length ZnSe lens.

The focal spot diameter was approximately 55 urn, measured between the 1/e 2 points of the Gaussian spatial intensity distribution. The pulse energy was

monitored with a pyroelectric energy meter calibrated with a thermopile.

Incident and transmitted waveforms were monitored with fast photon drag

detectors. This facility is described in detail elsewhere [4].

Table 1. Laser parameters.

Pulse energy Wavelength Pulse width e" 2 Focal spot

(J) (ym) (nsec) diameter (ym)

HF 0.2 2.7 125 53

DF 0.2 3.8 95 57

The thresholds for visible spark, increased visible scatter, and/or truncation

of the beam transmitted through the sample were measured in air. Changes in

the surfaces were monitored with a 20-power microscope. The threshold for a particular effect was defined as that flux which produced the effect in 50% of

the sites. Each site was irradiated only once; the thresholds determined under

these conditions are called the one-on-one thresholds. Once the one-on-one

thresholds were determined, an undamaged site was irradiated five times with a flux equal to 80% of the one-on-one threshold. After this conditioning the

power was increased until a given effect was observed. Thresholds determined

in this manner are referred to as N-on-one thresholds.

3. Results

The one-on-one thresholds for the diamond-machined and polished surfaces

are summarized in table 2. The most significant features of these results are:

A. The thresholds for the diamond-turned surfaces are of the same order

as those of the polished surfaces.

B. The thresholds for the polycrystalline MgF 2 (Irtran II) are of an

order of magnitude lower than those of the single-crystal CaF 2 and

SrF 2 surfaces.

Table 2. Summary of one-on-one damage thresholds.

Specimens D-T low scatter D-T high scatter Polished

HF thresholds (10 9 W/cm 2

) CaF 2 7. 04 4.4 6. 75

SrF 2 -4 2. 86 3. 86 6.74

SrF 2 -2 6. 84 3.95 4.68

MgF 2 (poly) 0. 32 0.34

DF thresholds (10 9 W/cm 2

) CaF 2 10.6 7. 99 9. 37

SrF 2 -4 5.22 3.97 12.0

SrF 2 -2 8. 20 6.34

MgF 2 (poly) 0.51 0.47

514
C. For both polished and machined surfaces on all the materials studied,

the HF (2.7 urn) thresholds are lower than the DF (3.8 urn) thresholds.

D. The thresholds are dependent upon the surface quality of the materials.

The center- turned single-crystal samples showed alternating azimuthal

regions of high and low visible scatter, which were apparently determined by

the relationship of the cutting direction to the crystallographic axes.

Thresholds were generally found to be 501 lower in the high-scatter regions.

Figure 1 is a micrograph of a damage site in the low-scatter region and a damage site in the high-scatter region. Note that the damage in the highscatter region is influenced by the scattering centers which are clustered

along the machining marks.

LOW SCATTER REGION HIGH SCATTER REGION

I I 200 iim

Figure 1. Surface damage on diamond-machined SrF 2 . In the lowscatter region, the damage region essentially maps the incident

intensity distribution. In the high-scatter region, damage is

seen to extend along the direction of the machining grooves

where scattering centers are concentrated.

The lower thresholds measured at 2.7 ym as compared to 3.8 ym are suggestive that absorbed surface water is contributing to the failure for both the

polished and diamond-machined surfaces. More evidence for suspecting that

surface water plays an important role in the damage is seen in the results of

the N-on-one experiments on these surfaces. These results are summarized in

table 3. Note that for all cases where an N-on-one effect is seen, the effect

(increase in damage threshold) is greater at 2.7 ym than at 3.8 ym.

Specimens

Table 3. N-onHF D-T

one effect ( DF D-T% increase)

HF

polished

DF

polished

CaF 2 54 40 47 26

SrF 2 -4 0

SrF 2 -2 21 13 0

MgF 2 (poly) 21 14 0

These results suggest that for the N-on-one irradiations surface contaminants are being removed from the surface, and that the contaminants are selectively absorbing for the HF lines. The peak of the water absorption band corresponds to the strongest HF line (2.7 ym) ; thus, absorbed surface water is a likely contributor to surface damage at HF frequencies.

As can be seen from table 2, the damage thresholds for the Irtran II are

an order of magnitude lower than for the single- crystal material. Near threshold the Irtran II failed by a change in visible scatter at the surface or

515
within the bulk near the surface. No truncation of the transmitted pulse was

observed near threshold. This evidence and the micrographs in figures 2 and

3 indicate that failure is initiated by defects within the material or at the

surface, and that the method of surface preparation has little or no effect.

The mechanism of failure for the Irtran II is probably heating of absorbing

defects within the focal volume.

SURFACE DAMAGE BULK DAMAGE

I I 100 jim

Figure 2. DF laser damage in polished Irtran II. Note that the

damage site does not map the incident beam. In all cases the

laser was focused on the surface. Note the high concentration

of defects. These defects are uniformly distributed throughout

the material.

SURFACE DAMAGE BULK DAMAGE

100 Aim

Figure 3. DF laser damage in diamond- turned Irtran II. Note the

high concentration of defects. These defects are uniformly distributed throughout the specimen.

For the polished and diamond- turned surfaces on the single-crystal materials, laser damage was accompanied by a bright, visible spark, an increase in

visible scattering, and an abrupt truncation of the beam transmitted through

the specimen. An example of the truncation of the transmitted beam is seen

in figure 4. Note that failure occurs at the peak of the pulse, and beam

truncation occurs within the response time of the detector (less than 1.5

nsec) . This type of behavior is strongly suggestive of electron avalanche

breakdown. The probable mechanism of failure is avalanche break-down initiated by surface defects and/or contaminants.

The effect of the orientation of the residual machining grooves with

respect to the laser electric field was investigated. The one-on-one threshol<

was measured for the cases where the field was parallel to the perpendicular

to the machining grooves. Polarizable defects perpendicular to the incident

field should enhance the coupling of the laser energy to the surface and lower

the threshold [5] . The threshold was found to be approximately 101 lower for

the case of the field perpendicular to the grooves than for the case of the

516
NO DAMAGE DAMAGE NEAR THRESHOLD

Figure 4. Transmitted waveform distortion. Note that nearthreshold damage occurs at the peak of the pulse within

the fall time of the detector (<1.5 nsec) .

field parallel to the grooves. This decrease in threshold is as predicted by

theory and is of little consequence for these low index materials. However,

for high index materials such as GaAs, Si, ZnSe, and Ge, the effect could be

significant since the induced polarization field caused by such defects is

proportional to the index of refraction. In figure 5 one can see the effects

of the interference of the induced polarization field and the incident laser

field for the case of the incident laser field perpendicular to the residual

machining grooves.

E E

I l 100 iim

Figure 5. Effects of laser polarization in diamond- turned SrF 2

. Note the "ripples" in the figure on the right. These ripples

are the result of interference between the impressed field

and the local field due to polarization charges induced at

surface defects by the incident field.

4. Summary and Conclusions

Diamond- machined surfaces on CaF 2 , SrF 2 , and Irtran II have HF/DF laser

damage thresholds of the same order as polished surfaces on these materials.

The mechanism of failure for s ingle - crystal CaF 2 and SrF 2 appears to be electron avalanche breakdown initiated by defects and impurities in polished and

diamond-machined surfaces. The mechanism of failure for Irtran II appears to

be absorption by bulk defects and is thus independent of the method of surface preparation. The lower thresholds at HF frequencies compared to those

at DF frequencies and the larger increase in thresholds (N-on-one effect) at

HF compared to DF indicates that- absorbed water has a significant effect on

the damage thresholds of these surfaces at these frequencies.

517
References

The following review article summarizes the work on diamond- turned metal

surfaces and is a good guide to work in this area; Porteus, J. 0. and

Decker, D. L. , "Metal Mirror Damage and Precision Diamond-Turning," in

Laser Damage of Optical Materials , A. J. Glass and A. H. Guenther, eds

.

(Academic Press, New YorFj 1979) , in press.

Soileau, M. J., Bennett, J. M. , Porteus, J. 0., Faith, W. N., Jernigan,

J. L. , and Saito, T. T. , "10.6 ym Pulsed Laser Studies of Diamond-Turned

KC1 Window Surfaces," in Laser Induced Damage in Optical Materials : 197 6 , A. J. Glass and A. H. Guenther, eds., National Bureau of Standards,

Washington, DC (NBS Spec. Publ. 462, 1976), pp. 149-157.

Decker, D. L. , Bennett, J. M. , and Grandjean, D. J., "Optical and Surface

Physical Characteristics of Diamond-Machined Infrared Windows," these

proceedings

.

Van Stryland, E. W. Bass, M. , Soileau, M. J., and Tang, C. C, "Pulsed

HF/DF Laser Damage in Window Materials," in Laser Induced Damage in

Optical Materials : 197 7 , A. J. Glass and A. H. Guenther, eds., National

Bureau of Standards, Washington, DC (NBS Spec. Publ. 509, 1977),

pp. 118-125.

Temple, P. A., and Soileau, M. J., "Resonant Defect Enhancement of the

Laser Electric Field," in Laser Induced Damage in Optical Materials : 1976 , A. J. Glass and A. H. Guenther , eds . , National Bureau of Standards,

Washington, DC (NBS Spec. Publ. 462, 1976), pp. 371-378.

518
JOINT MEETING - ACS BASIC SCIENCE DIVISION AND SYMPOSIUM

THE SCIENCE OF CERAMIC MACHINING AND SURFACE FINISHING I

Conference Attendees

A

Scott J_._Abel

Corning Glass Works

31501 Solon Road

Solon, OH 44139

S . Adachi

Hoya Optics USA, Inc.

2200 Sand Hill Road, Suite 200

Menlo Park, CA 94025

Marc A. Adams

Jet Propulsion Laboratory

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Art Ahlbrecht

3M Co.

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Rex Ashbrook

Kennametal Inc.

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Richard L . Ashbrook

NASA Lewis Research Center

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Charles F. Bersch

U.S. Navy

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Howard D. Blair

Ford Motor Co

P. O. Box 2053

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Tim Bornemann

Libbey-Owens-Ford Company

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R. Bowles

British Glass Industry

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ENGLAND

R. C. Bradt

Pennsylvania State Univ.

118 Mineral Industries Bldg.

University Park, PA 16802

Robert H. Bragg

Dept. of Matls. Science and

Mineral Engineering

Univ. of California

Berkeley, CA 94720

B

Francis I. Baratta

Army Materials and Mech.

Watertown, MA 0217 2

Michael Baxter

Fair-Rite Products Corp.

One Commercial Row

Wallkill, NY 12589

Harold E. Bennett

Naval Weapons Center

Code 38101

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Res. Ctr.

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Westinghouse R&D Center

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Philips Research Lab.

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Eindhoven, The Netherlands

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United Technologies

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AC Spark Plug - GM

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Leslie Bretland

Monsanto Research Corp.

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Philip L. Berneburg

GTE Sylvania, Inc./Chem. & Met. Div.

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Philips Research Laboratories

Eindhoven, 5600 MD

The Netherlands

519
S . D . Brown

Univ. of Illinois

Dept. of Ceramic Engineering

Urbana, IL 61801

Joseph E. Burke

General Electric Co.

P . 0 . Box 8

Schenectady, NY 12301

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Rowland Cannon

Mass. Inst, of Tech.

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Rockwell International

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Trans-Tech, Inc.

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Jet Propulsion Lab.

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Georgia Inst, of Technology

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Univ. of Rhode Island

Kingston, RI 02881

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Monsanto Research Corp.

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Rockwell International Sci. Ctr.

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Washington, NJ 07882

L. P. Cook

National Bureau of Standards

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Institut Cerac S. A.

Chemin Des Larges Pieces

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Univ. of Southern Calif.

Materials Science Dept.

Los Angeles, CA 90007

Robert E. Cuthrell

Sandia Laboratories

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Ivan B. Cutler

Univ. of Utah

Salt Lake City, UT 84112

D

Kiran H. Dalai

Cornell University

131 Bard Hall

Dept. of Materials Science & Eng.

Ithaca, NY 14850

D. L. Decker

Michelson Laboratory

Naval Weapons Center

China Lake, CA 93555

Arthur M. Diness

Office of Naval Research

Arlington, VA 22217

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General Electric Corp. Res. & Dev

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Schenectady, NY 12301

Louis P . Domingues

Trans-Tech, Inc.

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Alan L. Dragoo

National Bureau of Standards

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NASA - Lewis Res. Ctr.

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3M Co. - Industrial Abrasives Div

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Fa. Feldmuhhle AG

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T. E. Easier

Pennsylvania State Univ.

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University Park, PA 16802

George O. Ellis, Sr.

Ellis Ceramtek, Inc.

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A. G. Evans

Univ. of California

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Katherine T. Faber

The Carborundum Company

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Univ. of Illinois

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Ross Firestone

IIT Research Institute

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National Bureau of Standards

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TRW/Marlin Rockwell Div.

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National Bureau of Standards

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National Bureau of Standards

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GTE Sylvania Inc.

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Jack T. Gardner

TRW Systems

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G. E. Gazza

Army Materials & Mech.

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Richard L. Gentilman

Raytheon Res. Division

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Bell Telephone Lab., Inc.

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Univ. of Rhode Island

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National Micronetics Inc.

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University of Cambridge

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Pennsylvania State Univ.

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Barry L. Hammond

The David W. Taylor Naval

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Al Hare

Ellis Ceramtek, Inc.

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Lebanon, NJ 08833

D. P. H. Hasselman

Virginia Polytechnic Inst.

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Eaton Corp. - Engr. & Research Center

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Corning Glass Works

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Nairn Hemmat

Allied Chemical

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S . P . Hepburn

PPG Industries

153 Lingay Drive

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H . Herman

State Univ. of NY

Materials Science

Stony Brook, NY 11794

Arthur H. Heuer

Case Western Reserve Univ.

Olin 303, Univ. Circle

Cleveland, OH 44106

L. W. Hobbs

Case Western Reserve University

Cleveland, OH 44106

Michael Hoch

Univ. of Cincinnati, Loc 12

Dept. of Matls. Sci. and

Metallurgical Engineering

Cincinnati, OH 45221

Bernard Hockey

National Bureau of Standards

Center for Materials Science

Washington, D. C. 20234

H . Hohnke

Univ. of Michigan

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William D. Horn

Allied Chemical

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Charlotte, NC 28210

Robert S . Houston

Norton Co.

New Bond St.

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I

Osamu Imanaka

Kanazawa Univ., Faculty of

Engineering

40-20, 2 Chome, Kodatsuno

Kanazawa, JAPAN 920

522
Moshe Ish-Shalom

Israel Ceramic & Silicate Inst.

Technion City

Haifa, ISRAEL 32000

Marvin I to

California Research & Tech., Inc.

6269 Variel Ave., Suite 200

Woodland Hills, CA 91367

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Univ. of Rhode Island

Dept. of Mechanical Eng.

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Randall R. Kinnick

Detroit Diesel Allison of

General Motors Corp. Matls.

P. 0. Box 894 Speed Code W5

Indianapolis, IN 46206

J

Karl Jakus

Univ. of Massachusetts

Mechanical Engineering

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B . E . Johnson

Owens - Illinois, Inc.

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Toledo, OH 43666

Curtis A. Johnson

General Electric CR & D

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D. Lynn John

Northwestern Univ.

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Western Electric, Inc.

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Electric Power Res. Inst.

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Hillard M. Kahan

Ampex Corporation

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IBM Corporation

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Ceramic Finishing Co.

P. 0. Box 498

State College, PA 16801

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German Aerospace Res.

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Honeywell Inc.

10701 Lyndale Ave. S.

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Univ. of California,

Berkeley

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McMaster University

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Purdue University

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R. L. Landingham

Univ. of California

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F. F. Lange

Rockwell Science Center

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ITT Research Institute

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523
Kai-Hung Lau

SRI International

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D. William Lee

Arthur D. Little, Inc.

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Minydung Lee

General Electric Co.

Corp. R&D Center

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David Lewis

NYS College of Ceramics

Alfred University

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Harold Lombard

DAta Materials

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Victor Lou

General Electric

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R.P.I.

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Howard McCollister

Owens-Illinois

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Edward J. McCormack

CS Draper Lab.

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William J. McDonough

U. S. Naval Res. Laboratory

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Kelly D. McHenry

Honeywell

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Robert McKinley

Plessey/Frenchtown

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K. R. McKinney

Naval Research Lab.

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Edward A. Maguire

Raytheon Co.

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Stephen Malkin

Faculty of Mechanical Eng.

Technion - Israel Inst, of

Technology

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J. A. Mangels

Ford Motor Company

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William R. Manning

Champion Spark Plug Co.

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R. H. Marion

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Univ. of Maryland

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Pratt & Whitney Aircraft

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Alan D. Miller

University of Washington

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Drago Milovic

Western Gold and Platinum Co.

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T. E. Mitchell

Case Western Reserve Univ.

Cleveland, OH 44106

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Kanazawa University

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Ampex Corporation

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Univ. of Pittsburgh

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IBM

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Ford Motor Co.

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Ishverlal Naik

AiResearch Casting Co.

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Sumitomo Special Mtls. Co. LTD.

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IBM

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11380 Harrison

Romulus, MI 48174

R. E. Novak

RCA Laboratories

Princeton, NJ 08540

0

Linus U. Ogbuji

Case Western Reserve Univ.

1589 Coventry Rd. , #12

Cleveland, OH 44118

C. D. Olson

National Bureau of Standards

Center for Materials Science

Washington, DC 20234

Shinji Otsuki

Sumitoko Electronic Matls.

Co. LTD

282T-Ukumo Omachi-Machi

Saga, JAPAN

Peter Otten

Union Carbide Corp.

8888 Balboa Avenue

San Diego, CA 9212 3

P

John A. Palm

General Electric Co.

Corporate Res. & Dev. Ctr.

P. O. Box 8

Schenectady, NY 12301

Edward D. Parent

GTE Sylvania

Portsmouth Ave.

Exeter, NH 03833

Robert A. Penty

Hague International

3 Adams St.

South Portland, ME 04106

Harry Perry

Consolidated Beryllium Ltd.

Marble Hall Road, Milford Haven

Dyfed, UNITED KINGDOM

525
Ronald L. Peters

VR/Wesson Co. - A Division

of Fansteel

800 Market St.

Waukegan, IL 60085

Daniel R. Petrak

Harbison - Walker Ref. Co.

Div. Dresser Ind.

Garber Research Center

P. O. Box 98037

Pittsburgh, PA 15227

Bruce Pletka

Dept. of Metallurgical Eng.

Michigan Technological Univ.

Houghton, MI 49931

Robert Pohanka

Office of Naval Research

Code 471

Arlington, VA 22217

Q

CarrLane Quackenbush

GTE Labs.

40 Sylvan Road

Waltham, MA 02754

R

Bharat Rawal

AVX Ceramics Corp.

P. 0. Box 867

Myrtle Beach, SC 29577

Dennis W. Readey

Ohio State Univ.

2041 College Road

Columbus, OH 43017

Robert R. Reeber

Dept. of Energy

20 Massachusetts Ave., N.W.

Washington, D. C. 20545

Gary M. Renlund

General Electric CRD

P. 0. Box 8

Schenectady, NY 12301

William H. Rhodes

GTE Laboratories

40 Sylvan Road

Waltham, MA 02154

Paul E. Rempes

Champion Spark Plug Co.

Ceramic Division

20,000 Conner Ave.

Detroit, MI 48234

R. W. Rice

Naval Research Laboratory

Code 6360

Washington, DC 20375

Subhash Risbud

Lehigh Univ.

Coxe Lab. #32

Bethelem, PA 18015

J. E. Ritter

Univ. of Massachusetts

Dept. of Mechanical Eng.

Amherst, MA 01003

M. W. Robare

AiResearch Manufacturing Co.

Phoenix, AZ 85034

Carl Robbins

National Bureau of Standards

Center for Materials Science

Washington, D. C. 20234

Jack A. Rubin

Kyocera International

8611 Balboa Ave.

San Diego, CA 92123

Robert Ruh

Air Force Matls. Lab.

Wright Patterson AFB , OH 45433

C. K. Russell

United States Steel Corp.

125 Jamison Lane

Monroeville, PA 15146

Ralph A. Russell

IBM Corp.

Hwy 52 No.

Rochester, MN 55901

S

H. Sato

Purdue University

West Lafayette, IN 47906

R. W. Scheidecker

Ames Laboratory, USDOE

Iowa State University

Ames, IA 50011

526
Frederick Schmid

Crystal Systems, Inc.

35 Congress St.

Salem, MA 01970

Detlev Schultz

Fair-Rite Products Corp.

One Commercial Row

Wallkill, NY 12589

Susan R. Schuon

Michigan State Univ.

Dept. of Metallurgy Mech.

and Materials Science

East Lansing, MI 48824

Nirmal Sharma

Univ. of Illinois

Dept. of Ceramic Engineering

204 Ceramics

Urbana, IL 61801

J. E. Sheehan

General Atomic Co.

P. O. Box 81608

San Diego, CA 92138

L. B. Sibley

SKF Industries

1100 First Ave.

King of Prussia, PA 19406

Wilbur G. Simmons

AF Office of Scientific Res.

Bldg. 410

BAFB, DC 20332

S. C. Singhal

Westinghouse Res. & Dev. Ctr.

1310 Beulah Road

Pittsburgh, PA 15235

J . T . Smith

GTE Laboratories

40 Sylvan Rd.

Waltham, MA 02154

R . H . Smoak

The Carborundum Co.

R&D Div., P. O. Box 1054

Niagara Falls, NY 14302

W. J. Smothers

Bethlehem Steel Corp.

Bethlehem, PA 18016

M. J. Soileu

Michelson Lab

Naval Weapons Ctr.

China Lake, CA 93555

Lawrence R. Sparrow

COMSAT Laboratories

P. O. Box 115

Clarksburg, MD 20734

Richard M. Spriggs

Lehigh University

c/o President's Office

Bldg. 27

Bethlehem, PA 18015

M. Srinivasan

The Carborundum Co.

R&D Division

P. O. Box 1054

Niagara Falls, NY 14302

Read W. Steward

Eaton Corporation

26201 Northwestern Hwy

P. O. Box 766

Southfield, MI 48037

V. S. Sturican

The Pennsylvania State Univ.

328 MI Bldg.

University Park, PA 16802

Douglas L. Swanson

IBM

3605 Highway 52 North

Rochester, MN 55901

T

Donna L. Taylor

Corning Glass Works

Sullivan Park

Corning, NY 14830

Leslie L. Terner

Ford Motor Co.

7400 Finnegan Dr.

West Bloomfield, MI 48033

Bruce Thomson

United Technologies Res.

Silver Lane

E. Hartford, CT 06108

R. Thomson

National Bureau of Standards

Center for Materials Science

Washington, D. C. 20234

Nancy S. Tighe

National Bureau of Standards

Center for Materials Science

Washington, D. C. 20234

527
R. E. Tressler

Pennsylvania State Univ.

118 Mineral Industries Bldg.

University Park, PA 16802

Stuart Tuffs

Norton Co .

,

Industrial Ceramics Div.

1 New Bond Street

Worcester, MA 01601

V

Jim Vagt

IBM

HY 52, North

Rochester, MN 55901

John D. Venables

Martin Marietta Labs

.

1450 S. Rolling Road

Baltimore, MD 21227

John V. White

Alberox Corp.

Industrial Park

New Bedford, MA 02745

Sheldon M. Wiederhorn

National Bureau of Standards

Fracture and Deformation Div.

Washington, DC 20234

Ben A. Wilcox

National Science Foundation

1800 G St. , NW

Washington, DC 20550

W. S. Williams

Univ. of Illinois

Matls. Res. Lab.

Urbana-Champaign, IL 61801

Roger E. Wilson

Naval Surface Weapons Center

White Oak

Silver Spring, MD 20910

W

John B. Wachtman, Jr.

National Bureau of Standards

Center for Materials Science

Washington, D. C. 20234

Walter R. Wagner

IBM Corporation

P. 0. Box 12195

Dept. E81-B101

Research Triangle Park, NC 27709

James M. Wimmer

Univ. of Dayton

300 College Park

Dayton, OH 45469

T. W. Witkowski

Cast Western Reserve Univ.

Cleveland, OH 44106

C. M. Wu

Naval Research Lab.

4555 Overlook Avenue

Washington, DC 20375

G. C. Walther

ITT Research Institute

10 West 35 Street

Chicago, IL 60616

Gerald Q. Weaver

Norton Co. , Industrial Ceramics Div.

1 New Bond Street

Worcester, MA 01606

Francis H. Weikel

Corhart Refractories Co.

1600 W. Lee St.

Louisville, KY 40201

Richard D. Weltzin

IBM GSD

3605 Hy 52 North

Rochester, MN 55901

Nobuo Yasunaga

Etectrotechnical Laboratory

4-1/ 5-chome, Mukodai-machi

Tanashi-shi, Tokyo

JAPAN 188

Hun C. Yeh

Cleveland State Univ.

1983 East 24 Street

Cleveland, OH 44115

C. F. Yen.

Oak Ridge National Lab.

Oak Ridge, TN 37830

Thomas J. Whalen

Ford Motor Co.

Box 2053

Dearborn, MI 48121

528
AUTHOR INDEX

Andersson, C. A., 463

Adachi, S., 417

Bas, M. , 283

Bennett, H. , 337

Bowles, R., 489

Bradt, R. C, 455

Bratton, R. J., 463

Brehm, R. 61

Broese van Groenou, A., 61,43,147,191

Casey, A. W. , 231

Cook, L. P., 407

Cooke, F., 255

Cooper, G. A. , 115

Copley, S. M. , 283

Countermine, T. A. , 455

Cuthrell, R. E . , 139

Davidson, D. L. , 395

Dean, S. K. , 107

Decker, D. L. , 293, 315

Doyle, E. D., 107

Duwell, E. J., 201

Kino, G. S., 379

Kirchner, H. P., 23

Koepke , B . G . , 75

Kokaji, Y . , 325

Landingham, R. L. , 231

Lankford, J., 395

Levinger , R. , 305

Lindahl, R. 0. , 231

Lokken, R. C, 201

R. C, 201

Maan, N. , 43

Malkin S. , 93, 305

McCoy, L. G., 247

McKinney, K. R. , 477

Mecholsky, J. J., Jr.,

Messing, G. L .

, 247

Miles, P., 255

Miyade, E., 417

Obara, A. , 171

Okutomi, M. , 157

Olson, C. D., 407

Easier, T. E., 455

Evans, A. G. , 1, 185, 379

Farabaugh, E. N. , 407

Field, J. E., 15

Firestone, R. F., 261

Porteus, J. 0., 315

Rice, R. W., 185,351,42

Richard, D. M. , 23

Richardson, D. W. , 209

Robare, M. W. , 209

Rosenblatt, M. , 501

Gentilman, R. 255

Grandjean, D. J., 293

Grossman, D. G. , 221

Gruver, R. M. , 23

Hagan, J. T . , 15

Harada, S., 417

Hassler, C. R. , 247

Huerta, M. , 93

Ikeda, M. , 325

Imanaka, 0. , 157, 171, 317

Ito, Y. M. , 501

Izumitani, T. , 417

Kawakami, S., 317

Khuri-Yakub, B. T. , 379

Kimura, S., 325

Soileau, M. J., 315

Stokes, R. J., 75

Sugita, T. , 317

Swain, M. V., 15

Tarumi, N. , 171

Taylor, D. L. , 221

Terner , L. L. , 483

Tressler, R. E . , 4 55

Veldkamp, J. B. rj. , 43

Wallace, R. G., 283

Whalen, T. J., 483

Wu, C. Cm., 477/

Yamada, A. , 32 5 Yasunaga, N. , 171

529
A

SUBJECT INDEX

E

Abr.asi onless machining, 262

Abrasive machining, 1, 24, 210(Si), 216 (Si

dainacTe, 457

flaw distributions, 352

models 5/ If 71

rates, 4, 61, 210, 216

surfao e flaws, 351, 429

two sua :face technique, 255

Acoustic emission, 140, 142

Acoustic , surface waves, 379

A1 2 0 3 , 75, HI, 142, 157, 171, 205, 247, 2 317, 3 25, 351, 395, 429, 455

Auger speci troscopy, 345

B

Elastic emission machining, 159

Ni*) apparatus, 160, 161

model, 160

Electric discharge machining, 266

Electrochemical efficiency, 314

Electrochemical grinding, 305

rate, 306

surface damage, 311

Electron beam machining, 168, 267

Electron channeling, 397

<5, Electron diffraction, 180, 320

Electron spectroscopy for Chemical Analysis

(ESCA) , 180, 236, 332, 345

Ellipsometry , 346

Environmental effects, 44, 54, 115, 139, 186

Belt grindin g, 202

BeO, 429

B^C, 75, 351, 429

CaF 2 , 293, 351 , 513

Cam grinding, .214

Cemented carbides, 305

Chemical millir, >g, 266

Chemomechanical effects, 115, 139, 162

Coated abrasives?, 202

Contour grinding', 213

Cracks,

Hertzian cone, 23

lateral, 5, 16, 24

radial/median, 16, 24

D

Diamond sawing, 362

surface flaws, 36 3-365

Drilling, 116 (quart; 0 , 120 (granite) , 121 (marble) , 142' (glass, A1 2 0 3

, Si) , 224 (glass ceramic)

damage, 143

effect of:

additives, 122

drill sharpness,, 1?24

environment, 140, 142

lubrication, 224

tool geometry, 228

force, 225

rates, 116, 120, 122, 124, 134, 135

power, 226

power dissipation, 133

tool wear, 124

Ductile-to-Brittle transition, 143

Fracture mechanisms, 5, 16, 24, 158, 354, 431

Fracture origins, 355-374, 435-443, 468, 479,,

480

Fracture stress (see strength)

Fracture toughness, 3, 25, 71, 84, 140, 384,

469-471, 510

Fractography , 351

Ferrites, 44, 61, 75, 142, 147

Friction, 46, 117

Formed-wheel grinding, 214

GaAs, 395

Glass, 15, 27, 48, 61, 93, 109, 140, 255,

417, 489

chemical durability, 424

cleaning, 418

Glass-Ceramic, 221

Glass polishing,

compounds, 234

Granite, 120, 203

Green machining, 209

Grinding, (AI2O3) 79, 111, 205, 251;

(ferrites) 66, 149; (fire brick) 205;

(glass) 67, 93, 107, 192, 429; (granite)

I

203; (guartz) 206; (Si 3 NO 78, 216, 473,

477

correlation with material properties, 75

crack size, 77, 351, 473, 479; (see also,

fracture origins)

damage 68, 80, 99, 109, 112, 192, 200, 251/t

398, 498-500.

efficiency, 202

forces, 62, 65, 68, 78, 85, 86, 99

hardness, 66

530
SUBJECT INDEX (continued)

models, 1, 71, 103, 191

rates, 42, 192, 204-206, 216, 492

size effect, 69

specific energy, 309

techniques, 209

vibrations , 64

H

Hardness, 3, 49, 82, 185, 308, 402

relation to machining, 187

Hydrodynamic machining, 151, 263

rates (fire brick) , 265

Hydrogen embrittlement, 141

I

Impact cracking, 5, 24, 501

Indentation cracking, 5, 16, 24, 386

Infrared domes, 255

Infrared window materials, 293

Ion-beam machining, 167, 261

rates, 272

Ion microprobe, 164, 372

Interferometry, 259, 332, 339

L

Laser beam damage, 231, 348

effect of surface finish, 514-516

threshold, 302

Laser beam machining, 268, 283

groove depth, 285

groove profile, 286

specific cutting energy, 286

removal rate, 289

surface finish, 288

Laser glass, 231

Latent surface scratches, 417

M

Machining (see abrasive machining)

Marble, 121

MgAl204 , 351, 429

MgF 2 , 293, 351, 513

MgO, 48, 75, 118, 429

Microstructure, 186, 222, 360, 486

Mullite, 351

N

Numerical simulation, 501

0

Optical glass, 107

Optical quality, 231, 259, 279, 343

P

Pantographic grinding, 215

Plastic flow, 5, 16, 47, 103, 158, 398 zone depth, 401, 403, 404

Polishing, (A1 2 0 3 ) 162, 173, 317, 325; (glass)

231, 417; (Si) 178; (quartz) 181

apparatus, 172, 173, 318

damage, 166, 180, 238, 239, 331, 367, 398,

418-420

effect of environment, 163

hydration, 165

mechanochemical , 164, 171

model, 159, 174

rates, 165, 173, 242, 319, 330

surface analysis, 320

contamination, 231, 345

finish, 166, 180, 331, 411

Profile grinding, 214

Prosthetic devices, 247

PZT, 7 5

Q

Quartz, 115, 171

R

Raman spectroscopy, 167

Reaction rates, 175

Rehbinder effect, 115, 139

Residual stresses, 2 Rolling, 56

S

Sapphire (see AI2O3)

Scanning electron microscopy (SEM) , 80, 99,

109, 112, 126, 143, 176, 233, 248, 251,

285, 298, 386, 396, 407, 457, 498

Scratch test, (ferrites) 46; (glass) 27, 50;

(MgO) 52; (Si 3 NiJ 28

apparatus, 45, 326

crack size, 28, 29, 51

damage, 28, 29, 44, 46, 50, 53, 326

environmental effects, 54

hardness, 50

size effect, 49

Secondary electron emission, 407

Secondary ion mass spectrometry (SIM) , 236,

346

SiALON, 283

Silicon, 142, 171, 325, 395

Silicon carbide (SiC) , 283, 395, 429, 483

Silicon nitride (Si 3 Nit) 28, 75, 209, 283,

351, 385, 393, 463, 477

Single point grinding, 29 (glass) , 32 (Si 3 Nit)

crack size, 32, 34, 35, 37

531
SUBJECT INDEX (continued)

damage, 33, 35

forces, 31, 32, 35, 37

groove depth, 30, 34

Single-point diamond turning, 293, 513

material removal, 296

surface finish, 296

Slow-crack growth, 372

Specific energy, 48, 66, 96

SrF 2 , 293, 513

Strength, (Al 2 0 3 ) 250, 436, 441, 442, 458?

(BeO) 436, 445; (B 4 C) 358, 439; (CaF 2 ) 358; (glass) 102, 358, 494; (MgAl 2 0 4 ) 358; 438; (MgF 2 ) 358, 362; (MgO) 436;

(mullite) 358; (SiC) 429, 447, 484;

(SisNO 211, 213, 219, 465, 478

degradation 28, 29

effect of:

edge beveling, 459

flaw geometry, 358

grain size, 429

grinding direction, 213, 250, 358, 459

machining rate, 455, 459

machining technique, 219

oxidation, 478

surface coating, 484, 485

surface finish, 436, 459, 478, 484, 485,

493

ion beam machined, 273

size effects, 367, 375, 486

Surface characterization, 79, 99, 164, 180,

236, 299, 320, 331, 337-350, 379, 395,

407

Surface flaws, 501

Surface flaw detection, 379

Surface melting, 68, 204

Surface roughness, 228, 253, 300, 320, 340,

388, 411, 483

T

TEM, 109, 110, 340

Thermal stresses, 9

Tool materials, 210

Transducers, 381

Tungsten carbide, 305

Turbine components, 209, 390

U

Ultrasonic machining, 216

W

Water-drop impact, 501

Wear,

belt, 203

debris, 98, 110, 113, 174, 328

effect of additives, 132

grinding wheel, 147, 152, 154

profiles, 150

tools, 124, 126, 137, 210, 227

Weibull distribution, 455, 466

X

X-ray diffraction, 174, 176, 328

X-ray topography, 48, 331

Y

Y-deflection modulation, 407

Young's modulus, 185

Yttraloz (Y 2 0 3 ) , 351

Z

Zeta potential, 117, 140

Zinc selenide (ZnSe) , 501

532

6 U. S. GOVERNMENT PRINTING OFFICE : 1979 0 - 301-592
NBS-114A (REV. 9-78)

U.S. DEPT. OF COMM.

BIBLIOGRAPHIC DATA

SHEET

1. PUBLICATION OR REPORT NO.

NBS SP 562

2. Gov't Accession No. ; 3. feci

4. TITLE AND SUBTITLE

The Science of Ceramic Machining and Surface Finishing II

5. Publication Date October 1979

S. Perf«tto ispitation Code

7. AUTHOR(SJ Bernard

Roy W.

J. Hockey, National Bureau of Standards

Rice, Naval Research Laboratory, editors

8. Performing Organ. Report No.

9. PERFORMING ORGANIZATION NAME AND ADDRESS 'ask/Work itet No.

NATIONAL BUREAU OF STANDARDS

DEPARTMENT OF COMMERCE

WASHINGTON, DC 20234

11. Contract/Grant No.

12. SPONSORING ORGANIZATION NAME AND COMPLETE ADDRESS (Street, City, state, ZIP)

Sponsored in part by: Office of Naval Research, 800 N.

Juincy St., Arlington, VA 22217, Air Force Office of Scientific

Research, Boiling AFB, Washington, DC 20332, and American

:eramic Society, 65 Ceramic Dr., Columbus, OH 43214

13. Type of Report & Period Covered

N/A

14. Sponsoring Agency Code

i— _

15. SUPPLEMENTARY NOTES

Library of Congress Catalog Card Number: 79-600149

| | Document describes a computer program; SF-185, Fl PS Software Summary, is attached.

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 Second Symposium on the Science of

Zeramic Machining and Surface Finishing held at the National Bureau of Standards in

5aithersburg, Maryland, on November 13-15, 1978. The symposium was jointly sponsored

Dy the Air Force Office of Scientific Research, the American Ceramic Society, the

office of Naval Research, and the National Bureau of Standards. The purpose of the

symposium was to review recent progress by discussion of current problems and research

f forts. The proceedings contain forty of forty-three research papers, together with

floor discussions, presented at the Symposium. Specific topics covered include:

L) mechanisms of abrasive machining and surface finishing; 2) technology of machining

md surface finishing; 3) characterization of machined surfaces; and 4) effects of

machining on the properties of ceramics. Also included is an edited discussion on

lardness and its relation to machining.

17. KEY WORDS (six to twelve entries; alphabetical order; capitalize only the first letter of the first key word unless a proper name;

separated by semicolons)

brasion: ceramics; ceramic machining; effects of machining on ceramic properties;

lardness; surface characterizations; surface finishing

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BIBLIOGRAPHIC SUBSCRIPTION SERVICES

The following current-awareness and literature-survey bibli- ographies are issued periodically by the Bureau:

Cryogenic Data Center Current Awareness Service. A litera- ture survey issued biweekly. Annual subscription: Domestic, $25.00; Foreign, $30.00.

Liquefied Natural Gas. A literature survey issued quarterly.

Annual subscription: $20.00.

Superconducting Devices and Materials. A literature survey

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services to National Bureau of Standards, Cryogenic Data

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